Polímeros: Ciência e Tecnologia (Polimeros) 2nd. issue, vol 27, 2017

New materials like plastics are advancing vehicle technologies. But getting the best results is a big challenge.

To assist, SABIC offers industry-leading expertise in designing with a wide range of engineered materials. Across the entire vehicle. Because no matter what obstacles may hold our customers back, we’re there with ‘Chemistry that MattersTM’ to help them drive forward.

Polímeros

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VOLUME XXVII - Issue II - APR/JUNE - 2017

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ISSN 0104-1428 (printed) ISSN 1678-5169 (online)

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

and

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

Editorial Committee

Marco-Aurelio De Paoli (UNICAMP/IQ) - President

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

Members

A ss o ci at e E d i t o r s

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)

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

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

www.editoracubo.com.br

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

Financial support:

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

Quarterly v. 27, nº 2 (Abr./Maio/Jun. 2017) ISSN 0104-1428 ISSN 1678-5169 (electronic version)

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

1. Polímeros. l. Associação Brasileira de Polímeros.

Polímeros, 27(2), 2017

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

I I I I I I I I I I I I I I I I I

Editorial Section Agenda.................................................................................................................................................................................................E3 Funding Institutions.............................................................................................................................................................................E4 News....................................................................................................................................................................................................E6

O r i g in a l A r t ic l e Mixing process influence on thermal and rheological properties of NBR/SiO2 from rice husk ash Ana Maria Furtado de Sousa, Augusto Cesar de Carvalho Peres, Cristina Russi Guimarães Furtado and Leila Lea Yuan Visconte............. 93

Gelatin capsule waste: new source of protein to develop a biodegradable film Camila de Campo, Carlos Henrique Pagno, Tania Maria Haas Costa, Alessandro de Oliveira Rios and Simone Hickmann Flôres........... 100

Fractographic and rheological characterizations of CF/PP‑PE-copolymer composites tested in tensile Paula Helena da Silva Cirilo, Clara Leal Nogueira, Jane Maria Faulstich de Paiva, Lilia Müller Guerrini, Geraldo Maurício Cândido and Mirabel Cerqueira Rezende......................................................................................................................... 108

Evaluation of silanes in SBR 1502 / Telinne Monspessulana flour composites Oscar Buitrago, Oscar Palacio and Emilio Delgado...................................................................................................................................... 116

Reprocessability of PHB in extrusion: ATR-FTIR, tensile tests and thermal studies Leonardo Fábio Rivas, Suzan Aline Casarin, Neymara Cavalcante Nepomuceno, Marie Isabele Alencar, José Augusto Marcondes Agnelli, Eliton Souto de Medeiros, Alcides de Oliveira Wanderley Neto, Maurício Pinheiro de Oliveira, Antônio Marcos de Medeiros and Amélia Severino Ferreira e Santos......................................................... 122

Influence of microcrystalline cellulose in thermoplastic starch/polyester blown films Mônica Oliveira Reis, Juliana Bonametti Olivato, Juliano Zanela, Fábio Yamashita and Maria Victoria Eiras Grossmann........................ 129

Application of polyester derived from biomass in petroleum asphalt cement Fernando de Araújo, Ingrid Souza Vieira da Silva and Daniel Pasquini........................................................................................................ 136

Investigating the influence of conduit residues on polyurethane plates Rachel Faverzani Magnago, Nicolli Dayane Müller, Mayara Martins, Heloisa Regina Turatti Silva, Paola Egert and Luciano Silva........................................................................................................................................................................ 141

Extraction and properties of starches from the non‑traditional vegetables Yam and Taro Luan Alberto Andrade, Natália Alves Barbosa and Joelma Pereira............................................................................................................... 151

Whey protein-based films incorporated with oregano essential oil Sandra Prestes Lessa Fernandes Oliveira, Larissa Canhadas Bertan, Christiane Maciel Vasconcellos Barros De Rensis, Ana Paula Bilck and Priscila Cristina Bizam Vianna...................................................................................................................................... 158

Gamma radiation effect on sisal / polyurethane composites without coupling agents Marina Cardoso Vasco, Salvador Claro Neto, Eduardo Mauro Nascimento and Elaine Azevedo................................................................. 165

Understanding the water uptake in F-161 glass-epoxy composites using the techniques of luminescence spectroscopy and FT-NIR Rita Sales, Gilmar Thim and Deborah Brunelli.............................................................................................................................................. 171

Cover: Details of the fracture topography of CF/PP-PE (a)… e SEM micrographs of freeze-fractured surfaces of PU (A…).

Arts by Editora Cubo.

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Polímeros, 27(2), 2017

November Plastimagen Date: November 7-10, 2017 Location: Mexico City - Mexico Website: www.plastimagen.com.mx PLASTEC Minneapolis Date: November 8-9, 2017 Location: Minneapolis - USA Website: plastecminn.plasticstoday.com

PLASTEC West Date: February 6–8, 2018 Location: Anaheim – USA Website: plastecwest.plasticstoday.com 10th International Plastics Exhibition, Conference & Convention (PLASTINDIA 2018) Date: February 7–12, 2018 Location: Gujarat – India Website: www.plastindia.org/plastindia-2018/index.html

2nd International Conference and Exhibition on Polymer Chemistry Date: November 15-17, 2017 Location: San Antonio - USA Website: polymer.conferenceseries.com

March

30th International Plastics & Rubber Machinery, Processing & Materials Exhibition Date: November 15-18, 2017 Location: Jakarta - Indonesia Website: www.plasticsandrubberindonesia.com

Conductive Plastics Date: March 20–21, 2018 Location: Pittsburgh - USA Website: www.amiplastics.com/events/event?Code=C892

9th International Exhibition for Plastics Industry Date: November 22-24, 2017 Location: Almaty - Kazakhstan Website: plastworld.kz

December Polymers in Flooring Date: December 5-6, 2017 Location: Berlin - Germany Website: www.amiplastics.com/events/event?Code=C849 Fire Resistance in Plastics Date: December 5-7, 2017 Location: Cologne - Germany Website: www.amiplastics.com/events/event?Code=C847 33rd Annual Meeting of the Polymer Processing Society (PPS-33) Date: December 10-14, 2017 Location: Cancun - Mexico Website: pps-33.com Polymers in Footwear Date: December 11-12, 2017 Location: Duesseldorf - Germany Website: www.amiplastics.com/events/event?Code=C867

January 21st Thermoplastic Concentrates Date: January 23-25, 2018 Location: Coral Springs - USA Website: www.amiplastics.com/events/event?Code=C852 21st International Trade Fair Plastics and Rubber (INTERPLASTICA 2018) Date: January 23-26, 2018 Location: Moscow - Russia Website: www.interplastica.de Polyethylene Films Date: January 30 - February 1, 2018 Location: Coral Springs - USA Website: www.amiplastics.com/events/event?Code=C853

February Salone SamuPlast Date: February 1–3, 2018 Location: Pordenone – Italy Website: www.samuexpo.com/samuplast

Plastics Regulations Date: March 14–15, 2018 Location: Cologne - Germany Website: www.amiplastics.com/events/event?Code=C874

Polymers: Design, Function and Application Date: March 22–23, 2018 Location: Barcelona - Spain Website: sciforum.net/conference/polymers-2018 Plástico Brasil Date: March 25–29, 2018 Location: São Paulo - SP Website: www.plasticobrasil.com.br

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

May The International Plastic Showcase Date: May 7–11, 2018 Location: Orlando - USA Website: www.npe.org

June Polymers and Organic Chemistry (POC 2018) Date: June 4–7, 2018 Location: Montpellier - France Website: iupac.org/event/polymers-organic-chemistry-2018poc-2018 PLASTEC East Date: June 12-14, 2018 Location: New York – USA Website: plastec-east.plasticstoday.com/

Polímeros, 27(2), 2017 E3

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

ABPol Associates Sponsoring Partners

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

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PolĂ­meros, 27(2), 2017

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

Aditive Plásticos Ltda.

Avamplas – Polímeros da Amazônia Ltda.

CBE – Grupo Unigel Colorfix Itamaster Indústria de Masterbatches Ltda.

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

Formax Quimiplan Componentes para Calçados Ltda.

Imp. e Export. de Medidores Polimate Ltda.

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

Master Polymers Ltda.

Milliken do Brasil Comércio Ltda.

MMS-SP Indústria e Comércio de Plásticos Ltda.

Nexo International Ltda.

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

QP - Químicos e Plásticos Ltda.

Radici Plastics Ltda.

Replas Comércio de Termoplásticos Ltda.

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

Polímeros, 27(2), 2017 E5

New polymer can undulate and propel itself forward under the influence of light

N N N N N N N N N N N N N

Scientists at Eindhoven University of Technology and Kent State University have developed a new material that can undulate and therefore propel itself forward under the influence of light. To this end, they clamp a strip of this polymer material in a rectangular frame.

When illuminated it goes for a walk all on its own. This small device, the size of a paperclip, is the world’s first machine to convert light directly into walking, simply using one fixed light source. The researchers publish their findings on 29 June in the scientific journal Nature.

The maximum speed is equivalent to that of a caterpillar, about half a centimeter per second. The researchers think it can be used to transport small items in hard-to-reach places or to keep the surface of solar cells clean. They placed grains of sand on the strip and these were removed by the undulating movement.

The mechanism is so powerful that the strip can even transport an object that is much bigger and heavier than the device itself, uphill.

The motion of the new material is due to the fact that one side contracts in reaction to light, and the other one expands, causing it to bulge when illuminated.

That deformation disappears instantaneously once the light is gone. Although the material looks transparent to the human eye, it fully absorbs the violet light the researchers used, thus creating a shadow behind it.

The scientific team, led by professor Dick Broer of Eindhoven University of Technology, was able to create a continual undulating movement, using this ‘self-shadowing’ effect. They attached a strip of the material in a frame shorter than the strip itself, causing it to bulge. Then they shone a concentrated led light on it, from in front. The part of the strip that is in the light, starts to bulge downward, creating a ‘dent’ in the strip.

As a consequence, the next part of the strip comes in the light and starts to deform. This way the ‘dent’ moves backwards, creating a continual undulating movement.

This sets the device in motion, walking away from the light. When the device is placed upside down, the wave travels in the opposite direction, causing it to walk towards the light.

The research team managed to reach this specific behavior of the material using ‘liquid crystals’ (familiar

E6

in liquid crystal displays; lcd’s). The principle relies on the incorporation of a fast responding light-sensitive variant in a liquid crystalline polymer network. They engineered a material in such a way that this response is translated to an instantaneous deformation of the strip when illuminated, and relaxation directly when the light is gone.

Source: Plastemart - www.plastemart.com

Sabic Wins the Best Polymer Producers Award 2017 for Polypropylene SABIC has been named a winner of the prestigious Best Polymer Producer Award in the category of Polypropylene at an award ceremony held in conjunction with the European Plastics Converters (EuPC) Annual Event on 1 June 2017 in Madrid, Spain.

The winner is selected based on European-wide customer satisfaction survey conducted by the Polymers for Europe Alliance. Starting in February 2017 for in total three months, all users of polymers in Europe rated their suppliers’ performances by accessing the free of charge voting tool in strict confidence from mid last year up to end of May 2017. Companies are evaluated regarding the five criteria Polymer Quality, Regulatory Compliance, Delivery Reliability, Communication and Innovation.

The award is an affirmation of SABIC’s on-going efforts in plastics value chain through our customer engagement, commitment and investments in innovation and technology to provide solutions for the industry.

“SABIC is honored to be named The Best Polymer Producer winner for 2017. The award is tremendous validation that our commitment to the plastics industry – which to us, means listening to our customers and leveraging our global innovation expertise in response to their needs – is helping to solve their challenges,” said Sjoerd Zuidema, Director of Polypropylene Business Europe for SABIC. “The Best Polymers for Europe Awards represent a major opportunity for re-establishing a constructive dialogue and a good communication between suppliers and users of polymers in Europe,” added Ron Marsh, Chairman of the Polymers for Europe Alliance.

Source: SABIC - www.sabic.com

Polímeros, 27(2), 2017

http://dx.doi.org/10.1590/0104-1428.1959

Mixing process influence on thermal and rheological properties of NBR/SiO2 from rice husk ash Ana Maria Furtado de Sousa1*, Augusto Cesar de Carvalho Peres2, Cristina Russi Guimarães Furtado1 and Leila Lea Yuan Visconte3 Instituto de Química, Universidade do Estado do Rio de Janeiro – UERJ, Rio de Janeiro, RJ, Brazil Centro de Pesquisas e Desenvolvimento Leopoldo Américo Miguez de Mello, Petróleo Brasileiro SA, Rio de Janeiro, RJ, Brazil 3 Instituto de Macromoléculas Professora Eloisa Mano, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brazil 1

2

*ana.furtado.sousa@gmail.com

Abstract Silica was extracted from rice husk ash (RHA) by a sequence of reactions to produce nanosilica. Two laboratory routes, co-coagulation and spray drying, were used to incorporate the nanosilica into the rubber matrix. Samples were characterized regarding filler incorporation efficiency, thermal stability, rheological behavior and morphology.

Thermogravimetric analysis showed that spray-drying was the most efficient filler incorporation process and also the presence of silica increased the thermal resistance of the rubber compound when compared to the unfilled rubber.

The rheological behavior showed that NBR filled with silica presented higher elastic torque (S’), storage modulus (G’) and complex viscosity (η*) than unfilled rubber. The Payne effect was also observed for the composites produced by spray-drying. In addition, the thermal behavior and Payne effect results were supported by the comparison of morphology observed by FEG-SEM analysis.

Keywords: co-coagulation, latex, silica, spray-drying.

1. Introduction Silica has been used in many applications, including electronics, ceramics and polymer materials. The incorporation of well dispersed silica particles into polymer matrixes is a particularly important way to improve the thermal and mechanical properties of polymers. Furthermore, it is well known that rheological properties can change when fillers are incorporated in a polymer matrix, depending on the kind of filler and the way that it is incorporated. Thus, this knowledge is important to tailor polymer composites with potential applications in numerous technological fields[1-5].

Specifically in the rubber industry, silica is one of the reinforcing fillers widely used in compounds. It is alternative reinforcing filler used to improve some of the properties, such as to impart reduced heat buildup and high tear strength, tensile strength and abrasion resistance. However, a problem with conventional silica-reinforced rubber is that it is normally highly aggregated due to filler-to-filler interactions, resulting in a lower dispersion within the rubbery matrix and poor mechanical properties. The silicas normally used in rubber compounds can be of the precipitated or pyrogenic type, and are normally added to the compounds by open mixers or in internal mixers[6,7].

Recent studies have been devoted to mixtures of silica with rubber via a latex system, as an interesting incorporation option. In these processes, silica suspensions are prepared from commercial precipitated and pyrogenic silicas and then are incorporated in rubber by the co-coagulation method.

Examples of recent articles published in this respect are: obtainment of blends of epoxidized natural rubber (ENR),

Polímeros, 27(2), 93-99, 2017

styrene-butadiene (SBR) copolymer and silica[8]; investigation of the effect of pH and mixing time on the incorporation of silica in SBR rubber[9]; and preparation of masterbatches of NR and silica[10]. Besides the co-coagulation method, another technique for blending silica in rubber that also uses latex is the in situ generation of silica through the sol-gel reaction using tetraethoxysilane (TEOS). This method has proved to be highly efficient experimentally for dispersion of silica through reduced formation of aggregates[11,12].

The purpose of this study was to investigate the incorporation of nitrile rubber (NBR) via a latex system with silica sol, employing two mixing methods, co-coagulation and spray drying. In order to verify the process’s influence, samples from each process were characterized for filler incorporating efficiency, thermal stability, rheological behavior and morphology. Furthermore, the procedure to prepare silica sol from rice husk ash (RHA) and its characterization is also presented.

2. Materials and Methods 2.1 Materials The acrylonitrile butadiene copolymer (NTL615) latex with 33% acrylonitrile and solid content of 25% was kindly supplied by Nitriflex SA. Rice husk ash (RHA) was supplied by Empresa Brasileira de Pesquisa Agropecuária (Embrapa). The inorganic content of RHA was 72% and the chemical composition obtained by X-ray fluorescence

93

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

Sousa, A. M. F., Peres, A. C. C., Furtado, C. R. G., & Visconte, L. L. Y. (XRF, Bruker-AXS, model S4) was 95.30 SiO2; 1.70 K2O; 1.70 Na2O; 0.91 P2O5; 0.74 CaO; 0.61 MgO; 0.20 MnO; 0.12 SO3; 0.10 Fe2O3 and 0.01 Rb2O (% by weight).

This RHA has a high amount of silica (95.33%) and a few impurities, which were predominantly K2O and Na2O. Cation exchange resin Amberlite-120 was supplied by Vetec Quimica Fina. Analytical grade sodium and potassium hydroxide, sulfuric acid and aluminum sulfate were used as received. Sulfuric acid and aluminum sulfate were used to prepare the coagulant bath solution.

2.3.1 Co-coagulation (CC)

2.2 Silica sol production

Volumes of the prepared colloidal silica/latex were atomized and dried in a Büchi B290 Mini Spray Dryer.

The parameters for the spray drying process were as follows: air inlet temperature of 200 °C; air outlet temperature of 110 °C; aspirator set at 80%; pumper at 25% and nozzle cleaner 6.

The silica sol used in this work was synthesized in our laboratory using rice husk ash (RHA) as raw material.

The procedure used can be divided into three parts. First, RHA (150g) was reacted with sodium hydroxide (0.8 mol L-1) at 100 °C to produce a sodium silicate solution with SiO2/Na2O mole ratio of ~3. Five laboratory-sized batches were prepared and at the end, all reaction products were mixed to give just one water glass solution, which was characterized for SiO2/Na2O mole ratio, silica content, pH and density. Second, the water glass solution was passed through an ion exchange column filled with Ambertile-120 to allow the sodium ions to be replaced by hydrogen ions at the exchange sites of the cation resin, thus giving rise to an aqueous solution of active silicic acid with pH between 2 and 3. Finally, the active silicic acid was added (flow rate of 10 mL.min-1) to a 10 mL of KOH solution (10%wt) previously heated to 60 °C, in order to perform the nucleation, polymerization and particle growth of silica sol. In this part of the procedure, eight laboratory-sized batches of 100 mL were produced and all products were characterized for silica sol particle size distribution, silica content and pH.

2.3 Nitrile rubber (NBR) filled silica production To obtain the nitrile rubber filled with silica, 500 mL of NBR latex were mixed with the necessary amount of silica sol so as to have 2.5 phr and 5.0 phr of silica in the composition. First, the previously calculated latex and silica sol amounts were mixed and stirred for 2 hours.

Next, the colloidal suspensions were kept for maturation at 23 °C for 24 hours, a process which could be performed without disturbing the colloidal stability since the pH values of NBR latex and silica sol were in the range of 10-11.

The maturation step was necessary to give some time for the different particles present in these colloidal systems to adjust and co-exist in the new environment. Additionally, it allows air to escape and some particles to migrate into the micelle. As an example, in dipping and foam process, the maturation period is usually 10-20 hours and 2-3 days, respectively[13].

After that, two different processes, co-coagulation (CC) and spray drying (SD) were used as to occlude the silica particles inside the rubber during the destabilization of the system. In order to analyze the effects of silica content on the rheological and thermal properties of the rubber-silica blends, both procedures were also carried out with the NBR latex without any silica.

94

The co-coagulation experiments were carried out by adding the mixed colloidal silica/latex, at constant flow rate (50 mL.min-1) under stirring (2400 RPM), to a coagulant bath solution. At the end of the coagulation process, the powdered NBR-silica was vigorously washed with distillated water until all residual acid was completely removed.

Finally, the particles were separated by filtration and dried in an air-circulating oven at 100 °C for 1 hour.

2.3.2 Spray-drying (SD)

2.4 Characterization Dynamic light scattering (DLS) studies were conducted using a Malvern Zetasizer Nano ZS instrument.

Thermogravimetric analyses (TGA) were recorded with a TA Instruments analyzer at a heating rate of 20 °C/min, under air atmosphere.

Fourier-transform infrared (FTIR) spectra were record with a Perkin-Elmer Spectrum One spectrometer with ATR Spectrum software V5.3.1.

The rheology of the silica-rubber particles was investigated using rubber process analyzer (RPA 2000, Alpha Technologies, USA). For each test, two samples from each experiment were tested with two consecutive sweeps, one immediately following the other. Part of this test protocol was based on ASTM D6204. The procedure was as follows: (i) The sample was placed on the lower die and the cavity was closed; (ii) a pre-conditioned step was performed at 5 Hz, 2.8º, 100 °C during 4 minutes; (iii) the first strain sweep was carried out at 5 Hz and 100 °C, by applying a sequence of strain steps from 1 to 1259% (i.e., strains from 0.07º to 89.9º); (iv) the second strain sweep was carried out using the same parameters as the first strain sweep.

The morphology of the silica particles in the NBR matrix was investigated by field emission gun-scanning electron microscopy (FEG-SEM) performed with a FEI QUANTA 450 microscope. Samples were coated with platinum.

2.5 Data analysis Data analysis was performed using the software Statistica 8 (StatSoft).

3. Results and Discussions 3.1 Silica sol preparation from RHA The sodium silicate (water glass) produced from RHA had silica content of 6.9 ± 0.2 g/100 mL; density of 1.072 ± 0.001 g/cm3; pH of 11.8 and SiO2/Na2O mole ratio of 2.8 ± 0.2. These data indicated that SiO2/Na2O mole ratio and silica content were close to the recommended values for the water glass properties to be used for silica sol Polímeros, 27(2), 93-99, 2017

Mixing process influence on thermal and rheological properties of NBR/SiO2 from rice husk ash production, which are: silica content of 2-6% and SiO2:Na2O mole ratio close to 3.1[3].

Table 1 shows the individual results of the each batch run of silica sol produced. According to DLS measurements, all experimental runs showed monodisperse particles of silica sol, whose mean diameters ranged, respectively, from 3.3 to 7.6nm and from 2.6 to 7.1 nm for volume and number distributions.

3.2 Nitrile rubber (NBR) filled silica production The actual silica content of each experiment was determined

by TGA analysis (three replicated runs for each experiment) as the percent residue at 950 °C. Additionally, TGA was also used to determine the initial degradation temperature by using the extrapolated onset temperature method.

The actual silica amount present in the rubber matrix and the initial degradation temperature (the confidence interval for the mean) for each experiment are shown in Table 2.

Regarding silica content, the co-coagulation process was less efficient in the filler incorporation than the spray drying, since the amount of silica incorporated in the first process was lower than the theoretical predefined values.

We also observed that the amounts of silica incorporated by the co-coagulation process were exactly the same, regardless the initial amount used. This outcome suggests that only particles which have been adsorbed during the maturation step remained in the rubber composition, and probably the particles that were dragged and trapped during the latex coagulation did not have enough adhesion with the rubber and they were probably removed during the washing step.

Concerning the high efficiency of the spray-drying process, the results were as expected since it is a closed process, thus no material loss usually happens. Furthermore, there is no washing step in this process. An important comment that

should be highlighted is the fact that CC process showed a maximum limit of silica incorporation considering the experimental parameters used in the present research.

Because of that, it is not possible to discuss the effect of silica content increase in the products of CC process. Additionally, it is not also possible to perform comparisons between CD and SD processes regarding specifically the effect of silica content. However, each process can be analyzed individually and the SD and CC process can be compared regarding the unfilled rubber characteristics.

As for initial degradation temperatures, presented in Table 2, two comments can be made: (a) the unfilled NBR samples started to undergo thermal degradation at the same temperature intervals; regardless the process used, and (b) for each process the presence of silica in the rubber matrix increased the initial degradation temperature of the samples. The improvement of the thermal resistance of the CC and SD filled compounds compared to their respective unfilled ones can be explained by the fact that filler absorbed heat energy and, consequently, a retard of the heat transfer from the filler to the rubber chains occurs. It is also important to highlight this phenomena depends on the filler dimension and on the degree of the filler dispersion in the rubber matrix[14-17].

Changes in rubber behavior caused by the processability history can be evaluated by elastic torque response (S’) at high strains in the nonlinear viscoelastic region[18]. The elastic torque response (S’) of the first (1a Sw) and the second (2a Sw) strain sweeps of the all experiments produced by co-coagulation (CC) and spray-drying (SD) routes are shown in Figure 1.

Regarding the CC process, it can be seen from Figure 1a that co-coagulation produced unfilled and filled rubber compounds with the same rheological curve pattern. Here, the S’ torque rises with increasing strain, reaches a maximum

Table 1. Properties of silica sol obtained from RHA. Batch number

pH

Silica g.mL-1

1 2 3 4 5 6 7 8 Mean

11.0 10.5 10.1 10.4 10.1 10.0 10.3 10.3 10.0 ± 0.3

0.055 0.050 0.052 0.062 0.071 0.073 0.072 0.054 0.061 ± 0.008

Distribution of Particle Diameter, nm Volume Number 4.7 3.4 5.6 5.0 7.6 7.1 3.3 2.6 4.4 4.0 4.0 3.6 4.3 3.6 3.8 3.1 4.7 ± 1.1 4.1 ± 1.4

Table 2. Silica content and initial degradation temperature for the products of CC and SD processes.

Experiment Code CC-NBR CC-1 CC-2 SD-NBR SD-1 SD-2

Polímeros, 27(2), 93-99, 2017

Silica content, phr Theoretical values 0.0 2.5 5.0 0.0 2.5 5.0

Experimental values 1.2 ± 0.3 1.2 ± 0.1 2.3 ± 0.2 5.0 ± 0.1

Initial degradation temperature, °C 369-378 385-394 384-394 369-377 378-387 377-386

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Sousa, A. M. F., Peres, A. C. C., Furtado, C. R. G., & Visconte, L. L. Y. torque in both first and second sweeps at high strains, with overlap of the curves. These results indicate that the high deformations and amount of work applied to the samples during the first strain sweep did not modify their rheological properties. Moreover, the presence of silica in the rubber increased the S’ torque, reaching higher values than those observed for the unfilled rubber. Since all rubber samples were uncured, those higher values of S’ torque observed show that the presence of silica is responsible for making the rubber compound samples more elastic when they are submitted to high strains. Moreover, as expected, there is no difference on the S’ torque the curves of CC-1 and CC-2 since the amount of silica presented in both compositions are practically the same, as previously reported in Table 1.

According to the rheological behavior of the materials produced by spray drying, as shown in Figure 1b, two different curve patterns were found. For the unfilled NBR sample, in the first strain sweep, the S’ torque increased, reached a maximum near 300% strain and after that point it started to decrease. In addition, for this sample, the first and second strain sweep curves are not superimposed, indicating that the first high strain sweep altered the rheological property of unfilled NBR. This can be attributed to a reduction of the chain entanglement during the test. On the other hand, for the silica-rubber samples (SD-1 and SD-2), the S’ torque continually rises with increasing strain, reaches a maximum torque in both first and second high strain sweeps, with overlapping curves. The change of the curve pattern as well as the higher values of S’ torque of silica-rubber samples when compared to the unfilled NBR obtained from SD process show a reinforcing behavior and indicate an interaction between rubber and silica.

By comparing the unfilled compounds, Figure 1a (CC-NBR) and 1b (SD-NBR), it can be verified that the process affected the rheological properties of unfilled NBR. That result can be attributed to the fact that spray drying is a closed process, so additives used in the polymerization tend to remain in the rubber, contributing to the S’ torque reduction at high deformations. In order to investigate the above supposition, FTIR was performed for unfilled NBR from the co-coagulation and spray-drying processes, as shown in Figure 2. Looking at the two spectra, a new band at 1685cm-1, corresponding to the carbonyl stretching vibration band of C=O, appears in the spray-drying spectrum, which can indicate the presence of the polymerization additives.

Another possible reason, which could explain the appearance of presence of the band at 1685cm-1 in FTIR is the occurrence of the NBR thermal degradation during the spraying-dryer process. However, this hypothesis was refuted because the initial degradation temperatures of unfilled NBR produced

by CC (369-378 °C) and SD (369-377 °C) are practically the same, independently of the process used.

Measurement of the modulus decrease with increasing strain amplitude by dynamic mechanical testing is commonly used to investigate the reinforcing mechanism occurring in filled rubber compounds[19-21]. Figure 3 shows the storage shear modulus (G’) versus the strain amplitude for the co-coagulation and spray-drying experiments. Regarding the influence of the silica presence on the G’, as expected for co-coagulation process, the CC-1 and CC-2 showed similar values of G’ since they have the same amount of silica. For spray-drying process, the curves showed a relationship with 96

the silica amount, which means that, the higher the silica content is, the higher the G’ is.

For all unfilled and filled compounds studied in this work, two notable results were found: (a) unfilled NBR samples showed lower modulus values than silica/NBR, regardless of the process used; and (b) CC-NBR showed lower modulus values than SD-NBR. In general, the increase in storage shear modulus due to the presence of silica in a rubbery matrix can be explained by

Figure 1. Elastic torque response (S’) versus shear strain amplitude of the first (1a Sw) and the second (2a Sw) strain sweep for all samples prepared by (a) co-coagulation (CC) and (b) spray-drying (SD).

Figure 2. FT-IR spectra of unfilled NBR produced by co‑coagulation and spray-drying.

Polímeros, 27(2), 93-99, 2017

Mixing process influence on thermal and rheological properties of NBR/SiO2 from rice husk ash the contribution of the hydrodynamic effect and filler-filler interactions. The influences of polymer network generated by vulcanization and filler-polymer interaction were not considered in this case, since all samples tested were uncured and no coupling agent was used. About hydrodynamic effect, the increase in modulus results from the fact that fillers form a rigid phase that cannot be deformed. As for the filler-filler interaction, known as the Payne effect, it is characterized by a modulus decrease at low strain caused by the breakdown of filler inter-aggregates[18-20]. For silica filled NBR compounds, it is commonly understood that filler-filler networking takes place easily due to the polarity differences between polymer and silica.

From Figure 3, it can be seen that the Payne effect was not as pronounced for the silica/NBR samples produced by co-coagulation, since no decay of G’ modulus was observed until near 40% of strain. That result may indicate that the silica particles did not form aggregates with sufficient size to be detected.

As for the spray-drying process, the non-linearity of G’ modulus was clearly verified, principally for uncured NBR with 5.0 phr of silica, indicating that by using the spray-drying process, the filler-filler interaction increased as silica content rose.

Regarding the process influence on the unfilled rubber storage modulus (G’), it can be attributed to the fact that during spray drying the elastomer chains had a higher tendency to undergo physical entanglements than occurred in the co-coagulation process.

Figure 4 shows the complex viscosity (η*) as a function of strain amplitude at 100 °C for all co-coagulation and spray-drying experiments. Regarding the influence of the silica presence on the η*, similar outcomes of the storage shear modulus were observed.

Concerning the SD process, it is verified that SD-NBR sample was more sensitive to the strain amplitude deformation.

The reduction of the viscosity of the compounds in function of deformation applied can be explained by the reduced chain entanglements, reduced filler-rubber interaction and also the destruction of agglomerates, which facilitated the movement of the chains. In the specific case of the SD-NBR sample, as previously explained, the lower value of the viscosity at higher strains, when compared to the respective filled compounds, can be caused by the presence of additives from the emulsion polymerization (emulsifier/soap and other ingredients usually introduced into the polymerization vessels) that act as plasticizers.

The silica morphology and its dispersion in the rubber matrix were verified by FEG-SEM. Figure 5 (scale bar of 400 μm) shows the photomicrographs of the CC-1 (Figure 5a) and SD-1 (Figure 5b). Regarding CC-process, it can be noted that few silica particles are visible for CC-1. About the SD process, it is possible to see the presence of many large silica particles. Both morphologies support the results of thermal behavior and Payne effect presented.

The evaluation of silica size in sample CC-1 is shown in Figure 6, which presents two photomicrographs with scale bar of 3 μm (magnification of 25,000x) and 500 nm (magnification of 100,000x). According to these images, a broad range of silica particle sizes occurred, from the macro to the nano dimensions.

Figure 3. Storage shear modulus versus shear strain amplitude of samples prepared by co-coagulation and spray-drying processes.

The silica incorporation in the rubber was confirmed by the presence of a band corresponding to the siliceous material in the FEG-SEM/EDX spectrum of SD-1, as shown

Figure 4. Complex viscosity (η*) versus shear strain amplitude of samples prepared by co-coagulation and spray-drying processes.

Polímeros, 27(2), 93-99, 2017

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Sousa, A. M. F., Peres, A. C. C., Furtado, C. R. G., & Visconte, L. L. Y.

Figure 5. FEG-SEM photomicrograph with scale bar of 400 μm (magnification of 200x) of (a) CC-1 and (b) SD-1.

Figure 6. FEG-SEM photomicrograph of CC-1 with scale bar of (a) 3 μm and and (b) 500 nm.

in Figure 7. The presence of potassium is explained by the fact that it is a counter-ion of the colloidal silica particle.

4. Conclusions

Figure 7. FEG-SEM photomicrograph of SD-1 and the EDX spectrum of particle #3.

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Nitrile rubber filled silica from rice husk ash (RHA) was prepared by using co-coagulation and spray-drying processes.

Regarding the actual silica content of the compounds, under the experimental conditions used, only the spray drying was efficient concerning the silica content incorporation.

According to the thermal degradation temperatures, an improvement in thermal behavior of nitrile rubber occurred due to the presence of silica in rubber matrices.

About rheological behavior, all filled and unfilled rubber particles in the samples produced by co-coagulation exhibited the same rheological curve pattern and the presence of silica in the rubber matrix increased the S’ torque.

For the spray-drying process, two different S’ torque curve patterns were observed, being that difference ascribed to the by the presence of additives used in the polymerization and silica-rubber interaction. Also, the higher values of the storage modulus (G’) and Complex viscosity (η*) observed for unfilled spray-dried product was attributed to the physical entanglements produced during the process. In addition, Polímeros, 27(2), 93-99, 2017

Mixing process influence on thermal and rheological properties of NBR/SiO2 from rice husk ash the Payne effect increased as the silica content rose in the samples produced by spray drying.

FEG-SEM analysis showed that co-coagulation produces well silica dispersed compound with lower silica aggregate formation.

This study also demonstrated that RHA waste has potential as raw material to prepare silica sol, which is a nanomaterial with potential use in many applications.

5. Acknowledgements We thank the Instituto Nacional de Tecnologia (INT) for performing the FEG-SEM and Fundação de Amparo à Pesquisa do Estado do Rio de Janeiro (FAPERJ) for financial support.

6. References 1. Rocha, T. L. A., Jacobi, M. M., Samios, D., & Schuster, R. H. (2006). Evaluation of the influence of the polymer-filler interaction on compounds based on epoxidized elastomeric matrix and precipitated silica. Polímeros: Ciência e Tecnologia, 16(2), 111-115. http://dx.doi.org/10.1590/S0104-14282006000200010.

2. Pal, K., Rajasekar, R., Kang, D. J., Zhang, Z. X., Pal, S. K., Kim, J. K., & Das, C. K. (2010). Effect of fillers and nitrile blended PVC on natural rubber/high styrene rubber with nanosilica blends: Morphology and wear. Materials & Design, 31(1), 25-34. http://dx.doi.org/10.1016/j.matdes.2009.07.023.

3. Hassan, H. H., Ateia, E., Darwish, N. A., Halim, S. F., & Abd El-Aziz, A. K. (2012). Effect of filler concentration on the physico-mechanical properties of super abrasion furnace black and silica loaded styrene butadiene rubber. Materials & Design, 34, 533-540. http://dx.doi.org/10.1016/j.matdes.2011.05.005.

4. Taguet, A., Cassagnau, P., & Lopez-Cuesta, J. M. (2014).

Structuration, selective dispersion and compatibilizing effect of (nano)fillers in polymer blends. Progress in Polymer Science, 39(8), 1526-1563. http://dx.doi.org/10.1016/j. progpolymsci.2014.04.002.

5. Pourhossaini, M. R., & Razzaghi-Kashani, M. (2014). Effect of silica particle size on chain dynamics and frictional properties of styrene butadiene rubber nano and micro composites.

Polymer, 55(9), 2279-2284. http://dx.doi.org/10.1016/j. polymer.2014.03.026.

6. Leblanc, J. L. (2002). Rubber-filler interactions and rheological properties in filled compounds. Progress in Polymer Science, 27(4), 627-687. http://dx.doi.org/10.1016/S0079-6700(01)000405.

7. Zhang, C., Liu, L., Zhang, Z. X., Pal, K., & Kim, J. K. (2011).

Effect of silica and silicone oil on the mechanical and thermal properties of silicone rubber. Journal of Macromolecular Science, Part B: Physics, 50(6), 1144-1153. http://dx.doi.org /10.1080/08941920.2010.518533.

8. Pongdong, W., Nakason, C., Kummerlöwe, C., & Vennemann, N. (2015). Influence of filler from a renewable resource and silane coupling agent on the properties of epoxidized natural rubber vulcanizates. Journal of Chemistry, 2015, 1-15. http:// dx.doi.org/10.1155/2015/796459.

9. Wang, J., & Wu, Y. (2014). Preparation of silica-reinforced styrene–butadiene rubber via co-coagulation process. Journal

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of Elastomers and Plastics, 46(2), 144-155. http://dx.doi. org/10.1177/0095244312465277.

10. Prasertsri, S., & Rattanasom, N. (2012). Fumed and precipitated silica reinforced natural rubber composites prepared from latex system: mechanical and dynamic properties. Polymer Testing, 31(5), 593-605. http://dx.doi.org/10.1016/j. polymertesting.2012.03.003.

11. Sittiphan, T., Prasassarakich, P., & Poompradub, S. (2014). Styrene grafted natural rubber reinforced by in situ silica generated via sol–gel technique. Materials Science and Engineering B, 181, 39-45. http://dx.doi.org/10.1016/j.mseb.2013.11.018.

12. Lin, J., Wu, X., Zheng, C., Zhang, P., Huang, B., Guo, N., & Jin, L. Y. (2014). Synthesis and properties of epoxy-polyurethane/ silica nanocomposites by a novel sol method and in-situ solution polymerization route. Applied Surface Science, 303, 67-75. http://dx.doi.org/10.1016/j.apsusc.2014.02.075.

13. NPCS Board of Consultants & Engineers. (2010). The complete book on rubber processing and compounding. India: Asia Pacific Business Press.

14. Rao, Y. Q., Munro, J., Ge, S., & Garcia-Meitin, E. (2014).

PU elastomers comprising spherical nanosilicas: Balancing rheology and properties. Polymer, 55(23), 6076-6084. http:// dx.doi.org/10.1016/j.polymer.2014.09.065.

15. Šupová, M., Martynková, G. S., & Barabaszová, K. (2011).

Effect of Nanofillers Dispersion in Polymer Matrices: A Review. Science of Advanced Materials, 3(1), 1-25. http:// dx.doi.org/10.1166/sam.2011.1136.

16. Mokhothu, T. H., Luyt, A. S., & Messori, M. (2014).

Reinforcement of EPDM rubber with in situ generated silica particles in the presence of a coupling agent via a sol – gel route. Polymer Testing, 33, 97-106. http://dx.doi.org/10.1016/j. polymertesting.2013.11.009.

17. Poompradub, S., Thirakulrati, M., & Prasassarakich, P. (2014).

In situ generated silica in natural rubber latex via the sol–gel technique and properties of the silica rubber composites.

Materials Chemistry and Physics, 144(1-2), 122-131. http:// dx.doi.org/10.1016/j.matchemphys.2013.12.030.

18. Dick, J. S., Harmon, C., & Vare, A. (1999). Quality assurance of natural rubber using the rubber process analyzer. Polymer Testing, 18(5), 327-362. http://dx.doi.org/10.1016/S01429418(98)00026-9.

19. Fröhlich, J., Niedermeier, L. H. D., & Luginsland, H. D. (2005).

The effect of filler–filler and filler–elastomer interaction on rubber reinforcement. Composites. Part A, Applied Science and Manufacturing, 36(4), 449-460. http://dx.doi.org/10.1016/j. compositesa.2004.10.004.

20. Bezerra, F. O., Nunes, R. C. R., Gomes, A. S., Oliveira, M. G., & Ito, E. N. (2013). Efeito Payne em nanocompósitos de nbr com montmorilonita organofílica. Polímeros: Ciência e Tecnologia, 23(2), 223-228. http://dx.doi.org/10.4322/S010414282013005000022.

21. Ramier, J., Gauthier, C., Chazeau, L., Stelandre, L., & Guy, L. (2007). Payne effect in silica-filled styrene-butadiene rubber: Influence of surface treatment. Journal of Polymer Science. Part B, Polymer Physics, 45(3), 286-298. http://dx.doi.org/10.1002/ polb.21033.

Received: Nov. 14, 2014 Revised: Feb. 16, 2016 Accepted: Apr. 28, 2016

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http://dx.doi.org/10.1590/0104-1428.2371

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

Gelatin capsule waste: new source of protein to develop a biodegradable film Camila de Campo1, Carlos Henrique Pagno1, Tania Maria Haas Costa1,2, Alessandro de Oliveira Rios1 and Simone Hickmann Flôres1* Laboratório de Compostos Bioativos, Instituto de Ciência e Tecnologia de Alimentos, Universidade Federal do Rio Grande do Sul – UFRGS, Porto Alegre, RS, Brazil 2 Instituto de Química, Universidade Federal do Rio Grande do Sul – UFRGS, Porto Alegre, RS, Brazil 1

*simone.flores@ufrgs.br

Abstract This work aimed to develop biodegradable films utilizing a new source of gelatin derived from the nutraceutical capsule manufacture waste of coconut with safflower oil, coconut oil and safflower oil. The mechanical, physicochemical, barrier, optical, biodegradation, thermal and morphological properties were evaluated. All films showed low water vapor permeability, intermediate water solubility and high elongation at break. In addition, the films exhibited excellent barrier ability to ultraviolet light. After 15 days of soil burial degradation, the films lost over 68% of initial weight. Scanning electron microscopy showed an appearance free of pores, cracks or bubbles. Furthermore the films showed similar characteristics independent of the waste utilized. The results demonstrated that all the biodegradable films prepared presented appropriate characteristics to be used as substitute to synthetic packaging.

Keywords: gelatin films, food packaging, nutraceutical capsules waste, environmental impact.

1. Introduction The development of biodegradable film is an environmentally friendly technology that would permit a reduction in the impact and costs associated with polymers from non-renewable sources. There is growing interest, mainly from industry, regard to the preparation of packages that promote a greater product shelf life and having an environmental appeal, and there is reciprocal consumer interest in purchasing this type of product[1].

One of the alternatives for the biodegradable films production is the use of new materials. Polymers obtained from renewable resources or valorization of industrial wastes and by-products are considered a good alternative to reduce environmental impacts and costs[2]. The wastes may contain many substances of high value and if appropriate technology is employed, these materials can be converted into commercial products or raw materials for secondary processes. Numerous substances related to the food production process are suitable for separation and reuse[3], such as chitosan (derived from marine wastes) and soy by-products from soy oil industry[4].

An interesting waste to produce biodegradable films is obtained from the manufacture of nutraceutical capsules.

These wastes present a good source for the preparation of biodegradable films with desirable characteristics due to its composition (glycerol, gelatin and water). The production of nutraceutical capsules generates a large amount of waste that cannot be reused for the industry and is disposed of into the environment. Its use in other applications is of great interest to the nutraceutical industry because it would significantly reduce environmental problems as well as costs generated to its treatment. The gelatin waste from the manufacture of

100

nutraceutical capsules can be considered a new source of protein and has a high glycerol content naturally present in its composition. The glycerol provides better mechanical properties, forming films more stretchable and flexible[4], due to the reduction of intermolecular forces and increase on the mobility of polymeric chains. Vanin et al.[5] observed that glycerol was compatible with gelatin, and exhibited the highest plasticizing effect on the mechanical properties, producing a flexible film without phase separation.

Considering the possibility of use gelatin capsule waste as a promising material to produce biodegradable films and as an alternative for the reuse of this waste, the study aimed to develop a novel biodegradable film formulated from nutraceutical capsule manufacture waste of coconut with safflower oil, coconut oil and safflower oil and to characterize mechanical, physicochemical, barrier, optical, biodegradation, thermal and morphological properties, to develop a more sustainable material, adding value to this waste.

2. Materials and Methods The material used in this work was gelatin nutraceutical capsules waste provided by the Laboratory Chemical Pharmaceutical Tiaraju, located in Santo Angelo – RS. The gelatin used to produce the capsules was from bovine source. The wastes are basically composed of water (30%), glycerin (21.8%) and gelatin (48.2%). For the film preparation, the waste derived from the production of coconut with safflower oil, coconut and safflower nutraceutical capsules were utilized.

Polímeros, 27(2), 100-107, 2017

Gelatin capsule waste: new source of protein to develop a biodegradable film 2.1 Film preparation Gelatin films were prepared by casting. For the preparation of the filmogenic solution, 50 g of waste was dissolved in 70 mL of distilled water (conditions defined according to preliminary tests) in a water bath (Model 752A, Mark Fisatom) under constant stirring to melt the network at 60 °C for 30 minutes. The filmogenic solution was then placed in a vacuum desiccator for 2 minutes to remove air bubbles.

Then, 0.13 g/cm2 (corresponding to 20 g of the filmogenic solution), was weighed, placed in polystyrene Petri plates and dried in an oven with air circulation (Model B5AFD, Mark DeLeo) at 35 °C for 18 hours.

2.2 Film characterization The films were conditioned in desiccators under a controlled relative humidity of 58% at 25 °C, containing a saturated solution of sodium bromide (NaBr) for 48 h until their characterization[6].

2.2.1 Film thickness measurement The film thickness was determined using a digital micrometer (Model MDC-25, Mitutoyo Corp. Tokyo, Japan) with a range of 0 to 25 mm and a precision of 0.001 mm.

The values represent the average of five measurements made randomly along each sample evaluated and the final thickness is the arithmetic average of five points of each random sample.

the resulting material was weighed for the determination of the final dry weight (WF). The solubility was calculated

by Equation 1: WS = (%)

Wi − Wf ×100 Wi

(1)

where wi is the initial dry weight of the sample (g), and wf is the final dry weight of the sample (g).

2.2.5 Water Vapor Permeability (WVP) The WVP was determined gravimetrically, according to the method described by Mei et al.[10] with some modifications.

The samples were placed in permeation cells (inner diameter: 63 mm, height: 25 mm), filled with granular anhydrous calcium chloride and hermetically sealed. The permeation cells were placed in a glass chamber with a saturated sodium chloride solution, providing RH gradients of 75% at 25 °C. Mass gain was determined by weighing the permeation cells on an analytical balance (AY 220, Shimadzu) at intervals at 1 h, 12 h and 24 h. The water vapor permeability of the samples was determined in triplicate by Equation 2: WVP =

W .L

A.t.∆P

(2)

where W is the weight of water permeated through the film (g), L is the film thickness (m), A is the permeation area (m2), t is the time of permeation (h), and Δp is the water vapor pressure difference between the two sides of the film (Pa).

2.2.2 Mechanical properties

2.2.6 Opacity

The mechanical properties of films were evaluated by tensile strength (TS) [MPa], elongation at break (E) [%] using a Texture analyzer (Model TA.XT2i, Mark Stable Micro Systems, UK) with a load cell of 5 kg, using the A/TGT self-tightening roller grips fixture, according to ASTM[7]. Film specimen strips (80-25 mm) were cut and their thickness was measured using a micrometer at three random positions along each strip. Ten strips were cutted, and each one was held between the grips of the equipment for testing with the initial distance between the grip and test speed set to 50 mm and 0.8 mm s-1, respectively.

The opacity was determined by measuring the film absorbance at 210 and 500 nm using a UV spectrophotometer (model Shimadzu UV-1800). Films were cut into a rectangle piece and directly placed in a spectrophotometer test cell.

An empty test cell was used as reference. The film opacity was calculated dividing the absorbance values (nm) by the film thickness (mm). All determinations were performed in triplicate[11].

2.2.3 Moisture content Moisture content was determined according to Liu et al.[8].

The prepared film samples (2 cm in diameter) were dried in an oven (Model B5 AFD, Mark DeLeo) at 105 °C for 24 h, and their moisture content was analyzed gravimetrically.

2.2.4 Water solubility The water solubility was performed according to Colla et al.[9], with some modifications. The solubility was calculated as the percentage of dry matter of the film solubilized after immersion for 24 h in water at 25 °C. Discs of the film (2-cm diameter) were cut, weighed, immersed in 30 mL of distilled water, and slowly and periodically agitated. The amount of dry matter of the initial and final samples was determined by drying the samples at 105 °C for 24 h. Afterwards, the samples were filtered using desiccated pre‑weighed filter paper. The filter paper, containing undissolved fragments of film, was dried at 105 °C for 24 h in an oven (Mark DeLeo, model TLK 48, Brazil), and Polímeros, 27(2), 100-107, 2017

2.2.7 Thermal properties The gelatin film samples were submitted to thermogravimetric analysis (TGA), using nitrogen atmosphere, according to the methodology described by Tongnuanchan et al.[12], with some modifications. The equipment used was a Shimadzu model TGA-50. The samples were heated from room temperature to 650 °C at a rate of 10 °C min-1.

2.2.8 Morphological properties Surface and cross-section morphology of films were visualized using Scanning Electron Microscopy (SEM).

The dried film samples were mounted on aluminum stubs with double-sided adhesive tape, coated with a thin layer of platinum, and observed on a Scanning Electron Microscope (Model JSM 5800) at an acceleration voltage of 5 kV with a magnification of 200 times to the original specimen size.

2.2.9 Biodegradability: Indoor soil burial degradation The determination of soil burial degradation was performed according to Martucci and Ruseckaite[13]. Natural organic soil was used as the degradation environment for the films, which was added in plastic boxes (6 cm × 6 cm × 6.5 cm).

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Campo, C., Pagno, C. H., Costa, T. M. H., Rios, A. O., & Flôres, S. H. The film samples were cut into rectangles (2 cm × 3 cm) and dried at 60 °C in an oven (model TLK48, DeLeo, Brazil) until constant weight (m0), and placed into an aluminum mesh, that were buried at the depth of 4 cm from the surface of the natural organic soil. Every 2 days, water was added to the soil to maintain the humidity. The degradation degree of the films was determined after 15 days as the weight loss (WL; %) by Equation 3: WL = (%)

mt − m0 ×100 m0

(3)

where m0 is the initial mass and mt the remaining dried mass after 15 days.

2.3 Statistical analysis All analyses were performed in triplicate, and the results were evaluated by an analysis of variance (ANOVA) and Tukey’s test at a significance level of 0.05 using the software Statistica 12.0 (Statsoft Inc., São Paulo, Brazil).

3. Results and Discussion The resulting film emulsion of the three gelatin capsules waste analyzed were composed of glycerin (15%), gelatin (28.3%) and water (56.7%). Gelatin has been known to form clear, flexible, strong and oxygen-impermeable films when cast from aqueous solutions in the presence of plasticizer.

Edible films with gelatin reduce oxygen, moisture, and oil migration and can carry antioxidants or antimicrobial agents[14]. Due to the hydrophilic nature of these films, they can be used as good gas barriers, but they have poor water barriers.

Al-Hassan and Norziah[15], analyzed fish skin gelatin used for the development of films and found 81.3% for soluble protein. They developed films with starch and different protein mixtures, and the films with higher content of protein presented 21,6% of protein, lower values than those obtained in this study. The high protein content induced to form films with higher elongation at break. Thus, the gelatin capsules waste based films showed higher capacity to form more elastic films. Proteins have good adherence to hydrophilic surfaces and serve as good barriers against O2 and CO2[16], thus, the higher amount of it helps to improve these characteristics.

and soy protein films and found values lower than all films in this study (0.107-0.132 mm). Cozmuta et al.[18], developed gelatin films with hemp and sage oils and found similar values to those obtained in this work (0.182 to 0.211 mm) and observed increase in the thickness for films with higher oil content added.

The thickness values observed in the present study, compared to previous works, indicate that the waste studied allowed the formation of films with the appropriate thickness to be mechanically resistant and suitable for food biodegradable packaging.

3.2 Mechanical properties Adequate mechanical strength and extensibility are necessary for films to have resistance to external factors and suitable barrier properties for applications such as food packaging[19]. Representative stress and strain curves of the analyzed films are shown in Figure 1. The curves demonstrated the deformation behavior of the films, indicating that the films elaborated with gelatin capsule waste of safflower exhibited the higher values for elongation at break and the lower values for tensile strength, while the film developed with gelatin capsule waste of coconut with safflower, showed the higher values for tensile strength.

Table 1 shows the thickness, tensile strength (TS) and percentage elongation at break (E%), of the analyzed films.

The values obtained for tensile strength (TS) are similar to those found by Hosseini et al.[20] who studied films prepared from fish skin gelatin in cold water (2.17 ± 0.97 MPa); however, the film elongation at break in this study is greater than that found by these authors (82.61±20.11%). Al-Hassan and Norziah[15], also found lower values than those obtained in this work when developing films from starch with fish

3.1 Film thickness As observed in Table 1, the thickness of the films did not present a significant difference independent of the waste utilized. Carvalho[17], evaluated the thickness of cassava starch

Figure 1. Stress-strain curves of nutraceutical capsule waste based films.

Table 1. Film thickness, tensile strength (TS), and elongation at break (E) of nutraceutical capsule waste based films.a,b Film Coconut with safflower Coconut Safflower

Thickness (mm) 0.256 ± 0.005a 0.212 ± 0.04a 0.274 ± 0.05a

TSc (MPa) 2.41 ± 0.09a 2.14 ± 0.10b 1.96 ± 0.09c

Ed (%) 264.62 ± 1.57b 189.10 ± 5.79c 275.68 ± 3.97a

The results are represented as the means ± standard deviation; bValues with the same letter are not significantly different (p > 0.05); TS (MPa) = F max/A (F max, maximum load (N) needed to pull the sample apart; A, cross-sectional area (m 2) of the samples); d EAB (%) = (E/50) x 100 (E, film elongation (mm) at the moment of rupture; 50, initial grip length (mm) of samples). a c

102

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Gelatin capsule waste: new source of protein to develop a biodegradable film skin gelatin: 1.28 MPa and 1.67 MPa for tensile strength and values between 84 and 102% for percentage of elongation at break. Garrido et al.[21], found higher values for tensile strength (1.55-7.50 MPa) for films produced from soya by‑products with different ratios of sorbitol, but found lower values for elongation at break (3.95-117.83%).

Comparing to commercial polymers such as LDPE (9-17 MPa) and polystyrene (35-55 MPa) the values obtained in this study for tensile strength were lower, however the majority of films presents TS values lower than the synthetic commercial polymers. On the other hand, nutraceutical capsule waste based films showed high elongation at break and the values were higher than cellophane (20%) and polystyrene (1%)[22].

Dias et al.[23], observed that the highest Young’s modulus values (1053 ± 146 MPa) were found for rice starch and flour-based films with lower plasticizer content and these films with sorbitol were more inflexible than those with glycerol. The glycerol content of waste acted as a plasticizer and may have reduced the interactions between polymer chains, thereby increasing film flexibility[24]. Glycerol can be used in biodegradable films to increase their flexibility and elongation values, however reduce the values of tensile strength[25]. The studied films presented high percent elongation at break, providing films with higher flexibility and that take a longer time to break.

In addition, hydrophobic materials, such as lipids present in the waste, can decrease the strength of films. Shellhammer and Krochta[26], observed that increasing the lipid level, the strength of whey protein isolate (WPI) films reduced.

The incorporation of Candelilla wax gave the weakest WPI films, followed by beeswax, milk fat, and carnauba wax.

These authors found that the lipid also affected the tensile strength, elongation and elastic modulus of the composite films. In this work, the lipid at low concentrations present in the waste is homogeneously distributed in the polymeric matrix, resulting in good mechanical properties.

hydrophobic characteristics of the oil, that cause a decrease in water-protein interactions.

3.4 Water solubility The water solubility defines the tolerance to water and is determined by the chemical structure of the materials.

The desired value for the solubility depends to its application or intended use.

As observed in Table 2, the solubility in water for gelatin films was approximately 45%. The high glycerol content of nutraceutical capsule waste interacts with the film matrix by increasing the space between the chains, facilitating water migration into the film and, consequently, increasing solubility[25]. Nur Hanani et al.[30], found similar values of 40% for films prepared with fish gelatin. Hosseini et al.[20], found an average of 64% of solubility for films from skin gelatin derived from cold water fish. In the analysis of blend films from soy protein isolate and cod gelatin, Denavi et al.[31], observed values above 80%, and justified that this value would indicate a poor water resistance, however, for some applications, the high solubility could be advantageous: for example, as a carrier of bioactive compounds, and soluble film packaging is convenient to use in ready-to-eat products as they melt in boiled water or in the consumer’s mouth.

The results of this study indicate that the films exhibit intermediate solubility in water and may be applied in dry foods.

3.5 Water Vapor Permeability (WVP)

3.3 Moisture content

The Table 2 shows the results of water vapor permeability of nutraceutical capsule waste based films. Hosseini et al.[20], studying fish gelatin films, found higher values for WVP (0.826 ± 0.047 g.mm/h m2 kPa). However, the results of this study were higher than synthetic films, such as high‑density polyethylene film (HDPE) (0.0012 g.mm/h m2 kPa) and polyester film (0.0091 g.mm/kPa h m2)[32]. Dias et al.[33] obtained similar values to those found in this work (0.21 to 0.24 g.mm/h m2 kPa) for films of pig hide gelatin and glycerol containing yucca extract and lecithin respectively.

The moisture contents of gelatin films are shown in Table 2, and the values were around 20-21%. There was no significant difference in the moisture content between the three different wastes utilized. Arfat et al.[27], obtained similar values (20-22%) for films based on fish protein isolate. Bodini et al.[28], obtained lower values (14.7%) for films of fish gelatin. Increasing the glycerol amount, increases the film moisture content[29]. The oil present in waste may have caused a decrease in the film moisture because the

The known values of the water vapor permeability are essential for defining the possible film applications.

A polymer that is very permeable to water vapor may be suitable for fresh products packaging, whereas a slightly permeable film may be useful for the dehydrated products packaging[34].

The protein–protein and protein–lipid interactions forming the film matrix allowed nutraceutical capsule waste-based films to present adequate water vapor barrier properties for potential use as biodegradable packaging for dried foods.

Table 2. Water vapor permeability (WVP), water solubility, moisture content and opacity of nutraceutical capsule waste based films.a,b Waste Coconut with safflower Coconut Safflower

WVPc (g.mm/m2 h kPa) 0.17 ± 0.01a 0.22 ± 0.06a 0.21 ± 0.01a

Water solubility (%)

Moisture Content (%)

37.68 ± 1.08c 45.76 ± 4.81ab 46.23 ± 0.23a

21.36 ± 1.59a 20.44 ± 0.9a 20.07 ± 1.58a

Opacityd 210nm 59.01 ± 2.53a 62.20 ± 2.69a 62.43 ± 0.85a

500nm 2.68 ± 0.30a 1.63 ± 0.09b 2.29 ± 0.12b

The results are represented as the means ± standard deviation; bValues within each column with the same letter are not significantly different (p > 0.05); cWVP (g.mm/m2 h kPa) = (w, weight gain (g) of the cup; x, film thickness (m); A, area of exposed film (m2); t, time of gain (h); (P2 - P1), vapour pressure differential (Pa) across the film); dOpacity (%) at each wavelength; x, film thickness (mm). a

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Campo, C., Pagno, C. H., Costa, T. M. H., Rios, A. O., & Flôres, S. H. 3.6 Opacity properties The study of the UV light absorption capacity of the biodegradable films is important to determine their possible applications for food packaging. If these materials are able to absorb UV light, they could be used to package and extend the shelf life of fatty foods, which are susceptible to the oxidative degradation catalyzed by UV rays[35].

In the opacity, higher values indicate less transparency and high opacity[11]. According to the Table 2, when analyzed in the UV-visible (500 nm) region, the films showed low values, indicating that they had low absorption, which indicates greater transparency and lower opacity.

In the UV region at 210 nm, the films showed high values, indicating a high absorption, demonstrating that these films have a high ability to protect against UV radiation, which causes the oxidative deterioration of packaged foods, leading to nutrient losses, discoloration and off-flavors[36].

This result is in agreement with previous reports on gelatinbased films[2,37]. Both studies indicate that protein-based films are considered to have high UV barrier properties, owing to their high content of aromatic amino acids, which absorb UV light.

3.7 Thermal properties TGA results, showing the thermal degradation behavior of all films, are observed in Figure 2. The films showed a similar behavior considering the thermal properties studied.

Three main stages of weight loss were observed for all films.

The first stage was observed between 0 and 150 °C, where,

for both films, there was a loss of 20% by weight, which can be attributed to the moisture of the films. These results are in accordance with the moisture content presented before.

The second stage of weight loss appeared at the onset temperature of 150-200 °C most likely due to the degradation or decomposition of lower molecular weight protein fractions and glycerol compounds. Hoque et al.[38], also reported a degradation temperature in the range of 196-217 °C for cuttlefish skin gelatin film. For the third stage of weight loss (200-450 °C), there was a greater loss. This was possibly due to the decomposition of highly interacted proteins in the film matrix. The results indicated that the film degradation began at ≈ 200 °C. This result is higher than those found by Nuthong et al.[39], who reported that the initial temperature for the degradation of a porcine plasma protein-based film was observed at 170 °C. The results suggested that all the films studied showed high thermal resistance. Additionally, all films had residual mass at 650 °C, which represents a carbon residue of the the raw film decomposition.

3.8 Scanning Electron Microscopy (SEM) The SEM provides information about the film microstructure and the interactions between film components[40]. The scanning electron microscopy of the surface and cross section of the films is shown in Figure 3 and Figure 4 respectively.

The images obtained showed a surface without cracks, which suggests a cohesive matrix.

The obtained films were slightly yellow but still transparent and flexible. Their surfaces were without pores, cracks or bubbles and it was associated with the better mechanical and physical properties[23].

However, a rough surface with small particles was observed, that may be occurred due to aggregation phenomena of lipid droplets caused by the oil naturally present in the waste during the drying step, presenting irregularities that was visualized by crystals formed at the microscopic level.

Tongnuanchan et al.[12], studying films composed of fish skin gelatin, obtained films with a homogeneous appearance but also found films with a rough surface when essential oils were incorporated. Ma et al.[41], reported that the presence of olive oil led to the marked increases in the roughness of the films. The increase in the surface roughness is principally due to the adherence and formation of oil droplets[33].

Figure 2. Thermogravimetric Analysis (TGA) curves of nutraceutical capsules waste based films.

In the cross-section, some irregularities were also observed.

Interactions in protein-protein of the film matrix might be broken due to the oil present in the composition of the films,

Figure 3. Scanning electron microscopy (SEM) of the surface of nutraceutical capsule waste based films of (a) coconut with safflower; (b) safflower oil; and (c) coconut oil.

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Gelatin capsule waste: new source of protein to develop a biodegradable film

Figure 4. Scanning electron microscopy of the cross section of nutraceutical capsule waste based films of (a) coconut with safflower; (b) safflower oil; and (c) coconut oil.

Figure 5. Visual aspect of nutraceutical capsule waste based films prior and after 15 days of incubation in soil.

providing the roughness in the cross-section[42]. The oil droplets were more localized inside the film network, and on the macroscopic level those droplets was not perceptible.

3.9 Biodegradability: Indoor soil burial degradation Biodegradation study of nutraceutical capsules waste based films has been done by soil burial method, to reproduce the degradation conditions that happen in natural soil environment. The microbial population found in soil (bacteria, actinomycetes, fungi and protozoa) may act synergistically during degradation and can reproduce naturally occurring conditions[13]. The experiment was conducted for 15 days, when the films were almost totally degraded. After this period, the films could not be recovered and evaluated, due to the macroscopic deterioration observed. Pictures of the recovered samples before and after 15 days of exposure to soil burial are shown in Figure 5. The analyzed films showed similar visual aspect independent of the waste utilized, both prior and after 15 days to exposure to soil burial degradation.

The weight loss average of the films exposed to soil environment was considered as a degradation indicator.

After 15 days, all the films (independent of the waste utilized) reduced 68% of its initial weight. This lost was mostly attributed to the leaching of low-molar-mass compounds, such as glycerol. Degradation products of gelatin and glycerol might be eventually adsorbed by soil, being metabolized by microbes[43]. During soil burial the films also absorbed water, losing their initial shape.

The clear visual deterioration of the samples evidence the films susceptibility to degradation, and due this, they might be classified as rapidly degradable materials.

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4. Conclusion Nutraceutical capsule waste based films were successfully obtained, presenting a surface without bubbles or cracks.

The gelatin films presented good mechanical properties, with high elongation at break and showed low water vapor permeability and consequently can be applied in dry foods.

In addition, the films exhibited intermediate solubility in water and good absorption of ultraviolet radiation, which could provide increased protection to packaged food. The biodegradability test proved that the films are biodegradable in natural environmental conditions.

These results suggest a high potential for nutraceutical capsule manufacture waste to form biodegradable films with the appropriate characteristics for use in the food packaging industry, being an alternative to replacement non biodegradable packaging, decreasing the environmental impact. Moreover, the use of these wastes presents potential to contribute to the reduction of its amount sent to landfill, reducing environmental damage and costs, which is of great importance for industries and consumers. Further studies would be required to optimize the process condition, tensile strength and determine the specific use for films in commercial food systems.

5. Acknowledgements The authors are thankful to Laboratory Chemical Pharmaceutical Tiaraju, located in Santo Angelo – RS, for supplying raw material for this research, to Coordenação de Aperfeiçoamento de Pessoal de Nível Superior (CAPES) and Fundação de Amparo à Pesquisa no Estado do Rio 105

Campo, C., Pagno, C. H., Costa, T. M. H., Rios, A. O., & Flôres, S. H. Grande do Sul (FAPERGS) for the financial support.

The authors are also grateful to the Eletronic Microscopy Center of UFRGS, for their assistance in the use of SEM.

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Received: Nov. 09, 2015 Revised: May 18, 2016 Accepted: June 18, 2016

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http://dx.doi.org/10.1590/0104-1428.2355

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Fractographic and rheological characterizations of CF/PP‑PE-copolymer composites tested in tensile Paula Helena da Silva Cirilo1, Clara Leal Nogueira2, Jane Maria Faulstich de Paiva3, Lilia Müller Guerrini1, Geraldo Maurício Cândido1 and Mirabel Cerqueira Rezende1,2* Instituto de Ciência e Tecnologia, Universidade Federal de São Paulo – Unifesp, São José dos Campos, SP, Brazil 2 Divisão de Engenharia Aeronáutica e Mecânica, Instituto Tecnológico de Aeronáutica – ITA, São José dos Campos, SP, Brazil 3 Centro de Ciências em Gestão e Tecnologia, Universidade Federal de São Carlos – UFSCar, Sorocaba, SP, Brazil 1

*mirabelcr@gmail.com

Abstract This work shows the fractographic study of fractured surfaces resulted from tensile tests of thermoplastic composites based on poly(propylene-co-ethylene) (PP-PE) and modified PP-PE copolymers reinforced with continuous carbon fibers (CF). The PP-PE matrix was modified with two agents called AM1 (based on maleic anhydride) and AM2 (containing an elastomeric agent), respectively. Three different laminates - CF/PP-PE, CF/PP-PE(AM1) and CF/PP-PE(AM2) were manufactured. The best tensile strength and elastic modulus results were determined for the CF/PP-PE(AM1) laminate (507.6 ± 11.8 MPa and 54.7 ± 2.4 GPa, respectively). These results show that the AM1 agent contributed to increase the physicochemical interaction between the CF and the PP-PE matrix. This condition provided a better loading transfer from matrix to the reinforcement. Scanning electron microscopy analyses of the fracture surfaces show the fractographic aspects of the samples and allow evaluating the fiber/matrix-interfacial adhesion. Poor adhesion is observed for the CF/PP-PE and CF/PP-PE(AM2) laminates with the presence of fiber impressions on the polymer rich regions and fiber surfaces totally unprotected of polymer matrix. On the other side, a more consistent adhesion is observed for the CF/PP-PE(AM1) laminate. This result is in agreement with the tensile test data and show the presence of a good interaction between the laminate constituents. The correlation of the mechanical and fractographic results with the curves of complex viscosity versus temperature of the studied polymer matrices shows that the matrix viscosity did not affect the wettability of the reinforcement.

Keywords: fractography, thermoplastic composite, carbon fiber, PP-PE.

1. Introduction The technological advances and the tight requirements demanded by aerospace, marine, automotive and sports equipment industries have promoted the increase employment of structural polymeric composite reinforced with carbon fibers in these areas. This trend is attributable to the low density (1.4-1.6 g/cm3) of this class of materials associated with high values of both specific rigidity and mechanical resistance, characteristics that meet strict requirements in service. In addition, the processing of components with these materials is very versatile, capable of producing parts with large dimensions and complex shapes[1-5].

In this context, thermosetting composites have been established in a prominent place in the structural materials area[6-10]. But more recently, the study and use of thermoplastic composites reinforced with continuous fibers in academic and industrial segments have increased, considering their high structural performance and similar or superior mechanical properties to those obtained with thermosetting composites.

Beside this, they present high resistance to impact, better delamination resistance, and fracture toughness, greater resistance to environmental aging, they are non-flammable and can be stored for a long time at room temperature.

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These characteristics have motivated the manufacture of components with applications in aerospace, automotive, oil, gas and civil infrastructure[11-13].

In general, thermoplastic composites with continuous fibers are produced with less processing cycles, and can be hot pressed, thermoformed, pultruded, consolidated in autoclave or by automated methods on multiple complex shapes of large size[14-20]. Besides, they can be welded to other structures by electrical resistance, ultrasonic and induction technologies, for example[21-25].

Despite the advantages provided by polymeric composite materials, defects and damages may be present in the laminated structure, caused by different variables, which may occur during the processing of the component or during its useful life. Defects can happen due to design errors or originated in various stages of manufacture (for example, during the laminate stacking sequence). The damages can be resulted of impacts from inadequate conditions of transport and storage or due to difficulties in operation[26-31]. Defects and damages, even if not noticeable on a visual inspection, can contribute significantly to the reduction of the composite’s resistance. This situation can be aggravated when the

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Fractographic and rheological characterizations of CF/PP-PE-copolymer composites tested in tensile composite is exposed to harsh environmental conditions such as sudden changes in temperature, high relative humidity and ultraviolet radiation, which can affect its mechanical properties and consequently its integrity and durability[32-35].

Most of the structures manufactured in polymeric composites reinforced with continuous fibers are subjected to a diversity of loads. Accordingly, these structures are exposed to the occurrence of damages that may lead to the failure of material. Structural collapse may occur in different complicated shapes due to the anisotropic construction of the laminate, in which the possible failure modes to be developed are influenced by the orientation of fibers, number of layers, layer stacking sequence, load direction applied and environmental interactions.

The complete fracture of these materials typically occurs in three basic types, called: interlaminar, intralaminar and translaminar fractures, which generate fracture surfaces with different morphologies, consequently with different characteristics. When the fracture surfaces are properly interpreted, it is possible to identify the damage and failure processes. The fiber has influence on the matrix fracture process while the interfacial quality of fiber/matrix adhesion has a significant influence on the local where the failure starts. As consequence, the detailed fractographic analysis of fractured surfaces has fundamental importance on the qualitative evaluation of the processing-structure-property relationship of polymer composites. In this process it includes the analysis of resistance and loading conditions to which the material is subjected.

The fractography reveals the origin and the direction of crack propagation, the nature of loading that originated the crack, defects in material, failure mechanisms, and also determines the sequence of events of the failure and confirms or removes any suspicious on the failure modes present.

Moreover, the fractography presents itself as a powerful research tool that generates information that supports the improvement of materials processing and encourages communication between experimental and predictive areas. This knowledge can be used in the development and evaluation of theoretical models of behavior and growth of cracks[36-44]. Usually, the fractography is carried out with the aid of the scanning electron microscopy (SEM) technique.

The correct identification of fractographic aspects is not simple and immediate. However, if the fracture surface is accurately assessed, it is possible to identify and to analyze design data and composite processing parameters with potential to cause failures. This procedure is useful because it can be used to optimize the composite processing and also to prevent other damages that can occur in the future.

Thus, fractographic studies have been strongly disseminated and used not only in the area of metallic materials, but also in the polymeric composites reinforced with continuous fibers[45-53].

Due to the steady increase of thermoplastic composite materials in the manufacturing industry of structural components, this work aims to contribute to the fractographic area of impregnated carbon fiber laminates with polypropylene/ polyethylene (PP-PE) copolymer, fractured in tensile loading at room temperature. Correlation of the observed fractographic aspects with both the used processing technique and the individual characteristics of the laminate components is made. A brief rheological study of the used thermoplastic matrices is presented and correlated with the microscopic observations.

2. Materials and Methods 2.1 Materials In this study three different solid laminates were processed based on carbon fiber fabric (CF) with orientation (0°, 90°), style Plain Weave, with tows of 3000 filaments, from Hexcel Composites Co. This reinforcement was impregnated with thermoplastic matrices based on three PP-PE copolymer films, containing 7% (wt/wt) of ethylene monomer. Table 1 shows the specifications of the thermoplastic films used. Two of these films were modified by the manufacturer, Polibrasil Co., aiming to improve the mechanical behavior of the laminates reinforced with CF. For this, one of the PP-PE samples was modified with about 1% (wt/wt) of maleic anhydride. This sample was named PP-PE(AM1). The other sample was modified with 1% (wt/wt) of an elastomeric agent based on ethylene-octene copolymer, identified as AM2. Such elastomeric agent is used to improve the polymer flow and the impact resistance of PP-PE copolymers[54].

The processing of the thermoplastic laminates was based on the hot compression molding technique. The CF reinforcing layers and the polymeric films were stacked on the mold surface, alternating one layer of CF and two layers of the polymeric film, in respective order, totaling 15 layers of CF and films. This proportion corresponds in the final laminate at approximately 60 ± 1% by volume of CF. This value was determined by acid digestion, in triplicate, in accordance with ASTM D3171-11.

The hot compression molding was held in a hydraulic press, fitted with a mold of 400 mm × 400 mm. The heating rate was 2 °C/min up to the maximum temperature of 230 °C, holding at this temperature for 2 h. At this step, a pressure of 4 MPa was applied. The cooling was natural until to reach the room temperature after 8 h. This procedure ensured the consolidation of the laminates with a thickness of 3.0 ± 0.1 mm.

2.2 Characterization From the processed laminates, specimens were prepared for the tensile mechanical tests, in a longitudinal direction, according to ASTM D3039/D3039M-00. For this, it was used

Table 1. Thermoplastic matrices used.

Copolymer Matrices PP-PE (with 7% (wt/wt) of ethylene) PP-PE (with 7% (wt/wt) of ethylene) + agent AM1 PP-PE (with 7% (wt/wt) of ethylene) + agent AM2

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Code PP-PE PP-PE(AM1) PP-PE(AM2)

Melting Points (°C) 128/163 ± 0.5 126/164 ± 0.5 126/164 ± 0.5

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Cirilo, P. H. S., Nogueira, C. L., Paiva, J. M. F., Guerrini, L. M., Cândido, G. M., & Rezende, M. C. an universal testing machine MTS-810 model, at a constant speed of 2 mm/min at room temperature. The specimens presented dimensions of 250 mm × 25 mm × 3 mm thickness, with tabs of CF/epoxy resin bonded with epoxy resin adhesive film at their ends. This procedure was adopted to distribute uniformly the clamping forces of the specimen in the grips of the testing machine, to favor the smooth transfer of the load to the specimen and to protect the laminate surface from possible damages during the test.

After the tensile tests, the specimens had the fracture surfaces protected with a tape and cut with a diamond disc. The fracture surfaces were cleaned blowing dry air to remove any possible cutting debris. Then, the surfaces were coated with a gold film by sputtering process, making them conductive for the SEM analyses in a microscope model FEI INSPECT S50.

In order to evaluate the influence of viscosity of the matrices on the CF impregnation, rheological analyses of PP-PE films were carried out using a rheometer HAAKE model RS6000, adapted with parallel plate geometry and a gap of 1 mm. All films were heated up to 230 °C at a heating rate of 1 °C/min. Analyses were performed with tension of 500 Pa. For this, previous analyses to identify the linear viscoelastic regime were made. The maximum temperature for these analyses was determined previously from thermogravimetric analyses in air, where the degradation temperatures varied between 240-245 °C.

3. Results and Discussion When a laminate of polymeric matrix reinforced with continuous fibers is subjected to mechanical loading, the matrix begins to deform with the generation of strain on the fiber surface. As the reinforcement is sufficiently long, the load intensity necessary to lead them to fracture is higher, contributing to increase the composite strength. However, the conditions for an effective transfer of efforts from the matrix to the fibers depend on the nature of the fiber/matrix interface. Thus, the interface characteristics formed between the reinforcement and the polymer matrix have a great importance on the structural performance of the laminate, affecting the mechanical properties and the failure process of the composite.

AM1 agent contributed to have a more resistant laminate, about 15% higher than that determined for the CF/PP-PE. This behavior is attributed to the AM1 agent that improved the interaction between CF reinforcement and the polymer matrix. In this case, the AM1 agent contributed to increase the effort transference from the matrix to the fibers. On the other hand, the CF/PP-PE(AM2) laminate provides a lower mean value of tensile strength than that determined for the CF/PP-PE laminate, around 4% lower. This indicates that the AM2 agent did not favor the formation of a stronger reinforcement/matrix interface. However, considering the standard deviations of these measures it is possible to conclude, in a general way, that the film with the AM2 agent did not affect the tensile strength of the CF/PP-PE(AM2) laminate in relation to the CF/PP-PE one.

Regarding the elastic modulus results it is observed a descending order, where the best result was achieved for the CF/PP-PE(AM1) laminate (54.7 ± 2.4 GPa), followed

by CF/PP-PE(AM2) (48.6 ± 5.7 GPa) and CF/PP-PE (37.4 ± 3.4 GPa) laminates. These results reinforce that the CF/PP-PE(AM1) laminate shows the best mechanical behavior in tensile, evidencing that the AM1 agent provided greater chemical compatibility between the components, the PP-PE matrix and CF reinforcement. Probably, the existence of the largest chemical affinity conferred by maleic anhydride in the PP-PE polymer matrix assured a better interfacial adhesion fiber/polymeric matrix, which influenced positively the laminate strength. The results indicate also that the use of AM2 modifying agent, containing an elastomeric phase, increased the deformation stress of the laminate, in relation to the CF/PP-PE laminate.

The processing windows of polymers and their composites can be determined from the behavior of the complex viscosity curve (η*) versus temperature or time.

Figure 1 shows the complex viscosity graphic versus temperature of PP-PE, PP‑PE(AM1) and PP-PE(AM2) films.

As expected, Figure 1 shows that the viscosities of all films decrease with the temperature increasing. This behavior is due to the gradual destruction of existing interaction forces (van der Waals forces) with the temperature increasing.

Table 2 shows the results of strength and elasticity modulus in tensile of the bidirectional laminates (0°, 90°) tested in longitudinal direction. The analysis of this table shows that the CF/PP-PE(AM1) laminate presents higher tensile strength (507.6 ± 11.8 MPa), followed by CF/PP-PE and CF/PP-PE(AM2) laminates, with mean resistance values ​​of 440.1 ± 35.9 MPa and 422.8 ± 27.9 MPa, respectively. The correlation of these data shows that the Table 2. Results of strength and elastic modulus in tensile of the laminates.

Composite CF/PP-PE CF/PP-PE(AM1) CF/PP-PE(AM2)

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Tensile Strength (MPa) 440.1 ± 35.9 507.6 ± 11.8 422.8 ± 27.9

Elastic Modulus (GPa) 37.4 ± 3.4 54.7 ± 2.4 48.6 ± 5.7

Figure 1. Complex viscosity curves versus temperature of PP-PE films.

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Fractographic and rheological characterizations of CF/PP-PE-copolymer composites tested in tensile This structure can also be destroyed with the shear rate increasing. The PP-PE(AM1) film presents higher viscosity values ​​ranging from 1600 to 3100 Pa.s, in the temperature range used in the analysis (180 to 230°C). These higher viscosity values ​​can be attributed to the presence of maleic anhydride, which favored coupling reactions with the polyolefin, resulting in the viscosity increase and, possibly, the molar mass increase of the copolymer[55]. For this sample, it can also be observed the presence of a curve between 180 and 200 °C. This behavior evidences the action of maleic anhydride on the rheological behavior of the PP-PE in function of temperature.

The matrices of PP-PE and PP-PE(AM2) present the complex viscosity curves very close, with viscosity values ranging from 550 to 1100 Pa.s, in the temperature range between 180 and 230 °C. However, the PP-PE viscosity behavior presents a continuous drop in all temperature range, while the viscosity of the PP-PE(AM2), in the range of 180 to 200 °C, shows a sharp curve during the fall of viscosity, highlighting the AM2 agent behavior in this sample. The decrease of the viscosity of PE-PP(AM2) sample due to the elastomeric compatibilizer addition resulted in improved fluidity, in accordance with literature[54].

Figures 2-5 show representative images of the fracture surfaces resulted from the tensile tests of the studied

laminates. In this case, SEM analyses contributed to the observations and the capture of images of the fracture surfaces. The observed aspects were identified and correlated with the medium values of tensile strength, determined for each laminate studied.

The failure in tensile is one of the easiest failure modes to be found and understood in composite materials and has long been studied by many authors[42-44]. Generally, the fracture follows the development and propagation of cracks through the matrix and the fiber/polymeric matrix interface, because of the stress concentration in the material produced. Figure 2 shows detailed images of CF, obtained on the fracture surface of CF/PP-PE specimens tested in tensile. Figure 2a shows the longitudinal section of a CF, with typical superficial grooves of this type of reinforcement. This observed aspects is inherited from the polymeric precursor, polyacrylonitrile (PAN) fiber, as cited in the literature[56,57].

In this particular case, the CF image shows that this region presents a poor interfacial adhesion between fiber/matrix.

This observation can be attributed to the poor interaction of the components or due to the incomplete percolation of the polymeric matrix into the reinforcement. The considerable plastic deformation observed on the fracture surface allows characterizing the polymeric matrix as thermoplastic type, independent of previous information about the polymeric

Figure 2. Fractographic aspects of longitudinal section of CF (a) and fractured surface (CF top) representative of the CF/PP-PE(AM2) and CF/PP-PE laminates (b), respectively, tested in tensile.

Figure 3. Details of the fracture topography of CF/PP-PE (a), CF/PP-PE(AM1) (b) and CF/PP-PE(AM2) (c) laminates.

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Cirilo, P. H. S., Nogueira, C. L., Paiva, J. M. F., Guerrini, L. M., Cândido, G. M., & Rezende, M. C.

Figure 4. Partial views of CF/PP-PE (a), CF/PP-PE(AM1) (b) and CF/PP-PE(AM2) (c), with emphasis on reinforcement rich regions.

Figure 5. Details of impregnation of CF/PP-PE (a), CF/PP-PE(AM1) (b) and CF/PP-PE(AM2) (c).

matrix used in composite processing. It is also observed the presence of small voids dispersed within the matrix.

Generally, in thermoplastic composites, voids are nucleated at lower stress levels, and they are often developed in the fiber/matrix interface. With the continuing loading, the voids grow with the plastic deformation and viscoelastic behavior of the matrix, clumping together to form the fracture surface of the tested laminate[44].

Figure 2b shows another typical detail of CF laminates tested in tensile, which presents the end sections of three isolated carbon fiber filaments fractured, without the impregnation of polymer matrix. The identified fractographic aspects are highlighted by arrows on the fracture surface of each fiber.

Apparently, the failure beginning on the fiber is caused by the presence of a small defect in the fiber itself or in the fiber/matrix interface. Immediately next to the failure start point area, there is a fracture surface relatively flat, termed as mirror region, which indicates a slow fracture, in which there is only enough energy to propagate the crack. However, from the time that the fracture starts to propagate, the fracture velocity increases and the fracture topography becomes more rugged and rough with the fog aspect, forming radial line marks on the fiber surface. Radial lines present the aspect of an open range from an origin point. This aspect is often found on CF fracture surfaces. This aspect is considered relevant in the failure analysis of continuous fibers reinforced polymer composites, subjected to tensile stress and helps in the crack growth direction mapping. Similar aspects are reported in the literature[44].

Figure 3 shows representative images of the three thermoplastic laminates failed by delamination. Unlike of the fractographic aspects observed in thermoset composites, 112

where the fractographic aspects are more easily identified and explained[36,38,44], the fractographic evaluation of thermoplastic composites is much more complex. The more difficulty is due to the fact of thermoplastic matrices present few fractographic aspects due to their viscoelastic nature that generates deformations continuously during the application of mechanical loading. Thus, the investigator must know very well the characteristics of the individual components and the used processing technique in order to extract information from a complex and little elucidative scenario.

Despite this, Figure 3 shows that the three PP-PE laminates present different fracture morphologies, attributed to the modifying agents used in PP-PE films (Table 1). The CF/PP-PE laminate (Figure 3a), when compared to the other laminates, presents the worst impregnation, revealing regions of fiber oriented at 0° fully exposed, i.e., without impregnation of the polymer matrix. It is also observed extensive regions of polymer matrix positioned between CF fabric layers. These polymer rich regions show significant plastic deformation, as indicated by the fiber impressions at 90o, which are formed from the interlaminar failure propagation during the mechanical loading. The presence of these fractographic aspects on the fracture surface shows the weak interfacial adhesion fiber/matrix. It is also observed that the fracture plane of laminate presents broken fibers at 0°, occurred at various levels due to insufficient percolation of the polymer matrix into the reinforcement and poor fiber/ matrix adhesion.

Figure 3b shows a fractured region of CF/PP-PE(AM1) laminate. This figure shows a region with good consolidation, where the thermoplastic matrix is more cohesive and with aspects partially smooth and rough simultaneously.

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Fractographic and rheological characterizations of CF/PP-PE-copolymer composites tested in tensile In this case, it is observed that the reinforcement is better incorporated into the matrix, i.e., the CF is more impregnated by the polymeric matrix. These features indicate that the laminate presents an interface fiber/matrix stronger than that observed in Figure 3a. Figure 3c shows fractographic aspects identified in the region failed by delamination of the CF/PP-PE(AM2) laminate. Fiber impressions on the matrix and some loose fibers suggest the existence of a weak adhesion in the fiber/matrix interface, as noted in the CF/PP-PE laminate (Figure 3a). However, in this case, the polymer matrix presents extensive regions with smoother texture, different from that observed in the other two laminates. Probably, the manifestation of this feature was influenced by the elastomeric agent AM2 used the PP-PE(AM2) composition (Table 1).

Figure 4 shows images taken with smaller magnifications of regions with partial views of the CF reinforcement.

This figure shows clearly the presence of loose and broken CF, Figure 4a and Figure 4c, relative to CF/PP-PE and CF/PP-PE(AM2) laminates, respectively. In Figure 4a are observed matrix rich regions, presenting loose matrix layers with fibers impression, adjacent to the fracture plane. These observations confirm the weak interfacial adhesion between the components of CF/PP-PE laminate. In Figure 4c, the fibers are unprotected of matrix, misaligned and with many fragments of matrix on the surface. These evidences show weak interfacial adhesion and also the probably occurrence of problems in the composite processing stages.

Figure 4b shows the CF bonded to the polymeric matrix, indicating the wetting of the reinforcement by the matrix and therefore a better consolidation, with more consistent interface between the components of the composite. The presence of broken fibers in different sizes impregnated by the matrix consists of a fractographic aspect that indicates that the laminate failed at different stress levels with the applied load increasing. This behavior is further evidence that the use of AM1 coupling agent improved the reinforcement/matrix interaction, in relation to the other two laminates. This suggests a better interfacial adhesion, which provided a greater tensile strength as compared with PP-PE and PP-PE(AM2) laminates (Table 2).

In the hot compression molding of composites is essential the complete percolation of the polymer matrix among the reinforcing layers. This condition provides bonding of the fibers and a good consolidation of the laminate structure with orderly and consistent interfacial adherence of fiber/matrix.

Figure 5 shows a comparison of the impregnating features observed in the laminates studied. It is evident that the laminate CF/PP-PE(AM1) presents a much more efficient impregnation of the reinforcement (Figure 5b), than the other two laminates (Figure 5a and Figure 5c). In Figure 5a,c are observed misaligned fibers without impregnation, emphasizing the weak interfacial adhesion. These observations can be attributed to variations in the composite molding process, including difficulty in the matrix flow, which may prejudice the wetting and interactions between polymer and reinforcing under the actions of temperature and pressure. The matrix texture appearance, shown in Figure 5b, suggests that the fibers wetting quality promoted by the PP-PE(AM1) matrix increased the physicochemical contact between the composite components.

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The correlation of SEM observations with the tensile strength results presented in Table 2 is consistent.

This correlation shows that the CF/PP-PE(AM1) laminate has better impregnation of the reinforcement and the highest value ​​of both tensile strength (507.6 ± 11.8 MPa) and elastic modulus (54.7 ± 2.4 GPa), when compared to the other two laminates. The main influence of weak interfacial adhesion is the reduction of the mechanical properties, which compromises the final application of the composite.

The comparison of SEM observations with the rheological curves of the polymer matrices (Figure 1) shows that the greatest viscosity values of the PP-PE(AM1) matrix did not affect the wetting of the reinforcement. In this case, the processed laminate presents the best impregnation and the best mechanical performance. This result shows that the AM1 agent (maleic anhydride) acts as a good coupling agent, increasing the polarity of the PP-PE matrix and improving the reinforcement/matrix interaction. The efficiency of using the maleic anhydride as compatibilizing of copolymers and blends, aiming to improve adhesion and hydrophobicity of polyolefins, has been widely reported in literature[55,58].

Therefore, the interface feature has significant influence on both mechanical behavior and failure process of tensile tested-specimens.

4. Conclusions In the present work, the morphological aspects and tensile mechanical properties of CF/PP-PE laminates, with three different compositions of the PP-PE, were investigated. Two PP-PE matrices were modified with the agents designated AM1 (maleic anhydride) and AM2 (elastomeric agent). The CF/PP-PE(AM1) laminate showed the best tensile strength result (507.6 ± 11.8 MPa). While the CF/PP-PE(AM2) presented a reduction in strength at around 4% (422.8 ± 27.9 MPa) compared to CF/PP-PE (440.1 ± 35.9 MPa). Similarly, the highest elastic modulus values were determined for the CF/PP-PE(AM1) (54.7 ± 2.4 GPa), followed by the CF/PP-PE(AM2) (48.6 ± 5.7 GPa) and CF/PP-PE (37.4 ± 3.4 GPa) laminates.

These results show that the incorporation of AM1 agent in the PP-PE film contributed to increase the physicochemical interaction fiber/matrix and provided better conditions for the charge transference between matrix and reinforcement.

Fractographic analyses of the CF/PP-PE laminate shows poor adhesion between fiber and polymer matrix. This aspect was indicated by the presence of fiber impressions in polymer rich regions and CF surfaces totally non-impregnated by polymer matrix. The most consistent adhesion was observed for the CF/PP-PE(AM1) laminate, in accordance with the best mechanical performance in tensile. Plastic deformations are identified in the fracture morphology of the polymer rich regions attributed to the viscoelastic behavior of thermoplastic matrix. Other fractographic aspects, such as radial line marks, are observed in the fracture surface of the fibers oriented at 0°. Areas with significant volume of fiber non-impregnated suggest possible problems in the composite molding process. Complex viscosity results showed that this parameter did not influence the wetting of the carbon fiber reinforcement.

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5. Acknowledgements The authors thank CNPq for the financial supports (Process numbers: 150697/2014-7, 142314/2010-2 and 303287/2013-6) and CAPES/PVNS.

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Received: Dec. 23, 2015 Accepted: June 29, 2016

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http://dx.doi.org/10.1590/0104-1428.2206

O Evaluation of silanes in SBR 1502 / Telinne Monspessulana flour composites O Oscar Buitrago *, Oscar Palacio and Emilio Delgado O Facultad de Ingeniería, Universidad Militar Nueva Granada, Bogotá D.C., Colombia O O Abstract O A comparative study was performed on the effect of the addition of silane coupling agents (SCA), vinyltrimethoxysilane (VTMS) and 3-aminopropyltriethoxysilane (APTES) to a mixture of styrene butadiene rubber SBR1502 with Telinne O Monspessulana flour (TMF). SCA was directly added into the mixture using untreated and mercerized TMF. Also, TMF and SBR1502 mixing trials were conducted with the TMF previously mercerized and injected with each of the silanes. rubber compounds were subjected to tensile tests in order to evaluate the coupling power of both SCA. It was found O The that the vinyl silane type produced the best results in the tensile strength.

O Keywords: composite, natural fiber, SBR1502, silanes.

O 1. Introduction O Using natural fibers or wood flour as a potential filler to, conditions, and imparted stiffness and toughness to the at least partially, replace mineral fillers for the production of compound, although the excess in flour concentration leads O polymeric compounds is important because of the benefits to decrease of these mechanical properties . by plant fibers. These include the fact that their low Generally, natural fibers are subjected to mercerization O offered density allows the development of lighter compounds, that as a prelude to treatment with type SCA cupping agents; are non toxic, they are biodegradable, their production NaOH treatment allows the OH groups of cellulose to be O they can be sustainable, and they are less abrasive, increasing exposed directly and thus better fiber-SCA coupling can be the service life of machinery. achieved. The grafted fiber is then mixed with the polymer .

O SBR styrene butadiene rubber is an elastomer, widely used The coupling agent can also be added directly during the in the manufacture of items such as tires, hoses, footwear, mixing of the fiber with the polymer .

It is important to note that any SBR compound must O packing, conveyor belts, mats, and others .

1

1

1

1

*oscar.buitrago@unimilitar.edu.co

[5]

[6,7]

[8]

[1]

There are also numerous studies that examine the effect of incorporating natural fiber with SBR rubber.

For example, Kumar et al.[2] analyzed the rheological behavior of SBR1502 compounds with sisal fiber, finding that these acquire pseudo-plastic behavior upon chemical treatments owing to the strong interfacial adhesion between the fibre and the rubber matrix. Wang et al blended 50 parts of SBR1502 with 50 parts of linear low density polyethylene (LLPE) with rice husks and used maleic anhydride (MA) as a compatibilizer. The mechanical properties were optimal when the concentration of MA was 2.5 parts per hundred of rubber (phr). They also mention that when more than 25 phr of rice husk was used, the dielectric properties of the composite decreased[3]. Kumar and Thomas[4] analyzed SBR 1502 compounds with sisal fibers, finding that the tensile strength and tear resistance improved when the fiber orientation was longitudinal and its size was 6 mm in length. They also mention that the optimal concentration of sisal fiber was 35 phr.

Regarding the use of silanes to improve compatibility of natural fiber with SBR, Wang et al mixed Si69 silane with SBR 1502 and 20 phr of silica, varying the concentration of hemp flour. Adding the flour improved the vulcanization

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be cross-linked with sulfur (vulcanization), using peroxide or radiation in order to increase mechanical, chemical and thermal resistance to the respective compound[9-12].

In our research, we analyzed the effect of the incorporation of VTMS and APTES on composites prepared with untreated SBR1502 + TMF, SBR1502 + mercerized TMF and SBR1502 + grafted TMF with SCA. By measuring the tensile properties and hardness, we comparatively evaluated the coupling effect of both silanes.

2. Materials and Methods 2.1 Materials The following materials were used: Telinne Monspessulana flour with a particle diameter of 400-800 µm and a humidity of 0.20%. Rubber: SBR1502 manufactured by INSA S.A. Mexico. Coupling Agents: APTES YAC A110 of 98.41% purity produced by Nanjing Lanya Chem. Co. Ltd., VTMS Struktol SCA 971 of > 98.6% purity manufactured by Struktol. Processing aid: Struktol WB16 of Struktol. Antioxidant: Irganox 1076 manufactured by BASF. Cross-linking agent DCP Perkadox BC FF, purity 99%, produced by Akzo Nobel.

Polímeros, 27(2), 116-121, 2017

Evaluation of silanes in SBR 1502 / Telinne Monspessulana flour composites 2.2 Methods 2.2.1 Experimental design Seven (7) tests were conducted with the base formula as seen in Table 1. All quantities remained constant.

The “Control” test was defined for the compound of SBR1502 with TMF, untreated and free of SCA. Two (2) modes of silane addition were tested: 1) direct addition of SCA during mixing of SBR1502 with TMF with no surface treatment.

Direct addition of SCA during the mixing of SBR1502 with mercerized TMF. 2) Pretreatment of mercerized TMF with SCA and later mixing with SBR1502 (see Table 2).

The tensile properties of the silane were measured to assess its ability to couple the flour and the rubber.

2.2.2 Surface treatment TMF Mercerization: Telinne Monspessulana flour was mercerized in NaOH solution at 8% for 4 hours at 35 °C, according to the methodology used by Buitrago et al.[13].

Graft silane-TMF: The flour previously mercerized was grafted with silane. The amount of SCA was 10% based on the weight of TMF. Silane was pre-hydrolyzed for one hour, the time of TMF immersion was 4 hours and drying was performed at 45 °C for 24 hours. The above procedure was performed for the APTES and VTMS following the methodology used by Buitrago et al.[13].

2.2.3 Preparation of compounds Compounding: The preparation of the mixture was performed in a roller mill with a capacity of 600 cm3.

The SBR1502 was placed on rollers until band formation.

Then, immediately, 1/3 of the total TMF was added, next the SCA was added (see Table 2). TMF continued to be added slowly. When the incorporation of the flour was completed, the processing aid and antioxidant were added.

The DCP was added three (3) minutes before the end of the stage. The compounding temperature was 115 ± 5 °C; the processing time was 13 minutes.

Compression molding: We used a hydraulic press heated by electrical resistance. Mold dimensions: 17 cm*17 cm and 3.6 mm of thickness.

The operating parameters were: mold temperature 160 ± 1 °C, 5.24 MPa of specific pressure, pressing time of 6 minutes.

2.2.4 Methods of analysis Gel Percentage: It quantifies the degree of crosslinking of the compound. It was determined following the indications of ASTM D 2765-01, method A[14]. It is interesting to observe that soxhlet extraction diluted the non-cross-linked polymer, although it swelled the rubber compound. To determine the final weight, the samples that were subjected to extraction were dried at 50 °C for 4 hours. Subsequently they were cooled in liquid nitrogen for 24 hours, then they were pulverized and heated again at 80 °C for 4 hours to remove the solvent.

The sample calculation is shown in the Equation 1. % GEL = (WF − f *WI ) / (1 − f ) *WI  *100 (1)

WF and WI represent, respectively, the final and initial weight of the rubber compound, and f is the fraction of fiber in the composite.

Tensile tests: The specimen type B was used according to ASTM D412-06a standard[15]. The tests were performed on a universal machine Shimadzu AGS-X. Test Parameters: travel speed of 50 mm/min, temperature 24 °C, and relative humidity 45%. Five specimens were used for each formulation.

Hardness: Hardness was determined according to ASTM D2240-05 standard[16] at 24 °C and 45% of relative humidity. INSIZE durometer of A type was used. Five (5) measurements were performed for each test.

SEM: The fractured surface of the specimens under tensile strength were observed by scanning electron microscopy (SEM), the procedure was performed following the methodology used by Tobón et al.[17].

3. Results and Discussions 3.1 Gel content In Figure 1 an increase is observed in the percentage of gel when incorporating the SCA compared to the control, regardless of the means of addition.

Table 1. General formulation (all amounts in phr).

SBR 1502 100

TMF* 25

Antioxidant 1

Processing aid 1

DCP 1

SCA** 1

*Untreated / mercerized / grafted (see Table 2); **VTMS or APTES in direct addition to the mixture (see Table 2).

Table 2. Description of tests.

SCA Control VTMS APTES VTMS APTES VTMS APTES

Polímeros, 27(2), 116-121, 2017

Surface Treatment TMF

Method of incorporating the SCA Without silane

Without treatment

Direct addition of the SCA into the compound mixture

Mercerized

Direct addition of the SCA into the compound mixture

Mercerized

Surface treatment of graft SCA with fiber before mixing

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Buitrago, O., Palacio, O., & Delgado, E. This could mean that there was no competition between the silane coupling reactions with the flour and the crosslinking reaction of the polymer chains with DCP[17-19].

When comparing the two silanes difference in values ​​was observed between them, it shows that the percentage of gel APTES was higher independent of the method of addition.

3.2 Young’s modulus The modulus of elasticity was significantly higher in composites prepared with fiber previously grafted with APTES and VTMS silanes, with an increase in the modulus of elasticity of 28% and 37%, respectively, compared to the control (see Figure 2). These results in the modulus match the trends of other studies of hemp and wood with silane coupling agents[7,20-22] where the modulus of elasticity improved because the presence of the coupling agent allowed a more uniform dispersion in the rubber matrix allowing a strong interfacial bonding with fiber. The direct addition of the silanes did not significantly affect the modulus of elasticity regardless of the fiber treatment. Comparatively, it was not observed that the silane type influenced the Young’s modulus.

3.3 Tensile strength

Figure 1. % of Gel in SBR1502 compounds.

In Figure 3 the tensile strength results are presented.

The incorporation of APTES as a coupling agent did not affect the values of tensile strength compared to the control, while the incorporation of VTMS increased tensile strength regardless of the mode of addition; the maximum value was obtained with direct addition using untreated flour and was 30% higher compared to the control, due to the improvement of the transfer of stress in particle-polymer interface[20,22].

The other compounds, which have VTMS, increased traction by 15% compared to the control.

The above results support the conclusion that the vinyl type silane VTMS has superior performance compared to APTES in composites prepared with SBR1502. It is shown that mercerized flour affected the tensile strength negatively in both silanes when these were added directly to the mixture. This phenomenon may be attributed in particular to the conditions used in the mercerization process, where the concentration of NaOH at 8% could possibly generate the presence of large amounts of Na+ ions, which could in turn cause interference in the coupling[23].

3.4 SEM fractography analysis Figure 2. Young’s modulus for SBR1502 compounds.

Figure 3. Tensile strength for SBR1502 compounds.

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All SEM images show dark layers (SBR matrix), the wood flour particle is white; also voids formed by the release of the particle when breaking the specimen in the tensile test are observed.

Variation is observed in the shape and size of wood flour, for example many are below the nominal specification written at the beginning of the experimental part (400-800 µm).

This is because subsequent reduction or breaking of the particles during the shear mixing step.

There are morphological differences in fracture zones of some specimens, specifically on the surface of the wood particle. The SEM image of the composite VTMS With TMF-untreated (Figure 4a), has filament formation which corresponds to the formation of grafts between natural fiber and polymer, which is consistent because this specimen showed the highest value tensile strength (Figure 3).

Similarly in Figure 4b forming grafts shown in APTES with TMF-untreated compound, precisely this specimen showed the highest tensile strenght value when the amino silane is used.

Polímeros, 27(2), 116-121, 2017

Evaluation of silanes in SBR 1502 / Telinne Monspessulana flour composites No evidence of grafting is evident in the specimens where the silanes were added to the previously mercerized wood flour (Figure 4b,c), this is consistent as the results of tensile strength were low.

The SEM of specimens of wood flour previously grafted with silanes (Figure 5a, b) have different morphology compared to the control (Figure 6). It was observed that the particles have surface modification, but this change

was not sufficient to achieve grafting between the particle with SBR matrix.

3.5 Elongation at break Figure 7 shows no significant difference between both silanes. The compounds prepared with APTS yielded lower values compared to the control, confirming that the

Figure 4. SEM 1000x. Fractography of specimens. Direct adition of silanes. (a) VTMS-TMF untreated; (b) APTES-TMF untreated; (c) VTMS-TMF mercerized; (d) APTES-TMF mercerized.

Figure 5. SEM 1000x. Fractography of specimens TMF grafted with silane, (a) VTMS; (b) APTES. Polímeros, 27(2), 116-121, 2017

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Buitrago, O., Palacio, O., & Delgado, E.

Figure 6. SEM 1000x. Fractography specimen control.

Figure 8. Energy at break for SBR1502 compounds.

Figure 7. Elongation at break of SBR1502 compounds.

Figure 9. Hardness for SBR1502 compounds.

addition of the amino type silane does not contribute to the coupling between TMF with SBR1502. It is also noted that the compound prepared with VTMS grafted fiber was lower compared to the forms of addition.

4. Conclusions

3.6 Energy at break The results of the energy required to break rubber compounds with TMF indicate that it is not possible to determine any significant differences in all of the tests (Figure 8).

3.7 Hardness The results of hardness tests are seen in Figure 9; the compounds made with VTMS have a slight increase in hardness compared to the control. The maximum increase in hardness was achieved by the compound prepared with fiber grafted with APTES and VTMS. The addition of silanes can in some way increase the hardness, which is consistent with other studies into chitosan and sisal[7,20,21].

There is a relationship of hardness to elastic modulus, the highest hardness values also ​​ have greater Young’s modulus (Figure 2).

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Adding silanes did not decrease the percentage of gel; it was not possible to identify interference between the coupling reaction and cross-linking; comparatively gel percentage was higher in the compounds having APTES that VTMS. VTMS and APTES improve tensile strength in SBR-wood flour composites, the vinyl silane produced the highest values.

It is possible that the presence of the double bond of the vinyl group contributed to the formation of bonds with the fiber surface and rubber. This silane is recommended by the manufacturer for coupling silica type loads with SBR1502.

When the silanes were added directly on SBR-wood flour untreated, yielded values ​​greater of tensile strength compared to other modes of incorporation, which is confirmed by the observation of grafting on surface particle through of SEM images.

It was expected that mercerization of the fiber would help to improve the coupling of the vegetable flour with rubber, especially with VTMS silane. This did not happen, possibly due to the conditions used in the mercerization.

Polímeros, 27(2), 116-121, 2017

Evaluation of silanes in SBR 1502 / Telinne Monspessulana flour composites APTES amino type silane increased rigidity and hardness, but no positive effect was observed on tensile strength.

Further studies are recommended to treat TMF mercerizing, specifically by varying the concentration of NaOH. In the SEM images can be seen a change in morphology of wood flour mercerized vs untreated, but these effects did not ensure coupling between with SBR matrix.

5. Acknowledgements This study is derived from the project INV-ING-1547 sponsored by “Vicerrectoria de Investigaciones de la Universidad Militar Nueva Granada” - One year term 2014.

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11. Changjie, Y., Zhang, Q., Junwei, G., Junping, Z., Youqiang, S., & Yuhang, W. (2011). Cure characteristics and mechanical

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properties of styrene-butadiene rubber/hydrogenated acrylonitrile-butadiene rubber/silica composites. Journal of Polymer Research, 18(6), 2487-2494. http://dx.doi.org/10.1007/ s10965-011-9670-y.

12. Shih, R.-S., Kuo, S.-W., & Chang, F.-C. (2011). Thermal and mechanical properties of microcellular thermoplastic SBS/PS/ SBR blend: effect of crosslinking. Polymer, 52(3), 752-759. http://dx.doi.org/10.1016/j.polymer.2010.12.026.

13. Buitrago, O., Delgado, A., & Aperador, W. (2014). Surface treatment of straight Retamo Liso (Telinne monspessulana) by silane coupling agents (SCA). Ciência e Técnica Vitivinícola, 29(12), 11-23.

14. American Society for Testing and Materials – ASTM. (2001).

ASTM D2765-01: standard test methods for determination of gel content and swell ratio of crosslinked ethylene plastics.

West Conshohocken: ASTM. 15. American Society for Testing and Materials – ASTM. (2013).

ASTM D412-06a: standard test methods for vulcanized rubber and thermoplastic elastomers. West Conshohocken: ASTM. 16. American Society for Testing and Materials – ASTM. (2010).

ASTM D2240-05: standard test methods for rubber property: durometer hardness. West Conshohocken: ASTM. 17. Tobón, A. E. D., Chaparro, W. A. A., & Rivera, W. G. (2014).

Mejoramiento de las propiedades de tensión en WPC de LDPE: HIPS/fibra natural mediante entrecruzamiento con DCP. Polímeros: Ciência e Tecnologia, 24(3), 291-299. http:// dx.doi.org/10.4322/polimeros.2014.026.

18. Ahmad, E. E. M., & Luyt, A. S. (2012). Effects of organic peroxide and polymer chain structure on morphology and thermal properties of sisal fibre reinforced polyethylene composites.

Composites. Part A, Applied Science and Manufacturing, 43(4), 703-710. http://dx.doi.org/10.1016/j.compositesa.2011.12.011.

19. Mokoena, M. A., Djoković, V., & Luyt, A. S. (2004).

Composites of linear low density polyethylene and short sisal fibres: The effects of peroxide treatment. Journal of Materials Science, 39(10), 3403-3412. http://dx.doi. org/10.1023/B:JMSC.0000026943.47803.0b.

20. Ryu, S. R., & Lee, D. J. (2007). Effects of short-fiber shape on tensile properties of reinforced rubber. Journal of Materials Science, 42(3), 1019-1025. http://dx.doi.org/10.1007/s10853006-1397-5.

21. Ismail, H., Shaari, S. M., & Othman, N. (2011). The effect of chitosan loading on the curing characteristics, mechanical and morphological properties of chitosan-filled natural rubber (NR), epoxidised natural rubber (ENR) and styrene-butadiene rubber (SBR) compounds. Polymer Testing, 30(7), 784-790. http://dx.doi.org/10.1016/j.polymertesting.2011.07.003.

22. Wang, J., Wu, W., Wang, W., & Zhang, J. (2011). Effect of a coupling agent on the properties of hemp-hurd-powder-filled styrene–butadiene rubber. Journal of Applied Polymer Science, 121(2), 681-689. http://dx.doi.org/10.1002/app.33744.

23. Gwon, J. G., Lee, S. Y., Doh, G. H., & Kim, J. H. (2010).

Characterization of chemically modified wood fibers using FTIR spectroscopy for biocomposites. Journal of Applied Polymer Science, 116(6), 3212-3219. http://dx.doi.org/10.1002/ app.31746.

Received: June 19, 2015 Revised: Mar. 21, 2016 Accepted: June 29, 2016

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http://dx.doi.org/10.1590/0104-1428.2406

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

Reprocessability of PHB in extrusion: ATR-FTIR, tensile tests and thermal studies Leonardo Fábio Rivas1, Suzan Aline Casarin2, Neymara Cavalcante Nepomuceno3, Marie Isabele Alencar1, José Augusto Marcondes Agnelli2, Eliton Souto de Medeiros3, Alcides de Oliveira Wanderley Neto4, Maurício Pinheiro de Oliveira5, Antônio Marcos de Medeiros1 and Amélia Severino Ferreira e Santos3* Department of Materials Engineering – DEMat, Universidade Federal do Rio Grande do Norte – UFRN, Natal, RN, Brazil 2 Department of Materials Engineering – DEMa, Universidade Federal de São Carlos – UFSCar, São Carlos, SP, Brazil 3 Department of Materials Engineering – DEMAT, Universidade Federal da Paraíba – UFPB, João Pessoa, PB, Brazil 4 Department of Chemistry – DQ, Universidade Federal do Rio Grande do Norte – UFRN, Natal, RN, Brazil 5 Institute of Science and Technology – ICT, Universidade Federal de São Paulo – UNIFESP, São José dos Campos, SP, Brazil 1

*amelia@ct.ufpb.br

Abstract Mechanical recycling of biodegradable plastics has to be encouraged, since the consumption of energy and raw materials can be reduced towards a sustainable development in plastics materials. In this study, the evolution of thermal and mechanical properties, as well as structural changes of poly(hydroxybutyrate) (PHB) up to three extrusion cycles were investigated. Results indicated a significant reduction in mechanical properties already at the second extrusion cycle, with a reduction above 50% in the third cycle. An increase in the crystallinity index was observed due to chemicrystallization process during degradation by chain scission. On the other hand, significant changes in the chemical structure or in thermal stability of PHB cannot be detected by Fourier transform infrared spectroscopy (FTIR) and thermogravimetric analyses (TGA), respectively.

Keywords: biopolymer, degradation, PHB, recycling, reprocessing.

1. Introduction The rate of municipal solid waste (MSW) generation is rising more than the rate of urbanization around the world[1].

By 2025, the volume of MSW generated worldwide is expected to double, reaching an amount of 2.2 billion tonnes per year[1]. In general, plastics are of main concern because they represent between 3 and 14.3 wt% of the total MSW, most are from non-renewable sources and have relatively low recycling rates[1-3].

A key practice to minimizing the environmental problems associated with solid waste is the practice of the 3R (Reduce, Reuse and Recycle) concept, which contributes to reduce energy and natural resources consumption, extend life cycle of products and landfills, and store carbon for a longer period, all important initiatives within the concept of sustainable development. Another alternative that has gained importance is the replacement of petroleum-based plastics by biopolymers from renewable resources, which are generally compostable and able to close the carbon cycle.

According to the European Union Regulation, composting is a considered and accepted way of recovering biodegradable packaging wastes[3]. However, recycling of biopolymers can be an alternative to composting and to reduce carbon dioxide emission, thus meeting the 3R concept. This approach not only saves energy that would be used in polymer synthesis, but also reduces costs associated with bioresources (monomers) used to produce new bioplastics and spares carbon resources.

Since biopolymers are commonly used in the production

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of short-life goods, recycling of these polymers becomes even more important because of the large amount of waste generated in markets, as well as insumes consumed.

Nevertheless, in the case of biodegradable plastics, there are few studies related to their recycling[4-7]. Life cycle analysis[8] of poly(lactic acid) (PLA), which is one the most studied biopolymers, pointed out that despite of its low environmental impacts, this situation can worsen if recycling is not achieved. According to Shah et al.[9], pyrolysis, solvolysis and enzymatic monomer recycling are the most promising feedstock recycling technologies for biodegradable plastics because most biopolymers have heteroatoms in their main chain, which reduce thermal stability and increase susceptibility to hydrolysis. Besides that, studies of biopolymers recyclability by multiple reprocessing are also important due to the fact that the production of biopolymer-based products may involve more than one extrusion step and mechanical recycling is the most energy efficient recycling technology.

Polyhydroxyalcanoates (PHAs) have attracted a great deal of attention due to their types, properties similar to conventional plastics, biodegradability, biocompatibility, renewability, and synthesis by biotechnological processes[10,11].

Although their discovery occurred early, PHAs were rediscovered in the 1980s, and their use has gained increasing attraction, particularly in the last decades, as a

Polímeros, 27(2), 122-128, 2017

Reprocessability of PHB in extrusion: ATR-FTIR, tensile tests and thermal studies consequence of the reduction and limitation of the fossil fuel resources, allied with the increasing concern about the environmental issues caused by improper manipulation, use and disposal of such non-renewable materials. Among these polymers, biosynthesized polyhydroxybutyrate (PHB), the most common PHA, has a high potential to compete with commodity polymers due to its similarity in crystallinity, melting temperature, tensile strength and Young’s modulus, and competitive costs[12]. Nevertheless, high stiffness, brittleness and poor thermal stability above the melting temperature (a narrow processability window) are some of its drawbacks compared to many conventional polymers[13-15].

Feedstock recycling of PHB has already been considered as a viable route to produce end products such as crotonic acid, linear oligomers having crotonate end groups and a cyclic trimer, as well as to produce plasticized PLA by reactive extrusion grafting of PHB degradation products onto PLA chains[16-19]. At temperatures a little above the melting point, there is a rapid decrease in molecular weight and a subsequent production of crotonic acid as the main volatile product[20-23]. The presence of moisture, fermentation residues, oxygen, metal and alkali catalysts also is known to favor thermal degradation[14,24-28]. This is an important limitation for processing and consequently, to mechanical recycling. Therefore, there has been a great deal of interest in studying the thermal degradation behavior of PHB and other related poly(hydroxyalkanoate)s[29-35]. Also studies seeking to improve the thermal stability of PHB by grafting chemicals in PHB chain[33,36], and adding polymeric additives in PHB matrix[37,38] have been developed.

The thermal degradation behavior of PHB has been discussed in many works[14,20,21,30,39-42], in which a random chain scission reaction (β-elimination) involving a six-membered ring transition state (Figure 1) has been considered as the main mechanism based on typical structures of pyrolysis products, i.e., crotonic acid and oligomers with a crotonate end group, i. e., unsaturated end groups. Since the proposed mechanism is a non-radical random chain scission the conventional stabilizers and antioxidant was not efficient to prevent PHB degradation[38].

In this work, the effect of multiple reprocessing cycles on the properties of PHB products obtained by more than one reprocessing cycle, as well as in their recyclability was evaluated. Thermal and mechanical properties were measured after each extrusion cycle, and correlated with polymer structural changes.

2. Materials and Methods 2.1 Materials Poly(3-hydroxybutyrate) (PHB) powder, produced by bacterial fermentation, was kindly supplied by PHB Industrial S/A, Serrana-SP, Brazil, registered under the brand BIOCYCLE and used as received. This polymer has a melt flow index of 6.5 g/10 min (190 °C, 2.16 kg) and

number-average molecular weight (Mn) of 56,650 g.mol-1, weight-average molecular weight (Mw) of 167,223 g.mol-1, and polydispersity index (PDI) of 2.95, as measured by gel permeation chromatography.

2.2 Reprocessing cycles Extrusions were carried out using an AX Plásticos (São Paulo, Brazil) single screw extruder (ϕ = 16 mm, L/D = 26 and 90 rpm) with Maddock mixing section between the compression zone and the flow control zone up to three times. Temperature profile used was: 165, 170 and 170 °C. The neat PHB powder (virgin polymer) or pellets (extruded polymer) was dried at 60 °C under vacuum for 12 h before each processing cycle. Since it is a first approach on recyclability of PHB, no other materials/additives, such as nucleating agent or plasticizers, were used.

2.3 Compression molding Compression molded films were prepared in a hydraulic press (capacity of 24kgf) at 170 °C for 1 min at about 10 kgf, followed by quenching in an ice-water bath.

2.4 Mechanical properties Tensile tests were carried out according to ASTM D 882 at 25 °C and strain rate of 1%.min-1 using a TA Instruments DMA, Q800. The test-specimens with dimensions around (50 × 10 × 0.075) mm were cut from films prepared by compression molding. At least 5 samples were carried out for virgin PHB and after each extrusion cycle. Results were averaged arithmetically.

2.5 Fourier transform infrared (FTIR) spectroscopy Infrared spectra were obtained by a Nexus 470 Nicolet FTIR spectrophotometer. 32 scans, resolution of 4 cm-1 and interval of 2 cm-1 were used. Analyses were performed in the attenuated total reflectance mode (ATR) by direct analysis of films on SnZn crystal.

2.6 Thermogravimetric analysis (TGA) Thermogravimetric analyses (TGA) were conducted under 50 mL.min-1 of nitrogen flow, using a Shimadzu TGA 60 apparatus. The samples were heated to 600 °C at 20 °C.min-1. The characteristic degradation temperatures: temperature at the maximum of the DTG curve (Tmax) and the temperature (Tx%) at which the sample loses x% of its initial weight were determined.

2.7 Differential scanning calorimetry (DSC) Thermograms were obtained using a Shimadzu DSC 60 differential scanning calorimeter (DSC). Calibration was performed with indium and tin in the temperature range from 0 to 350 oC. The sample weight was approximately

Figure 1. PHB random chain scission mechanism.

Polímeros, 27(2), 122-128, 2017

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Rivas, L. F., Casarin, S. A., Nepomuceno, N. C., Alencar, M. I., Agnelli, J. A. M., Medeiros, E. S., Wanderley, A. O., No., Oliveira, M. P., Medeiros, A. M., & Santos, A. S. F. 6-10 mg. All the samples were heated from 25 °C to 220 °C at 10 °C.min-1 in purge flow of nitrogen at 50 mL/min. Melting enthalpies were determined using constant integration limits.

The degree of crystallinity (Xc) was determined using the following equation: Xc(%) =

∆Hm x100 (1) ∆H100%

where: ΔHm is the melting enthalpy per unit of weight of PHB samples and ΔH100% denotes enthalpy per unit weight of the 100% crystalline PHB, which is assumed to be 146 J/g[43]. All DSC analysis was taken from the films obtained by compression molding in order to best represent the material undergoing mechanical test. This procedure has been already adopted by Srubar et al.[44].

3. Results and Discussions 3.1 Mechanical properties In order to assess the industrial possibility of PHB recycling, its mechanical properties must stay as stable as possible along the processing cycles. A representative tensile curve of virgin PHB was depicted in Figure 2, which shows its brittle nature. Therefore, in this case, the tensile strength is coincident with tensile stress at break. Figure 3 shows the values of tensile strength of PHB, before and after each extrusion cycle. It can be observed that the tensile stress at break decreased from 32.1 MPa to 13.4 MPa after three extrusion cycles. At the 3rd cycle, the tensile stress at break had the same order of magnitude of the tensile strength of LDPE[45]. However, the elongation at break of these two polymers is very distinct and the tactile sensation of PHB at the 3rd cycle is like a brittle material, which hindered additional extrusion cycles for further characterization of samples.

According to Pillin et al.[5], a similar decrease in tensile strength values was observed only at the 6th injection cycle for PLA. Nevertheless, the properties of PLA stay useful,

Figure 2. Representative tensile stress-strain curve of virgin PHB determined at 1 mm.min-1 by using a universal machine Shimadzu AG-X 10kN model, based on the ISO standard 527- 2.

The test‑specimens, type 5B, were prepared by laser cutting from films prepared by compression molding.

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i.e., above about 10 MPa, until the 7th cycle, the maximum number of cycles evaluated.

The color change of PHB from an off-white powder to opaque brown color was observed after first extrusion, as reported in the literature[46,47]. Such behavior is accounted for PHB chromophoric carbonyl groups.

The strong decrease in tensile strength is probably ascribed to a reduction in the molecular weight of PHB due to chain scission reactions caused by thermal degradation[15,20,31] during extrusion cycles. It is well established that the former properties are strongly affected by chain scission, since the broken chains are confined in the amorphous regions between the lamellae, where tie molecules, which are responsible for the mechanical integrity of semicrystalline polymers, are located[48].

The hypothesis of mechanical properties decrease with molecular weight was in good agreement with the results achieved by Sadi et al.[47] and Renstad et al.[49]. In both works, the tensile strength at break for PHB and (poly-3-hydroxy butyrate-co-valerate) (PHBV), respectively, decreased with molecular weight depression. Nevertheless, none of these studies have determined until how many extrusion cycles the mechanical properties of PHB still useful and the rate that decay in the mechanical properties with extrusion cycles occurs. Kendall[50], for example, in its life cycle assessment for the production of PHB from material recovery facilities considers that there is no current technology for PHB mechanical recycling, but in this study results show that it is possible to recycle PHB up to three extrusion cycles without any additive to inhibit its degradation or improve its properties. For conventional plastics, it is known that during thermomechanical recycling[51-53], mechanical properties decrease with increasing multiple extrusion or injection molding cycles. This trend depends upon the type and chemical nature of the polymer. Nevertheless, blends of recycled polymers with virgin ones, reprocessing with stabilizers and incorporation of reinforcing fillers are alternatives to improve recycled plastics properties[54-57].

Furthermore, changes in mechanical properties could be also attributed to changes in the structure and stability of crystalline state or rearrangement of interlamellar amorphous state[44,58-60]. Nevertheless, changes in molecular weight or crystalline and amorphous morphologies are beyond the scope of this study.

Figure 3. Results of tensile strength for PHB before and after each extrusion cycle.

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Reprocessability of PHB in extrusion: ATR-FTIR, tensile tests and thermal studies 3.2 Fourier transform infrared (FTIR) spectroscopy The original chemical structure of PHB consists of molecules terminated by a hydroxyl and a carboxyl group.

The hydroxyl and carboxyl end groups are observed at approximately 3600 cm-1 and 1720 cm-1, respectively.

Other characteristics PHB vibrations appear at around 1277cm−1 and 970 cm-1. The peak at 1277 cm-1 denotes the –C-O-C- group and at 970 cm-1 is assigned to bending vibrations of olefinc -C-H[61,62].

As the thermal degradation proceeds, it incorporates vinyl (crotonate) ester and carboxyl groups end groups in PHB structure[20,30,31,41]. Therefore, a gradual increase in crotonate ester groups with extrusions paths can be expected, as well as a decrease in hydroxyl groups present in the original polymer. The absorption band assigned to stretching vibrations of double carbon/carbon bond, -C=C-, is expected to appear at around 1660 cm-1[63].

FTIR analyses of PHB samples during extrusion cycles are shown in Figure 4, which presents the spectra of virgin PHB and PHB after each extrusion cycle.

The presence of absorption bands associated to the formation of new chemical groups due to degradation mechanisms of polymer was not noticed. According to Yu et al.[64], the band at 1700-1720 cm-1, assigned to the carbonyl absorption band in the infrared spectra, is shifted to 1654 cm-1 when conjugated with vinyl end groups. Nevertheless, no shift can be observed for this band (Figure 5), corroborating the results of Yu et al.[44] that observed no difference in infrared absorption of PHB film surface even after up to 40 wt% of the PHB film submitted to alkaline hydrolysis had been decomposed into crotonic acid and 3-hydrobutyric acid.

In relation to the intensity of absorption band at 1720 cm-1, nothing can be commented, since absorption intensity in ATR/FTIR spectroscopy depends largely on the quality of contact between sample and crystal[65].

3.3 Thermogravimetric analysis (TGA) Table 1 shows the values of temperature at the peak of the DTG curve (Tmax) and the temperature at which the sample loss 10% of its initial weight (T10%). A peak on a DTG curve characterizes the temperature at the highest rate of thermal degradation. Figure 6 depict the TGA curves of PHB before and after multiple extrusion cycles.

According to the literature[20,30,31,41], as PHB degradation proceeds, an ester chain with an unsaturated ester as end group, and an ester chain with a carboxylic acid as end group are formed. Each of these two types of ester may react at a distinct rate. However, COOH and COOR groups have very similar inductive effects[66], such that any structural changes suffered by PHB during extrusion cycles were not enough to change the thermal behavior of PHB. The TGA curves are all superposed and only a slight shift between them at the beginning of the curve could be verified (Figure 6).

No mass loss was observed until about 225 °C and the temperature at DTG peak (Tmax) was about 300 °C. The thermal decomposition of PHB takes place within a narrow temperature range, i. e., 235 ~ 315 °C. For all samples, the residual weight is lower than 0.5 wt%. Each TGA curve indicates a single step degradation, which means that the Polímeros, 27(2), 122-128, 2017

Figure 4. FTIR spectra of PHB as a function of each extrusion cycle.

Figure 5. FTIR spectra in stretching region of -C=O of PHB before and after each extrusion cycle.

Table 1. Temperature data from TGA curves of virgin and reprocessed PHB. Extrusion cycle

T10% (°C)*

Tmax (°C)**

0 1st 2nd 3rd

271.0 271.0 271.1 270.7

297.2 297.1 299.5 297.4

*T10% - temperature at 10 wt% of weight loss; **Tmax - temperature at the inflection point of the peak of DTG curve.

Figure 6. Thermogravimetric curves of PHB as a function of extrusion cycles.

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Rivas, L. F., Casarin, S. A., Nepomuceno, N. C., Alencar, M. I., Agnelli, J. A. M., Medeiros, E. S., Wanderley, A. O., No., Oliveira, M. P., Medeiros, A. M., & Santos, A. S. F. Table 2. Thermal properties of PHB as a function of extrusion cycles.

Extrusion cycle 0 1st 2nd 3rd

Xc (%) 57.7 58.8 62.7 62.2

Tm (°C) 174 172 172 172

thermal degradation of PHB polymer chains occurs by only one mechanism[39].

3.4 Differential scanning calorimetry (DSC) The results of the crystallinity degree (Xc) and melting temperature (Tm) are shown in Table 2.

According to these results, the crystallinity degree increased from 57.7 to 62.2%, this increase was more noticeable from the first to the second extrusion cycle. This evolution of the crystallinity with the number of extrusion cycles is also likely to be ascribed to a gradual decrease of the molecular weight of PHB which enhances the mobility and increases crystallization during the cooling step. This phenomenon is also known as chemicrystallization process[47]. Corroborating with this hypothesis, there is the tensile strength results that conversely decreased with extrusion cycles, indicating that the increase in crystallinity does not occur with the increase of tie molecules, responsible for transferring stress between two crystalline lamellae. The same behavior was observed for PLA at multiple injections cycle[5]. Furthermore, contributions of physical processes related to constraints imposed by amorphous chains[60] due to progressive crystallization process of PHB can be disregarded because each measurement was carried out at the same aging time.

The melting temperature decreases slightly at the first extrusion cycle and remains constant for further extrusion cycles. This modification in melting temperature reflects the formation of less perfect crystalline regions, which corroborates with the possibility of molecular weight reduction[25].

4. Conclusions Along multiple extrusion cycles, PHB suffered several changes such as, decrease in mechanical properties and increase in the degree of crystallinity. According to these results, a loss in tensile strength at break of PHB above 50% was observed at the third extrusion cycle. Nevertheless, that decrement in mechanical properties was significant at the secord cycle, i.e., ~ 40%. Also an increase in crystallinity degree was noticed mainly from the first to the second extrusion cycle, which probably could be associated to chemicrystallization process. FTIR results did not show any significant changes in polymeric structures associated to the formation of new chemical groups. Similarly, the thermal stability of PHB along processing cycles exhibited only a trend to decrease the thermal stability with extrusions paths, as evidenced by TGA curves. It is worth to note that the recyclability potential explored in this work refers to property retention up to three extrusion cycles under extreme recycling conditions (100% recycling with no added virgin polymer 126

and no additive or reinforcement). Therefore, improvement in PHB recyclability can be achieved by mixtures with virgin PHB, or incorporation of stabilizers and/or chain extenders to control its degradation during reprocessing.

5. Acknowledgements The authors gratefully acknowledge to PHB Industrial S.A. for PHB providing the polymer used in this work, to National Foundation for Science and Technology Development (CNPq) for the scholarships and profs.

Dr. Daniel D. Melo (UFRN), Dr. Edson N. Ito (UFRN) and Dr. Luiz Fernando Meneses Carvalho (IFPI-Teresina) for their kind help.

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Received: Dec. 11, 2015 Revised: May 23, 2016 Accepted: June 29, 2016

Polímeros, 27(2), 122-128, 2017

http://dx.doi.org/10.1590/0104-1428.2338

Influence of microcrystalline cellulose in thermoplastic starch/polyester blown films Mônica Oliveira Reis1*, Juliana Bonametti Olivato1, Juliano Zanela1, Fábio Yamashita1 and Maria Victoria Eiras Grossmann1 Departamento de Ciência e Tecnologia de Alimentos, Centro de Ciências Agrárias, Universidade Estadual de Londrina – UEL, Londrina, PR, Brazil

1

*monicareis09@gmail.com

Abstract This work investigated the influence of microcrystalline cellulose (MCC) in thermoplastic starch/poly (butylene adipate-co-terephthalate) films produced by blown extrusion, using different MCC contents (4, 7 and 10 g.100 g-1).

The films were characterised for their mechanical, structural and barrier properties. Increasing fibres concentration reduced the tensile strength (6.9 to 4.6 MPa), the elongation at break (568 to 147%) and weight loss in water (12.8 to 11.1%) of the films. The rigidity of the films increased from 19.8 MPa (without MCC) to 79.2 MPa in the samples with 10 g.100 g-1 of MCC. SEM images showed the occurrence of some agglomerates in this sample. The water vapour permeability of the films was not affected by the presence of MCC. The production of starch/PBAT/MCC films by blown extrusion was successful; however some adjustments are necessary to improve the dispersion of the particles at the polymeric matrix.

Keywords: extrusion, cellulosic fibres, biodegradable films, polyester.

1. Introduction In the last years, there was an increase in the researches focused in the development of biodegradable materials, due the growing accumulation of the conventional plastic materials, which are hard to be decomposed. Biodegradable polymeric packaging derived from cellulose, proteins and starch[1-4] have gained a great impulse because they allow the reduction of the use of materials derived from petroleum.

The starch is the most used agro-resource to produce biodegradable films due their biodegradability, low cost and wide availability[5]. According to Teixeira et al.[6] and Tang and Alavi[7], the native starch is not a real thermoplastic, but however, in the presence of a plasticiser, high temperatures and shear, it melts and flows, forming a material called thermoplastic starch (TPS).

The TPS, however, when used as a single polymer to produce biodegradable materials presents some restrictions that limit their use, including their hydrophilic character with high water vapour permeability and deficient mechanical properties, which are dependent of the relative humidity of the environment[7,8]. To overcome these drawbacks, TPS is frequently blended with biodegradable synthetic polymers[9-12], as poly (butylene adipate-co-terephthalate) (PBAT), an aliphatic-aromatic copolyester that combines desirable performance properties with biodegradability[13,14].

The inclusion of cellulosic fibres in biodegradable matrices has been the focus of numerous studies, with the aim to improve the mechanical and barrier properties of the materials[4,15-19]. Müller et al.[16] studied the effect of cellulose fibres (0.10, 0.30 and 0.50 g fibres/g starch) in the mechanical and physicochemical properties of starch-based films produced by casting. Their results showed that more resistant (8.39 MPa) and rigid (Young’s modulus of 217 MPa) films were produced with the addition of the fibres, but with a lower elongation at break (22%), when compared to

Polímeros, 27(2), 129-135, 2017

the material containing no fibres (1.59 MPa, 21 MPa and 83%, respectively).

Microcrystalline cellulose (MCC) presents as an alternative to the vegetal fibres used in biodegradable films. The MCC is a purified and insoluble cellulose produced by partial acid hydrolysis of the wood cellulose. The hydrolysis occurs in the amorphous areas of the polymeric chain, followed by the separation of the released microcrystals, which corresponds to a highly crystalline cellulose[20-22]. The use of MCC in different biodegradable polymeric materials has been studied[18,23-25].

Sun et al.[18] developed poly (vinyl alcohol) (PVA)/MCC composites by injection moulding. The Young’s modulus of the plasticised PVA increased from 204.8 MPa to 731.2 MPa, with the addition of 20% of MCC. The tensile strength increased from 37.8 MPa (PVA/MCC-0) to 46.5 MPa (PVA/MCC-20), which confirmed the strong interfacial interaction and good dispersion between MCC and PVA. While the injection moulding and the production of films by casting are less exigent process in relation to the structuration of the polymeric matrix and their interaction with the fillers, this does not occur with the blown film production. In this case, the requirements concerning the compatibility, dispersion, fibre/matrix adherence and the capacity of the fibres to transmit the tension along the material are greater, and the content of the added fibres has an important effect. To the best of our knowledge, our previous work[24] is the only study about blown films with MCC in blends of thermoplastic starch/polyester. The aim of this work was to produce biodegradable films based on thermoplastic starch/PBAT/MCC by blown extrusion and evaluate the influence of higher concentrations of MCC in their properties.

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Reis, M. O., Olivato, J. B., Zanela, J., Yamashita, F., & Grossmann, M. V. E.

2. Materials and Methods 2.1 Materials The blown films were produced with native cassava starch (17% wt amylose, 13 g.100 g-1 moisture), (Indemil, Guaíra, Brazil), poly (butylene adipate-co-terephthalate) (PBAT), supplied by BASF (Ludwigshafen, Germany) under the commercial name Ecoflex®, microcrystalline cellulose (MCC) M 101 Avicel®, supplied by Hexus Food Ingredients (Portão, Brazil), with a particle size of 10-15 µm and aspect ratio around 1, and glycerol, supplied by Dinâmica (Diadema, Brazil).

2.2 Methods 2.2.1 Blown films production The materials were processing using a laboratory co-rotating twin-screw extruder (BGM, model D20, Taboão da Serra, Brazil) with screw diameter (D) of 20 mm and length/diameter (L/D) ratio of 34. The TPS/MCC/PBAT pellets were produced in two steps. In the first step, TPS/MCC pellets were produced by dispersing the MCC in distilled water for 3 min with the aid of an Ultra Turrax homogeniser (MARCONI, model MA 102, Piracicaba, Brazil), which was then added to the starch and glycerol and manually mixed before extrusion. The mixture was pelletised using a temperature profile of 90/120/120/120/120 °C, screw speed of 100 rpm, and a matrix with five 2-mm holes. The venting ports were maintained closed to avoid water evaporation.

In the second step, the produced TPS/MCC pellets were processed with the PBAT with the same process conditions to produce TPS/MCC/PBAT pellets.

To obtain the blown films, TPS/MCC/PBAT pellets were fed to a single-screw extruder (BGM, model EL-25, Taboão da Serra, Brazil) composed of a 25-mm diameter screw with an L/D ratio of 30. The barrel temperature profile used was 90/120/120/130 °C for the four zones and 130 °C for the 50-mm film-blowing die and screw speed of 40 rpm.

The process conditions (screw speed and temperature profiles) were defined according to Reis et al.[24]. The pellets and the blown films were produced in duplicate.

The proportion of the components in each sample is presented in Table 1. The control sample (MCC0), without MCC, contained 56 g.100 g-1 of TPS (27 g glycerol/100 g starch) and 44 g.100 g-1 of PBAT. The concentrations of MCC (expressed as % w/w, TPS/PBAT basis) were selected in preliminary tests. The inclusion of water (2.4 mL/g MCC) was necessary to improve the dispersion of MCC and the used content was the minimum required to promote this

dispersion. Before the mixture with starch and glycerol, none free water was visible being the MCC hydrated and swollen.

The film thickness was controlled by the roll speed and the air-flow rate. The average thickness of the films was of 164 ± 17 to 195 ± 6 µm.

2.2.2 Mechanical properties A texture analyser model TA.XT2i (Stable Micro Systems, Goldaming, England) fitted with a 50 kg load cell was used to conduct the tensile tests of the films. Tensile tests were based on ASTM method D-882-02[26]. Ten samples from each formulation were cut along the longitudinal direction (50 mm in length and 20 mm in width) and fit in the tensile grips. The crosshead speed was set at 50 mm.min-1 and the initial distance between the grips was 30 mm. Before testing, the samples were conditioned at 23 ± 2 °C and 53 ± 2% relative humidity (RH) (saturated solution of Mg(NO3)2) for 48 h. The tensile strength (MPa), elongation at break (%) and Young’s modulus (MPa) were determined.

2.2.3 Density To determine the films density, three samples from each formulation with 20 mm × 20 mm were kept in a desiccator with anhydrous calcium chloride (CaCl2/0% RH) for 2 weeks and then, were weighed according to the procedure described

by Müller et al.[27].

2.2.4 Water Vapour Permeability (WVP) The tests were conducted using the American Society for Testing and Materials ASTM E-96-00[28] standard, with some modifications. Before the analysis, the samples were stored at 25 °C and 53% RH for 48 h. Each film sample was fixed in the circular opening of a permeation cell with a 60 mm internal diameter, and silicone grease was applied to ensure that humidity migration occurred only through the film. The interior of the cell was filled with a magnesium chloride solution (MgCl2/32.8% RH) and the device was stored at 25 °C in a desiccator to maintain a 42% RH gradient across the film. A saturated sodium chloride solution (NaCl) was used in the desiccator to provide 75.3% RH. The samples were weighed every 3 h during the 72 h of testing time. Changes in the weight of the cell or mass gain (m) were plotted as a function of time (t). The slope of the line was calculated by linear regression (R2> 0.98), and the water vapour permeation ratio (WVPR) was obtained with Equation 1: m  1  WVPR =   .   (1)  t   A

Table 1. Concentration of the components in the formulations.

Components Samples

TPSa

PBAT

MCC

Waterb

(g.100 g-1)

(g.100 g-1)

(g.100 g-1 TPS/PBAT)

(mL)

MCC0 56 44 0.0 0.0 MCC4 56 44 4.0 9.6 MCC7 56 44 7.0 16.8 MCC10 56 44 10.0 24.0 a Containing 27g glycerol/100g starch; bCalculated as 2.4 mL/g MCC; The numbers in the codes of the samples represent the level of added MCC.

130

Polímeros, 27(2), 129-135, 2017

Influence of microcrystalline cellulose in thermoplastic starch/polyester blown films where m/t is the angular coefficient of the curve and A is the sample permeation area.

The WVP (g.s-1.m-1.Pa-1) was calculated using Equation 2: = WVP WVPR.st / sp ( RH1 − RH 2 ) (2)

where st is the mean sample thickness (m), sp is the water vapour saturation pressure at the assay temperature (Pa), RH1 is the relative humidity of the desiccator and RH2 is the relative humidity in the interior of the permeation cell.

These tests were conducted in duplicate.

2.2.5 Weight Loss in Water (WLW) The weight loss in water was determined according to Olivato et al.[29]. Samples were previously dried for three days in a desiccator containing anhydrous CaCl2 (0% RH).

After weighing, the films were immersed in distilled water, maintaining a proportion of 30:1 (water/sample), for 48 h at 25 °C. The samples were then removed and dried at 105 °C for 4 h, and the weight of the conditioned specimen after treatment was used to determine the % weight loss in water.

These tests were conducted in triplicate.

2.2.6 Scanning Electron Microscopy (SEM) A scanning electron microscope FEI, model Quanta 200 (Hillsboro, USA) was used to observe the fractured surface of the blown film samples. The samples were submerged in liquid nitrogen and then broken (cryogenic fracture). Before coating with a gold layer, the samples were stored at 25 °C in a desiccator with silica gel (≈0% RH) for 3 days. The coating was produced with a Sputter Coater (BAL-TEC SCD 050).

Images were taken of the fractured surface and surface of films at a magnification of 1600x and 800x, respectively.

2.2.7 Statistical analysis The data were analysed using STATISTICA 8.0 software (Statsoft, Oklahoma), with analysis of variance (ANOVA) and Tukey’s test at a 5% significance level.

Comparing the samples without MCC (MCC0) and MCC4, no significant differences between these samples were observed, for all the analysed parameters.

The increase of MCC content from 4 to 10% wt reduces the tensile strength and elongation at break of the films, from 6.5 ± 0.1 to 4.6 ± 0.1 MPa and from 579 ± 27 to 147 ± 34%, respectively. This fact occurred as a result of the formation of cellulose agglomerates when higher concentration of this component was used, which reduced the reinforcing efficiency. The SEM images (Figure 1) confirm these results.

The rigidity of the material, represented for the Young’s modulus, increased from 19.8 MPa (MCC0) for 79.2 MPa (MCC10). This effect can be related to the formation of a network structure above the percolation threshold produced by the cellulose fibres by means of hydrogen bond interactions[18,30].

According to Ibrahim et al.[31], the Young’s modulus of the polymers may increase by the reinforcement effect of the cellulose fibres while the tensile strength may not be improved, or even be reduced, due the flocculation of the cellulose fibres, which is in accordance with the results of the present work.

Similar effects were reported by Ma et al.[32] for thermoplastic pea starch composites containing 12% wt of MCC. Mathew et al.[21] analysed the effect of MCC in poly (lactic acid) (PLA) matrix and observed that greater concentrations of MCC in the samples (from 10 to 25% wt) reduced the tensile strength and elongation at break and increased the Young’s modulus.

Considering the density of the films, no significant differences were recorded between the formulations. Despite the lower density of the fibres, the concentrations used in this work ranged from 4-10% wt and were not sufficient to affect the density of the materials. A possible cause of this is related to the absence of compaction of the polymeric chains with the inclusion of MCC (as can be observed in the Figure 1).

3.2 Water Vapour Permeability and Weight Loss in Water

3. Results and Discussions 3.1 Mechanical properties and density The mechanical properties and density of the films are expressed at Table 2. With a separate behaviour, the sample with greater proportions of MCC (MCC10) presented the lower tensile strength (4.6 ± 0.1 MPa) and elongation at break (147 ± 34%) and higher Young’s modulus (79.2 ± 13.6 MPa).

Opposite to the effects of MCC in the mechanical properties, the addition of these fibres did not influence the water vapour permeability (WVP) of the films, even at highest concentrations, as showed at Figure 2.

The presence of high concentrations of the fibres probably introduces a tortuous path, which makes more difficult the water diffusion through the polymeric matrix[33]. According to

Table 2. Mechanical properties and density of the films.

Properties Formulations*

Tensile strength

Elongation at break

Young’s modulus

Density

MCC0 MCC4

(MPa) 6.9 ± 0.3a 6.5 ± 0.1a

(%) 568 ± 53a 579 ± 27a

(MPa) 19.8 ± 0.6b 23.3 ± 0.9b

(g/cm3) 1.03 ± 0.01ª 1.04 ± 0.02ª

MCC7 MCC10

5.7 ± 0.6b 4.6 ± 0.1c

448 ± 86b 147 ± 34c

25.3 ± 1.4b 79.2 ± 13.6a

0.96 ± 0.06ª 0.98 ± 0.05a

*Numbers in the formulation codes are concerning MCC content (g.100g-1 TPS/PBAT); Results express in (mean ± standard deviation); Different letters in the same column indicate significant differences (p≤0.05) according to the Tukey’s test.

Polímeros, 27(2), 129-135, 2017

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Reis, M. O., Olivato, J. B., Zanela, J., Yamashita, F., & Grossmann, M. V. E.

Figure 1. SEM images of the fractures (1) (1600x magnification) and surfaces (2) (800x magnification) of the samples: (a) MCC0; (b) MCC4; (c) MCC7; (d) MCC10.

Dogan and McHugh[34], with the incorporation of cellulose fibres a reduction of the coefficient of water diffusion is expected since that the diffusion process into the matrix depends on the available pathways for the water molecules.

However, a slight reduction on the diffusion coefficient was not sufficient to change the water vapour permeability of the films, as evidenced in the results of the present work.

On the other hand, Kunanopparat et al.[35] reported that the addition of hemp and wood fibres in wheat gluten/glycerol composites reduced the water sensibility of the materials.

Also Ma et al.[32] found similar results for thermoplastic pea starch composites containing MCC. The various processing techniques (extrusion, compression moulding, and casting), polymeric matrices, type and concentrations of the fibres can explain the different effects of the cellulose fibres in the water vapour permeability of the composites.

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The Figure 2 presents the % of weight loss in water (WLW) of the samples. All the samples presented significant distinct results. While the formulation MCC0, without microcrystalline cellulose, exhibited a weight loss in water of 12.75 ± 0.09%, the inclusion of MCC in concentrations of 7 and 10% wt reduced the WLW of the films. This reduction was dependent of the concentration of the MCC at the formulations, with greatest values for WLW observed for the MCC10 (containing 10% wt of MCC). This behaviour could be associated to the less hydrophilic character of the cellulose, when compared to the starch[36,37]. The molecules of cellulose, due their linear structure, are able to perform intra and intermolecular hydrogen bonds, resulting in the formation of a crystalline structure which is totally insoluble in water[22,38]. In the samples containing greater proportions of MCC, the concentration of starch is lower, i.e., there Polímeros, 27(2), 129-135, 2017

Influence of microcrystalline cellulose in thermoplastic starch/polyester blown films molecules and the hydrophobic nature of the polymeric matrix. In the present work, the polymeric matrix contains thermoplastic starch (hydrophilic) and also PBAT, with hydrophobic character, which contributed to a low interfacial adhesion of the MCC in this matrix.

4. Conclusion The addition of MCC in starch/PBAT films was successful, considering the low thickness characteristic of the materials produced by blown extrusion, which makes more difficult the processability of the films, since the MCC particles represents more fragile points in the matrix structure.

Figure 2. Water Vapour Permeability (WVP) and Weight Loss in Water (WLW) of the films.

are a substitution of a more soluble molecule (starch) for a less soluble one (cellulose), which results in a reduction of WLW of the samples.

On the other hand, the MCC4 sample showed an opposite behaviour, i.e., the weight loss in water (13.18 ± 0.15%) was greater than the sample MCC0. A possible reason for this could be related to the easier removal of the MCC from the matrix structure, due to the starch solubilisation.

Thus, even being insoluble, MCC stayed dispersed at the water. This explanation is valid only for low MCC contents (≤ 5%), as reported by Reis et al.[24].

3.3 Scanning Electron Microscopy (SEM) The fracture and surface SEM images are presented at Figure 1, with magnification of 1600x and 800x, respectively.

Under high temperature, shear and in the presence of a plasticiser, the starch granules were completely disrupted and no residual granules could be observed at SEM images.

It is possible to notice a uniform dispersion of MCC in the starch/PBAT polymeric matrix, in the films with lower MCC concentration (Figure 1b.2 and c.2), however, the inclusion of MCC changed the films structure.

Greater concentrations of MCC added to the films led to the occurrence of micropores, as evidenced at SEM images (Figure 1c.1 and d.1). This effect was more pronounced in the films containing 7 and 10% wt of MCC (MCC7 and MCC10). Oriented beams were observed at surface images (Figure 1a.2, b.2 and c.2), corresponding to the MCC0, MCC4 and MCC7, respectively. In a different way, the beams disappeared at surface image of the sample MCC10 (Figure 1d.2), which, on the other hand, presented some aggregates.

In the Figure 1c.1 and d.1 images, some empty spaces were identified in the starch/PBAT matrix, which indicates weak or none interfacial adhesion between the matrix and MCC, when it was added in higher concentrations. A similar result was demonstrated by Mathew et al.[21], studying PLA/ MCC films, indicating no interfacial adhesion of MCC in the PLA matrix. Santos and Tavares[39] also developed PLA/MCC films and found in the SEM images that the dispersion of the MCC is relatively poor, due the self-aggregation of their Polímeros, 27(2), 129-135, 2017

Using greater proportions of MCC, a significant increase of the rigidity of starch/PBAT blown films was observed, but nevertheless, no contribution can be noticed in the tensile strength.

The occurrence of micropores, identified at SEM images, mostly at the samples with higher MCC content, did not affect the water vapour permeability of the films. The SEM images also showed that MCC tended to form some agglomerates when used at higher concentrations (10% wt).

The blown extrusion process could have influenced the results of the films, considering that most of the materials studied and produced with MCC frequently use other techniques, as casting or compression moulding, to obtain biodegradable sheets. So, to make viable the use of MCC in the blown films, some adjustments in the process must be focused in future studies, such as a treatment in the MCC particles to improve their dispersion/compatibility when used at content higher than 5% wt and, consequently, improve the tensile strength and water vapour permeability of the biodegradable films.

5. Acknowledgements The authors thank CAPES, CNPq and Fundação Araucária for financial and fellowship support.

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15. Follain, N., Joly, C., Dole, P., Roge, B., & Mathlouthi, M. (2006). Quaternary starch based blends: influence of fourth component addition to the starch/water/glycerol system.

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Received: Aug. 14, 2015 Revised: May 31, 2016 Accepted: June 29, 2016

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http://dx.doi.org/10.1590/0104-1428.2401

O Application of polyester derived from biomass in petroleum asphalt cement O Fernando de Araújo , Ingrid Souza Vieira da Silva and Daniel Pasquini * O Institute of Chemistry, Universidade Federal de Uberlândia – UFU, Uberlândia, MG, Brazil O O Abstract O This study evaluated the effects of the incorporation of a new additive to asphalt cement oil (CAP). A polyol product was obtained through the oxypropylation reaction of sugarcane bagasse. This polyol was polymerised with pyromellitic in order to obtain a polyester (BCP) to test its suitability in terms of the material properties to be applied as O anhydride additives. FTIR spectra of the polymerised material (BCP) confirmed the occurrence of chemical modification due to appearance of a new band at 1750 cm , characteristic of ester groups. The TGA data showed that the BCP product O the had higher thermal stability than the polyol. According to the softening point and elastic recovery tests, the incorporation O of 11% and 16% w/w BCP in conventional CAP met the specifications of regulatory standards.

O Keywords: petroleum asphalt cement, polyol, polyesters, sugarcane bagasse.

O 1. Introduction For most road applications, conventional asphalt exhibits As CAP constitutes 25 to 40% of the coating cost , behaviour which satisfies the requirements necessary it is feasible to study the applicability of the polymerised O good for proper performance in traffic and under normal weather material (BCP) as an alternative source to reduce the However, as the volume of commercial vehicles, production cost of the petrochemical CAP. Furthermore O conditions. weight, and increasing axle size, continues to grow year by it offers environmental benefits, due to the added value on special highways or airports, or in corridors with of the materials that until now, did not offer noble O year, heavy traffic and adverse weather conditions, with large applications. differences between winter and summer, it has O temperature become increasingly necessary to modify and improveing the properties of asphalt . Among the modifications that 2. Materials and Methods been investigated are natural asphalt, gilsonita or O have asphaltite, but especially polymers of various kinds that 2.1 Obtaining sugarcane bagasse 1

1

1

1

*pasquini@iqufu.ufu.br

-1

[1]

[1]

improve the performance of the binder.

A petroleum derivative, known as petroleum asphalt cement (CAP) in Brazil, is often used as the binder of the mineral aggregates. It is a semi-solid material, dark brown to black in colour, waterproof, viscoelastic, is slightly reactive and exhibits adhesive and thermoplastic properties.

The addition of polymers to the CAP tends to improve the viscoelastic properties, providing greater stability to the road surface[2]. The use of polymer-modified asphalts can reduce the frequency of maintenance and increase the life-time of local service roads which are difficult to access or suffer from a high-cost penalty for traffic interruption, should repairs be required[1]. Due to its low cost, abundance and easy availability, sugarcane bagasse (BC) is often used as a raw material for oxypropylation reactions with a view to producing a viscous polyol (BCO). This reaction makes the hydroxyls of the starting biomass (BC) more accessible for further reactions. Due to the high reactivity of anhydrides, the condensation polymerisation reaction of BCO and pyromellitic anhydride were carried out, which were aimed at producing an elastomeric copolymer type polyester (BCP) with desirable characteristics comparable to conventional CAP.

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The BC was obtained from a plantation located in Monte Alegre-MG. The sample was ground in a Willey mill and subsequently sieved through sieves with a 35 and 80 mesh.

The fractions collected between these screens, with dimensions of 0.425 by 0.180 mm, were used for the oxypropylation reaction.

2.2 Oxypropylation of sugarcane bagasse For the oxypropylation reaction, 5 g of the BC sample was added to 50 mL of an ethanol solution containing 0.5 g of KOH and then maintained in an oven for 12 h at 105 °C to allow for sample drying and solvent evaporation. After ethanol evaporation, 25 mL of propylene oxide (OP) was added to the BC sample in a 300 mL stainless steel autoclave.

The sealed autoclave, equipped with a thermocouple, a pressure gauge, and a heater controller system, was then heated at a heating rate of 5 °C min-1, to 200 °C while the increasing corresponding pressure was monitored. The finalization of the oxypropylation reaction was revealed by increasing the pressure followed by a reduction to atmospheric pressure which results in the total consumption of OP. The resulting material was a viscous polyol (BCO).

Polímeros, 27(2), 136-140, 2017

Application of polyester derived from biomass in petroleum asphalt cement 2.3 Procedure for the polymerisation of the elastomeric copolymer For this polymerisation reaction, 20 g of BCO, 1 g of sodium acetate (catalyst) and 8 g of pyromellitic anhydride (AP), were added to a stainless steel autoclave and allowed to react for 30 minutes at 200 °C. The elastomeric copolymer type polyester product was identified as BCP.

2.4 Analyses of Fourier transform infrared spectroscopy (FTIR) The samples of BC, BCO, AP and BCP were characterised by the FTIR analysis of KBr disks prepared with 1 mg of the sample per 100 mg of KBr, using a Shimadzu IR-21 PRESTIGE spectrometer. Infrared spectra were obtained in the range of 4000 to 400 cm-1, with a spectral resolution of 4 cm-1 and 32 scans.

2.5 Solubility test Solubility tests of BCP were performed in different solvents. For each solvent (25 mL) 0.15 g of BCP was added, and the mixture was kept for 24 hours at room temperature, after which its solubility was analysed. The solvents used were: acetone, water, chloroform, ethanol and hexane.

2.6 Thermogravimetric analysis (TGA) The thermal stabilities of the BC, BCO, AP and BCP samples were tested with a Shimadzu DTG-60H machine.

Around 5 to 7 mg of each sample was put into an aluminium pan and heated from 25 to 600 °C at a heating rate of 10 °C min-1 under a nitrogen atmosphere with a flow rate of 50 mL min-1.

2.7 Preparation of the modified CAP with BCP Analysis of the modified CAP with BCP as well as the characterization of conventional CAP (30/45) and the modified CAP were made in the quality control lab of BT Construction Ltd., located in Uberlândia MG. The BCP and CAP samples (30/45) supplied by the company were maintained at 105 °C for a period of 8 h to evaporate any residual water, which would interfere with the preparation procedure of the premixture of both. Then, both CAP and BCP were heated to approximately 150 °C on a heating plate in separate containers for approximately 40 minutes. 11%, 16% and 21% w/w of BCP, in relation to a total CAP mass of 600g, was added in order to evaluate how the incorporation of the additive would affect the properties of the modified CAP. Then, to prepare the premixture, both the CAP (30/45) and the incorporated BCP were subjected to vigorous stirring for approximately 40 minutes on a heating plate with a temperature controller set to 150 °C. The premix was stored in an oven at 150 °C for 8 hours before subsequent softening point and ductility test characterization.

2.8 Softening point test The softening point test refers to an empirical measure that correlates to the temperature at which materials, in this case asphalt, soften when heated under certain specific Polímeros, 27(2), 136-140, 2017

conditions and reach a predetermined flow level. In this assay, the samples were CAP (30/45), which is one of the materials used in the production of asphalt by the BT construction company, and was compared to the modified BCP-CAP prepared for this study.

The test samples in the form of ring and ball, were completed with the modified BCP-CAP. The test samples were kept at room temperature, approximately 25 °C, and were immersed in a glass beaker for 30 minutes. They were then placed on a controlled heating plate and the glass beaker containing the test samples was heated at a rate of 5 °C per minute. These assays were performed in duplicate and conformed to DNER 382/99 standards[3]. Conventional CAP (30/45) was subjected to the same analysis.

2.9 Ductility test Ductility refers to the ability of a material to stretch in the form of a filament. This test was performed to evaluate the cohesion of asphalt. To perform this test, both of the test samples were analysed with the CAP and the modified BCP-CAP, respectively. The test samples were maintained at room temperature, approximately 25 °C. Then, they were cut with tweezers so that the surface of both of the test samples were the same, thus ensuring that there were no alterations to the diameter of the samples, when stretched for the test run.

Prior to the ductility test, a preliminary step of correcting the density of the water present inside the machine was carried out, with sugar (sucrose), with respect to the density of the modified CAP, whose density was determined by pycnometry. Adjusting the density of the water present inside the machine above the density of the CAP was necessary to ensure complete immersion of the test samples in the machine, so that when stretched, the samples still remained fully immersed. The ductility metre was maintained at 25 °C and then the test samples were immersed in the equipment and were programmed to be stretched to 20 cm. When this 20 cm stretching value was reached, the equipment automatically disconnected.

The system was then left undisturbed for 90 minutes, after which the manual recoverable strain times for the immersed test samples were given in cm. After these measurements, the system was left undisturbed again for 60 minutes, and then another recoverable deformation cycle was measured.

The elastic recovery of each test sample was calculated as the ratio between the value found after the 60 minutes and the value recorded in the first measurement (20 cm). These assays were performed in duplicate and were performed in accordance with DNER 382/99 standards[3].

2.10 Proof body moulding To shape the test samples, the norm, as defined in DNER ME 043/95 was adopted[4], known as the Marshall dosage method of asphalt mixtures. A standard trace was used, as determined by the standards, to calculate the quantity of aggregates needed to shape the test samples using the modified BCP-CAP, as shown in Table 1. This approach was similarly applied for the CAP (30/45) samples.

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Araújo, F., Silva, I. S. V., & Pasquini, D. For the mixture detailed in Table 1, 4.8% w/w of conventional CAP or modified BCP-CAP, was added.

Marshall compaction was used in an effort consisting of 25 blows with a socket proctor, followed by the application of a static load of 5,000 pounds for two minutes, with the intention of levelling the test sample surfaces.

2.11 Proof body characterization 2.11.1 Percentage of empty volume This parameter is important for characterising the behaviour of asphalt mixtures because it has a significant influence on their cohesion and stiffness properties. The percentage of empty (% Vv) is defined as the relationship between the empty volume and the total volume of the mixture, and was calculated according to Equation 1[1]: % Emptyvolumes =

Theoretical density − Specific density Theoretical density

(1)

2.11.2 Percentage of minerals empty and Bitumen empty relation The percentage of mineral empty (VAM) represents the volume of intergranular empty spaces between the aggregate particles and the compacted mixture. This parameter was calculated according to Equation 2[1]: = VAM

%VV xVolumeCAP ×100 (2) Volumetotal

Additionally, this parameter was in-turn used to calculate a parameter called the Bitumen Empty Relation (RBV) defined in Equation 3[1]: R= BV

%VV ×100 (3) VAM

2.11.3 Fluency and stability This assay was performed for both of the test samples prepared with CAP (30/45) and the modified BCP-CAP in order to determine the mechanical parameters, fluency and stability. In this assay, the test samples were stored in a thermostatic bath with a temperature controller set to 60 °C for 40 minutes. Then, the test samples were placed in a compression mould used for this assay. The compression mould was coupled to the Marshall press to determine the maximum load which the test samples could withstand before rupture, defined as stability, as well as the vertical displacement corresponding to the application of the maximum load supported by test samples, defined as fluency[1].

3. Results and Discussion Figure 1 shows the FTIR spectra for the samples denoted BC, BCO, AP and BCP. The major peaks observed in the spectra were the BC bands at approximately 3300-3900 cm-1 which can be attributed to the OH stretch of cellulose, hemicellulose and lignin. The band observed at around 2800-3000 cm-1 can be attributed to aliphatic CH bonds, the band around 1675-1759 cm-1 is most likely associated with the carbonyl groups (C=O) present in the lignins and hemicelluloses, bands around 1000-1250 cm-1 which are characterisitic of C-O bonds that correspond to ethers link structures present in lignin, cellulose and hemicellulose, and a band at 1515 cm-1, which is indicative of the presence of lignin and can be attributed to the C=C vibration of aromatic rings[5]. For the oxypropylated BCO samples, an increase in CH aliphatic bands from of 2800 to 3000 cm-1 was observed, with the appearance of a new peak at 2970 cm-1 associated with methyl groups due to the grafted OP units.

The increase of the peak associated with the CO groups, as observed from 1000-1100 cm-1, can be attributed to the insertion of ether groups of OP. The FTIR spectra shown in Figure 1 confirm BC chemical modification when compared to the modified residue of BCO. These spectra also confirm the chemical modification process by oxypropylation, confirming the presence of the OP homopolymer formed during the oxypropylation reaction[6].

The AP FTIR spectrum, shown in Figure 1 shows AP characteristic bands near 1775 cm-1 which can be assigned to C=O stretching, and the regions 1225 cm-1 and 926 cm-1 can be associated with the C-O stretch of the cyclic anhydrides[7].

From the BCP spectrum it was possible to observe the appearance of a new peak at around 1759 cm-1,which confirms the polymerisation of BCO with AP, leading to the formation of the BCP polyester copolymer, due to the introduction of C=O groups, which can be observed for wavelengths associated with carbonyl esters.

From the thermograms of Figure 2, obtained in the thermogravimetric analysis, it is observed that BC has the higher comparative thermal stability, and the degradation profile is typical of a lignocellulosic material. The BCO has low thermal stability due to its viscous liquid state characteristics. After the polymerisation reaction, there was an increase in the stability of the resulting BCP polyester.

Table 1. Aggregate type for shape test samples of the modified BCP-CAP and CAP (30/45).

Aggregate type Crushed stone 3/4” Crushed stone 0 Stone powder Limestone

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Quantity (%) 18.0 32.0 45.0 5.0

Figure 1. FTIR spectra of BC, BCO, AP and BCP. Polímeros, 27(2), 136-140, 2017

Application of polyester derived from biomass in petroleum asphalt cement Degradation temperatures (T on set), obtained by the method of tangents in thermograms, were 260 °C for BC, 95 °C for BCO and 187 °C for BCP. In the solubility tests, conducted with BCP in different solvents, it was observed that BCP was not soluble in water, chloroform and hexane, and was partially soluble in ethanol and acetone. Solubility in water is a critical factor for possible applications of CAP formulations, as it essential to ensure that BCP will not solubilise with rain water, ensuring that premature repairs are not required. The solubility tests also showed that BCP is probably a crosslinked polymer due to its lack of solubility in the solvents tested. The results obtained, in terms of thermal stability and solubility, show that BCP has characteristics which are compatible with improved performance upon addition to CAP formulations.

The potential of BCP as an additive for conventional CAP was evaluated using the softening point and elastic recovery assays, and compared to conventional CAP. The values were compared with current standards (DNER 382/385)[3]. These parameters are considered essential in identifying the potential of the materials studied for the intended applications, and provide grants in accordance with the consulted standards (DNER 382/385)[3]. The incorporation of this additive would cause changes in these parameters, and the application of this additive would hamper use of the modified CAP for the desired application if the results were deemed not to align with the recognised standards.

Here, the effect on the properties of conventional CAP upon the incorporation of 11% and 16% w/w of BCP, as well as compliance of these parameters, in terms of aligning with the specifications of existing rules, was investigated and is reported. The addition of 21% w/w of BPC was incompatible with the CAP and a phase separation was observed. So it was not possible to carry out the tests for this composition.

Table 2 presents the results obtained during the characterisation of the conventional CAP and modified CAP with the additive BCP. According to the data shown in Table 2, the values found during tests, the BCP-CAP showed comparable values when compared to conventional CAP and followed the specifications of existing rules. Rose and co-workers reported that modified asphalt containing 12% recycled

rubber tyres had a softening point of 57.5 ºC and an elastic recovery of 79.8%[8]. D’Antona and co-workers reported that when 4% of a copolymer of ethylene and vinyl acetate (EVA) was added to conventional CAP, the softening point value was 62.3 °C[9].

Considering the above findings, the application of BCP as an additive does not prevent the use of modified CAP for the preparation of asphalt, and corroborates the data previously reported in the literature[8,9].

In addition, the test samples were moulded using the Marshall compactor manual following the norm according to DNER ME 043/95[4]. Figure 3 shows the test sample of BCP-CAP produced in cylindrical form.

Table 3 shows the results obtained from the Marshall test, which present the specific parameters used to characterise the proof bodies produced using the conventional CAP and the modified BCP-CAP. Tayfur and co-workers investigated the addition of 7% of a styrene butadiene copolymer (SB) and 7% of cellulose fibre (CF) to asphalts. The values found for stability were

Figure 2. Thermograms of BC, BCO and BCP samples.

Table 2. Parameters used in the characterisation of CAP. Softening Point (°C)

Elastic Recovery (%)

DNER 382/385 standards

60.0 ± 1.0

85.0 ± 1.0

Conventional CAP with 11% w/w BCP additive CAP with 16% w/w BCP additive

63.9 ± 3.0 64.0 ± 3.0 62.8 ± 3.0

83.7 ± 1.5 83.5 ± 1.5 83.8 ± 1.5

CAP

Figure 3. (a) BCP modified CAP proof bodies in cylindrical form - top view; (b) BCP modified CAP proof bodies in cylindrical form - side view.

Table 3. Marshall test results.

Parameters

Conventional CAP 30/45

Empty (%) Relation empty bitumen (%) Stability (kgf) Fluency (mm)

4.3 73.4 976 3.2

Polímeros, 27(2), 136-140, 2017

CAP Modified with BCP CAP Modified with BCP (11% w/w) 5.1 70.1 965 3.2

(16% w/w) 4.1 74.7 917 3.2

Standard Marshall 3-5 65-75 > 500 2-4.5

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Araújo, F., Silva, I. S. V., & Pasquini, D. 695 kgf for SB and 690 kgf for CF, the Vv to the percentage of SB was 3.80% and was 3.90% for CF, and finally the values found for fluency were 4.35 mm for SB and 3.90 mm for CF. In view of these results, the values found for the modified asphalt with BCP, in terms of stability, was on average 33% to 39% higher than the values found by Tayfur et al.[10]. The percentage of Vv was on average 8% to 32% higher than the values found by Tayfur et al.[10].

A study by D’Antona and co-workers reported on the incorporation of 4% of an ethylene and vinyl acetate copolymer (EVA) to conventional CAP, and reveals a Bitumen empty relation (%) value of 65.5%[9]. Thus, the modified asphalt with the BCP additive, in terms of this parameter, was on average 7% to 14% higher than the values found

by D’Antona et al..

In this sense, the comparative analysis, in terms of the parameters studied with the asphalt produced using the BCP additive, has properties superior to previously published data with other additives[9,10]. The comparative analysis of the data shown in Table 3 indicates that the incorporation of the BCP additive to CAP is in accordance with the specifications of the DNER ME 043/95 standards[4] (Marshall method).

4. Conclusion Based on these results it was observed that the oxypropylation reaction of BC was effective in producing a viscous polyol, BCO, which can be used in the synthesis of new polymers. This reaction can be regarded as a green chemistry process as all of the material that was added to the autoclave was removed as a product in the form of a polyol.

The FTIR results confirm the polymerisation reaction of BCO with AP, generating a type of polyester copolymer referred to as BCP. TGA tests showed that the BCP copolymer obtained exhibited increased thermal stability compared to the BCO monomer used in the polymerisation. Solubility tests indicate that the BCP is not soluble or is partially soluble in solvents with a wide range of polarity, indicating that BCP is most likely a crosslinked polymer.

It was possible to use BCP-CAP for the production of asphalt, and according to the specific characterisations used (Marshall test), the asphalt produced meets the specifications of regulatory standards, which enables the use of the new additive for asphalt production.

Whereas both mixtures with 11% and 16% w/w BCP presented results according to the standards, and considering the environmental aspect, we conclude that the addition of 16% w/w BCP is the best condition due to the fact of using a larger volume of renewable source of material as additve in CAP.

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5. Acknowledgements The authors acknowledge CNPq, FAPEMIG and Capes for financial support, and BT Construction forworking in partnership.

6. References 1. Bernucci, L. B., & Motta, L. M. G. (2008). Pavimentação asfáltica: formação básica para engenheiros. Rio de Janeiro: Universidade Petrobras.

2. Amaral, S. C. (2000). Estudo de misturas densas com agregados do estado do Pará, utilizando asfalto convencional (CAP-40) e asfalto modificado com polímero SBS (Betuflex B 65/60) (Master’s dissertation). Universidade de São Paulo, São Carlos.

3. Departamento Nacional de Estradas e Rodagem – DNER. (1999). Normas 382/1999: determinação da recuperação elástica de asfálticos modificados por polímeros, pelo método do ductilometro. Rio de Janeiro: DNER. 4. Departamento Nacional de Estradas e Rodagem – DNER. (1995). Normas 043/1995: misturas betuminosas a quente: ensaio Marshal. Rio de Janeiro: DNER. 5. Silvério, H. A., Flauzino, W. P., No., Dantas, N. O., & Pasquini, D. (2013). Extraction and characterization of cellulose nanocrystals from corncob for application as reinforcing agent in nanocomposites. Industrial Crops and Products, 44, 427-436. http://dx.doi.org/10.1016/j.indcrop.2012.10.014.

6. Menezes, A. J., Pasquini, D., Curvelo, A. A. S., & Gandini, A. (2009). Self-reinforced composites obtained by the partial oxypropylation of cellulose fibers. 1. Characterization of the materials obtained with different types of fibers. Carbohydrate Polymers, 76(3), 437-442. http://dx.doi.org/10.1016/j. carbpol.2008.11.006.

7. Oliveira, V. A. (2008). Síntese e caracterização de géis de acetato de celulose reticulados com dianidrido piromelítico e dianidrido do ácido 3,3´,4,4´benzofenona tetracarboxílico (Master’s dissertation). Universidade Federal de Ouro Preto, Ouro Preto.

8. Rosa, A. P. G., Santos, R. A., Crispim, F. A., & Riva, R. D. D. (2012). Análise comparativa entre asfalto modificado com borracha reciclada de pneus e asfalto modificado com polímeros. Teoria e Prática na Engenharia Civil, 20, 31-38.

Retrieved in 23 December 2016, from http://www.editoradunas. com.br/revistatpec/Art4_N20.pdf 9. D’Antona, D. M. G., & Frota, C. A. (2011). Estudo de Misturas Asfálticas com Ligante Modificado pelo Polímero EVA para Pavimentos Urbanos de Manaus - AM. Polímeros: Ciência e Tecnologia, 21(1), 13-18. http://dx.doi.org/10.1590/S010414282011005000007.

10. Tayfur, S., Ozen, H., & Atakan, A. (2007). Investigation of rutting performance of asphalt mixtures containing polymer modifiers. Carbohydrate Polymers, 21(2), 328-337. http:// dx.doi.org/10.1016/j.conbuildmat.2005.08.014.

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

Polímeros, 27(2), 136-140, 2017

http://dx.doi.org/10.1590/0104-1428.05616

Investigating the influence of conduit residues on polyurethane plates Rachel Faverzani Magnago1*, Nicolli Dayane Müller1, Mayara Martins1, Heloisa Regina Turatti Silva1, Paola Egert1 and Luciano Silva2 Una Produção, Construção e Agroindustria, Universidade do Sul de Santa Catarina – UNISUL, Palhoça, SC, Brazil 2 Arranjo Promotor de Inovação em Nanotecnologia, Universidade Federal de Santa Catarina – UFSC, Florianópolis, SC, Brazil 1

*rachel.magnago@unisul.br

Abstract Converting waste into a product similar to the original one or into another useful product is to save energy, protect natural resources, and bring back to the production cycle what was discarded. In that direction, new polyurethane-based composites have been developed by incorporating 5%, 10%, 15%, and 20% PVC conduit discarded by the construction industry. The objective of this study was to investigate the interaction between the phases of waste incorporation and the effect upon the new material properties. The samples were produced by the polycondensation process. Microstructural analysis revealed a reduction in pore size across the polymer matrix. However, there were no changes in thermal insulation, water absorption, compressive strength, and burning rate tests and in the thermogravimetric analysis and differential scanning calorimetry. The results from this study showed that the replacement of raw material by waste did not affect its properties.

Keywords: conduit, insulation, polyurethane, residue, thermal.

1. Introduction The supply chain of the construction industry is responsible for many positive impacts on the built environment, given that it enables the implementation of urban infrastructure and provides quality of life for its inhabitants. However, it consumes a significant amount of natural resources and generates tons of waste each year[1-4].

Construction waste consists of various types of materials, such as concrete, mortar, wood, and plastic, accounting for about 60% of the entire residue collected each year in Brazil. In general, these materials are used for embankment, which can cause solubilization or leaching of certain harmful substances present in these materials.

Furthermore, Mancini et al.[2,5-7] claims that plastics have low biodegradability and occupy a large amount of space, decreasing the useful life of landfill areas. The effects of degradation can appear after 20-30 years[5,8]. Building companies are improving their construction techniques and optimizing proper waste disposal to meet the current Brazilian legislation (CONAMA Resolution No. 307)[9].

The environmental impact of building materials can be reduced signicantly by increasing the amount of recycled materials.

Polyvinyl chloride is a thermoplastic polymer widely used in construction works as electrical conduit pipes. PVC conduit can be embedded, buried, or apparent, typically used in electrical installations for buildings, with a rated voltage not exceeding 1000 V for alternating current, and frequency below 400 Hz, or 1500 V for direct current. The conduit function is to protect and route electrical wiring in a building

Polímeros, 27(2), 141-150, 2017

structure. It should withstand temperatures from -5 °C to 60 °C for 24 hours without showing any deformity, not allowing the passage of electric current above 100 mA. PVC conduit was chosen as the focus of this research because the main properties of the material encompass thermal insulation, low flammability, and self‑extinguishing flame[6,10-12].

Polyurethane is a thermoplastic polymer often used in construction works due to its excellent performance as thermal and acoustic insulation[7,8,13-17]. The physico-chemical properties of polyurethane are important for that purpose[5].

The material has suitable mechanical strength to be used as thermal insulation in applications such as precast insulated sandwich wall panels, in ceiling lining or floors, and inside wood frame walls[18-21]. Singh[8], however, warns that this is not suitable for surface coating and interior finishes in buildings because it is not resistant to fire. The ignition of polyurethane foams includes all gas-phase processes that occur between the fuel production step and the occurrence of a visible hot flame. The ignition of polyurethane foams occurs by the interdiffusion of the flammable gases with air.

The basic physical and chemical aspects of gas-phase ignition reactions have been studied by several researchers[7,8,22].

Thus, the introduction of flame retardants is required in those situations[7,8,16,23]. Flame retardants are intended to inhibit or stop the polymer combustion process. Depending on their nature, they can act physically (by cooling, forming a dilution layer or fuel protection) or chemically (gas or condensed phase reaction). They can interfere with different processes involved in the polymer combustion (heating, pyrolysis, ignition, and propagation of thermal degradation)[13,16,23].

141

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

Magnago, R. F., Müller, N. D., Martins, M., Silva, H. R. T., Egert, P., & Silva, L. There is a large number of flame-retardant additives, which are divided into halogenated and non-halogenated compounds[3,13,16,23-28]. The halogenated flame retardants are more efficient than the non-halogenated composites, although they are more toxic during the burning process[23,26,27].

Halogenated flame retardants are organic compounds containing chlorine or bromine (or fluorine, or iodine). When they decompose, they interfere with the combustion process, both in the gaseous combustion and in the condensed phase.

However, fluorine- and iodine-based retardants are not commonly used because none of them interferes with the exact point of ignition of the combustion process[22,23,25,26].

In 2013, a tragedy of epic proportions was reported in the national and international media: the Kiss nightclub fire in the southern Brazilian city of Santa Maria[29]. This was not an isolated case, though. In 2003, there was a deadly Rhode Island nightclub fire in the United States as well. Several security failures were decisive for the former event, but it should be noted that polyurethane foam used for acoustic insulation has high flammability properties.

Flame retardants for polyurethane are generally utilized to retard ignition timing and reduce the burning rate and smoke formation. The type of flame retardant for PU depends on the application and rules that govern that application. Halogenated flame retardants are commonly used for rigid foams, and include tetrabromophthalic

anhydride, aliphatic chlorine, tetrabromophthalate, and 2,3-dibromo-2-butene-1,4-diol[16,25,30].

For that purpose, a polyurethane composite was developed with the incorporation of polyvinyl chloride (PVC) of flexible conduit residues[31] from the construction industry. PVC was incorporated into the foam components by simple mechanical mixing at the compounding stage.

This is still a relatively unexplored method, and this study is intended to contribute to analyze PVC compatibility with PU[32-40]. This paper presents an analysis of the conduit residue (PVC) incorporation into polyurethane to obtain plates, in order to investigate whether this material presents flammability reduction and improvement of thermal and mechanical properties.

2. Materials and Methods 2.1 Preparation of the composites PU/PVC Formulated polyether polyol, and toluene diisocyanate reagents were purchased from Arinos company (Arinos Química, São Paulo, Brazil) to manufacture the composite.

All reagents were used as received, and safety guidelines were followed.

Figure 1 exhibits a package of PVC conduit in yellow as marketed (Figure 1A). PVC conduit installation during construction generates residual cuts that are usually discarded.

After collection, the residues were washed, dried in an oven

Figure 1. (A) Flexible PVC conduit; (B) pieces of conduit cuts; (C) pieces for grinding; (D) triturated conduit and sieves for particle size classification; (E) conduit powder with particle size less than 1.18 mm.

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Investigating the influence of conduit residues on polyurethane plates at 40 °C for 48 hours, and ground in a blender (Figure 1B) and classified by particle size (Figure 1C). A number 16 ASTM stainless steel sieve was used for the classification of grain sizes. The material used for the production of composites presented particle sizes passing a 1.18 mm sieve (Figure 1D) and Mw 108.511 (Mw/Mn 2.233).

Molecular mass distribution of the PU and PVC was estimated by size exclusion chromatography (SEC), using a GPC Viscotek VE2001 with an ultraviolet detector at the wavelength of 280 nm and bimodal exclusion columns PVDF. Samples weight was 0.01-0.02 g and tetrahydrofuran was used as an eluent. The calibration of columns was carried out with a polystyrene standard of the Mw 30,000 and polydispersity Mw/Mn 1.0. The model polyurethane used as a standard was synthesized from the same polyether polyol and the same toluene diisocyanate (component molar ratio was 1:1.1, Table 1).

Specimens with a mass percentage of 5%, 10%, 15%, and 20% of PVC waste were prepared to substitute for polyurethane (polyol and isocyanate) with a total mass of 12.85 g. PVC was incorporated into the foam components by simple mechanical mixing at the compounding stage.

Table 1 shows the mass of the reagents used in each test.

The composites were obtained by mixing conduit residues (PVC) to formulated polyether polyol, and adding toluene diisocyanate. The mixture was stirred for 0.5 min and poured into a mold 0.05 m in diameter and 0.1 m height, cut in half vertically to facilitate demolding. A lid 0.05 m in diameter was used to close the molds. The specimens were demolded after 24 hours[38,41,42].

The curing time consisted in the reaction between hydroxyl (formulated polyether polyol) and isocyanate groups (isocyanate polymer), leading to the formation of a polymeric matrix

(Mw 278 and Mw/Mn 1.133). The chemical reaction that occurs between isocyanate and polyol forms a chemical bond by urethane links (-O-CO-NH-), which forms polyurethane because of polycondensation[14,43,44]. Isocyanate worked as a crosslinking agent, promoting solidification of the mixture, thus the proportions of the reactants were maintained for all composites[21,22]. Immediately after adding the curing agent, the mixture viscosity was suitable for being poured into the molds. The incorporation of the conduit residue occurred in different proportions with mass reduction of the reagents at baseline, being incorporated into PU. Figure 2 shows the PU/PVC composites with 5%, 10%, 15%, and 20% conduit. They had a uniform surface, good visual appearance, and showed no deformation after demolding.

2.2 Scanning Electron Microscopy (SEM) The micrographs of PU and PU/PVC plates were obtained by scanning electron microscopy using a JEOL JSM-6390LV device, at a voltage of 15 kV, and the samples were fractured in liquid nitrogen.

2.3 Water resistance test The procedure followed the gravimetric method recommended by the Standard Test Method for Water Absorption of Plastics D 570-98[45].

Initially, the samples were oven-dried at 50 °C for 24 hours. The dry material was weighed on a digital scale (Shimadzu BL-3200H) with a precision of 0.01 g. Then, the samples were immersed in a distilled-water bath at 25 °C for 24 hours. After that, they were removed and dried with paper towels and weighed again. The water absorption rate was obtained by the following formula:

Table 1. Amounts of raw materials for the production of PU = A specimens and PU/PVC composites with 5%, 10%, 15%, and 20% of PVC from conduit residues.

Poliol (g) PU PU/5PVC PU/10PVC PU/15PVC PU/20PVC

6.12 5.85 5.50 5.20 4.90

Toluene diisocyanate (g) 6.73 6.39 6.06 5.72 5.39

PVC residues (g) 0.64 1.29 1.93 2.56

Msat − Ms ×100 Ms

(1)

where: A: Water absorption (%); Msat: Mass of saturated sample (g); Ms: Mass of dry sample (g).

2.4 Thermal insulation The insulation test consisted of a qualitative evaluation for comparison of heating curves between the PU and the developed materials. A specimen made from the composite containing an internal cavity was developed to perform the assay. This apparatus allowed placing a beaker inside with 10 ml of previously cooled water, as shown in Figure 3.

A heating curve for water was then provided by using a digital thermometer (Hanna HL2221) at 10-minute time intervals, up to a maximum time of 60 minutes.

2.5 Mechanical compression strength

Figure 2. PU/PVC composites with 5%, 10%, 15%, and 20% conduit.

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Mechanical compressive strength tests followed the requirements of ASTM D695-15[47]. The mechanical tests were carried out by compressing 5 kN load cell in a universal testing equipment (EMIC DL-30000), according to the requirements of ASTM D-63890. In these tests, the specimens were subjected to pressure increments until the plastic deformation occurred, at room temperature.

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Magnago, R. F., Müller, N. D., Martins, M., Silva, H. R. T., Egert, P., & Silva, L.

Figure 3. Tests run and molds used. Source: Tudo Sobre Plasticos[46].

2.6 Differential Scanning Calorimetry (DSC) The analysis of differential scanning calorimetry was performed by using the Q2000 Thermal Analyst Instruments, Universal (TA Instruments). The temperature of the samples ranged from -50 °C to 150 °C, with heating-cooling rate of 5 °C min-1 under nitrogen atmosphere.

2.7 Thermogravimetry Thermogravimetric analysis (TGA) of samples was performed under nitrogen atmosphere at a heating rate of 10 °C/min within a temperature range of 20 °C to 900 °C, using a TGA Q5000 (TA Instruments) device.

2.8 Burning rate Tests to measuring the burning characteristics followed the requirements of NBR 9178:2015[48] and ASTM D3801.

Five samples of the composite measuring 102 mm in width × 356 mm in length × 13 mm in thickness were used for the tests. Each sample was placed in contact with the flame, and measurements were made of the time it took for the material to combust, and the distance traveled by the flame along the board. Figure 4 shows a schemating drawing of the test.

3. Results and Discussions 3.1 Scanning Electron Microscopy Scanning Electron Microscopy (SEM) analysis made it possible to evaluate the effect of incorporating conduit residues into the morphology of polyurethane foam. Figure 5 shows images of PU (5A and 5B) and images of PU containing 20% conduit residue (5C and 5D). There is no certified or standard micrograph for PU foams, given that they vary with the type and percentage of their components[3,4,27,28,31,49].

Therefore, the PU sample used as reference was prepared under the same conditions of the composite.

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Figure 4. Schematic drawing for polyurethane combustion test.

Figure 5 images show that cryogenically fractured surfaces exhibited a closed-cell structure, probably due to the type of cross-link during the polymerization reaction, which helps reduce cell disruption during the expansion process. Both the PU (Figure 5A and 5B) and the PU/PVC composite (Figure 5C and 5D) had an approximately spherical shape. There was a PVC dispersion in PU, and the interaction between the polymers occured at the interfacial level. There were hard PVC particles in the PU foam, which led to foam contraction[31,49]. Cell growth was hampered by the incorporation of the conduit as it can be seen by comparing 5A (PU) with 5C (PU/PVC), at a 15‑fold magnification, and 5B (PU) with 5D (PU/PVC), at a 30-fold magnification. PU cells have a diameter of about 0.800 mm (Figure 5A), whereas PU/20PVC composite cells (Figure 5C) have a diameter of about 0.300 mm, corresponding to an approximately 2.5-fold reduction. PU has cell size around 0.800 mm (5A), whereas in the image 5C there are up to three cells in 0.800 mm. The embedded PVC presented a maximum size of 1.18 mm, and the size reduction of PU foam cells was higher in the composite containing 20% PVC. There was good adhesion between PU and PVC resulting from intermolecular interactions between Cl groups (PVC) and N-H groups (PU)[50], given that PU foam was observed in PVC particle on the fracture surface, as circled in Figures 5C and 5D. Another indication of this interaction is that the images were obtained from the fracture surface, and there were no signs that the conduit was pulled out of the matrix at the time of the fracture. In fact, this homogeneity between conduit and matrix is essential for mechanical reinforcement[51].

The presence of predominantly closed-cell structures prevented moisture absorption, blocked the passage of gases and vapors, and did not communicate with the external environment. These are important characteristics for materials to be used as thermal insulation[52,53]. This behavior was maintained for all prepared PU/PVC composites, regardless of the percentage of conduit residues (PVC) incorporated into PU. Polímeros, 27(2), 141-150, 2017

Investigating the influence of conduit residues on polyurethane plates

Figure 5. SEM micrographs of freeze-fractured surfaces of PU (A and B) and PU with 20% conduit residue (C and D), at 15-fold (A and C) and 30-fold magnification (B and D), respectively.

3.2 Water Resistance Test This test was intended to determine the percentage of water absorbed by the material when subjected to immersion.

After weighing each specimen, water absorption percentage was obtained for each composite in the different conduit percentages, as shown in Table 2.

Considering that PU and composites presented predominantly closed-cell structures, as shown in the microstructure analysis, there was no significant variation in water absorption between the samples. Most study samples showed about 13% water absorption in the immersion test.

Then, when hold in environmental condition, they lost the absorbed water and returned to the original mass within 6 days, thereby not accumulating moisture. These results indicate that the material keeps its integrity even in contact with water[5,6].

3.3 Thermal insulation Figure 6 shows the results of the mass heating curve of water thermally isolated in PU and PU/PVC composites.

Figure 6 shows that all samples had the same thermal behavior, i.e., the incorporation of PVC residues did not alter the profile of the heating curve already shown by the PU apparatus. The temperature change rate for the mass of water was approximately 0.18 °C/min for all samples in Figure 6. Therefore, the incorporation of conduit residues in the percentage indicated in this study did not hamper the product’s performance related to thermal insulation as already shown by PU. Polímeros, 27(2), 141-150, 2017

Table 2. Average immersion test results for PU and PU-conduit composites (PU/PVC).

Sample PU PU/5PVC PU/10PVC PU/15PVC PU/20PVC

Water absorption percentage (%) 13.12 (0.18) 11.86 (0.22) 13.31 (0.17) 13.27 (0.18) 13.71 (0.19)

( ) standard deviation

These results confirm those from the microstructural analysis that identified the existence of closed-cell structures in the composites[3,4,27,28,31,49]. The existence of closed cells in the composite is particularly important for materials to be used as thermal insulation, given that conductivity occurs by radiation, which is hampered regardless of the cellular size[20]. These facts may explain the similar scores found for the heating curves of thermal-insulated water mass.

3.4 Mechanical compressive strength The effect of PVC on the compressive strength of the PU/PVC composites was studied. Figure 7 shows the compressive strength-deformation profile for the specimens of PU and PU/PVC composites.

The composites studied had an increase in the elasticity modulus and maximum strength as compared to the PU matrix.

These behaviors are due to the optimal level of compatibility reached between the components, which indicates that PVC 145

Magnago, R. F., Müller, N. D., Martins, M., Silva, H. R. T., Egert, P., & Silva, L. As it can be seen in Figure 7, the material is elastic in the initial stage of load application. This feature is important because the composite with elastic characteristics makes it difficult for the membrane of closed cells to break, which is vital for thermal insulation. Otherwise, air expansion within cell structures, for example, could facilitate thermal radiation[20].

3.5 Differential Scanning Calorimetry (DSC) Figure 8 shows the thermograms obtained for PU plate and PU/PVC composite with 15% conduit residue.

The curves for all blends show similar profile.

Figure 6. Profile of temperature variations of PU/PVC composites with 5%, 10%, 15%, and 20% conduit and pure PU.

Tg and Tm scores were similar for both PU plate and PU/PVC composite. For the PU, Tg was observed at -9.89 °C and Tm at 104.43 °C. Tg and Tm scores of the PU/PVC blend were similar to the PU component alone, since the interaction occurs at the interfacial level.

3.6 Thermogravimetry Figure 9 shows the TGA results for PU, PVC and PU/20PVC composite developed in this study. Degradation curves occurred in all samples at a temperature range between 20 °C and 900 °C. The heating rate was 10 °C/min with nitrogen atmosphere.

For the PU sample, mass loss began at 293.40 °C. At 325.00 °C, the mass loss was 45.43%, and complete degradation occurred at 770 °C. The first and second PU decomposition steps correspond to urethane bond-breaking and polyol decomposition, respectively[8,39,50,51].

Figure 7. Compressive strength-deformation profile for the specimens of PU/PVC composites.

has contributed to the mechanical reinforcement of PU/PVC composites, because PVC presented a uniform distribution in the PU matrix, as well as good interfacial interaction between the polymers[31,49]. The addition of PVC to the PU matrix causes two opposite intermolecular effects. Firstly, it reduces hydrogen interactions in the PU matrix, leading to a reduction in the mechanical properties of the material.

Secondly, the intermolecular interactions between the Cl groups (PVC) and NH groups (PU) led to an increase in the mechanical properties. The second effect prevailed, leading to a slight increase in the mechanical properties of the material. These results confirm what was observed in microscopy studies, since it shows good adhesion between the PVC and the PU matrix.

It is worth mentioning that the strength-deformation curves for the composites showed elastic and plastic characteristics regardless of the conduit percentage added to the PU matrix.

The elastic and plastic deformations are attributed to the presence of polyurethane, wherein PU/PVC composites were less elastic and had higher maximum strength scores as compared to those of PU. 146

Dehydrochlorination of PVC starts at the temperature the 32.87 °C, HCl is released from the polymer backbone in sequential stages, to yield long, conjugated polyenes.

At the temperature of 52.31 °C, there was a PVC mass loss of 57.68%, and at a temperature of 208.13 °C, mass loss was 87.49%[33,35,39], completing the decomposition around 300 °C. For the PU/PVC composite with 20% conduit residue, degradation also occurred in two stages. The degradation started at 22.5 °C and then a decrease in the values of degradation temperatures occurred, which can be attributed to the PU/PVC interface. Dehydrochlorination occurred mainly in the first stage of degradation and began at a lower temperature than that observed for PVC. HCl formation may have been catalyzed by PU amino groups[33,35,39,50]. At a temperature of 48.39 °C, there was a 51.39% loss of the total mass of the sample, which means that urethane bond-breaking also occurred at a lower temperature than that observed for the PU. There were two effects on the temperature scores for the composite degradation. The first effect consisted of the partial destruction of PU hydrogen bonds caused by the addition of PVC in the PU matrix. Thus, the temperature score for the PU degradation decreases by increasing the PVC content. The second effect is that the increase in the PVC content causes greater interaction between the Cl groups and NH groups, leading to catalyze C-Cl and C-NH bond breaking at lower temperatures. The results indicate that these effects led to decrease the degradation scores for the PU/PVC composite. Degradation of the composite samples was completed at 300 °C, whereas for the PU it occurred at 770 °C. Polímeros, 27(2), 141-150, 2017

Investigating the influence of conduit residues on polyurethane plates The interaction between polymers can influence the ignition and burning rate of the materials, given that HCl formation at a temperature lower than that of the PU decomposition act as a retardant in the gaseous phase of the composite burning. For that purpose, a vertical burning rate test was performed, which is designed for a diagnosis of real burning conditions.

3.7 Burning rate Figure 8. DSC thermogram for PU plate and PU/PVC composite with 15% conduit residue.

Figure 9. TGA spectrum for PU, PVC and PU/PVC composite with 20% conduit residue.

The composites under study were assessed according to the NBR 9178:2015[48] that regulates flexible polyurethane foam testing to determine burning characteristics and investigate the effect of conduit residue incorporation into the PU matrix. According to the NBR 9178:2015[48] and ASTM D3801[54], by comparing the burning times of the two composites it is possible to estimate burning speed and burning length, i.e., the distance traveled by the flame along the board.

Figure 10 exhibits a sequence of images (A-F) for the PU/20PVC composite burning, where (A) 2 s, (B) 4 s, (C) 5 s, (D) 6 s and (F) 7 s from the start of the burn.

In the burning test for PU plate, ignition was followed by self-sustained combustion, with complete burning of the specimen in 7 s, that is, the burning rate was 50.9 mm/s[13,54,55].

Figure 10 (A-F) exhibits the PU/20PVC composite burning test; the other samples showed similar behavior to the PU/20PVC. There was immediate ignition and complete burning of the specimens in the time range between 3 s and 8 s; thus, the burning rate ranged from 44.5 to 118.7 mm/s[56].

Figure 10. Burning test for the PU/20PVC composite, where (A) 2 s, (B) 4 s, (C) 5 s, (D) 6 s and (F) 7 s from the start of the burn.

Polímeros, 27(2), 141-150, 2017

147

Magnago, R. F., Müller, N. D., Martins, M., Silva, H. R. T., Egert, P., & Silva, L. The presence of PVC in the composite did not slow the burning rate.

According Vilar[57], when polyurethane foam is burned in the open air, a complete combustion occurs, forming carbon dioxide (CO2), nitrogen oxides (NOx), and water (H2O).

However, in a closed environment, such as in buildings, the effect of burning becomes even more dangerous because of the drop in atmospheric oxygen, which leads to an incomplete combustion and creates poisonous carbon monoxide (CO), in addition to producing hydrogen cyanide (HCN), also known as hydrocyanic acid or cyanide gas[26,58].

It is widely accepted that PVC dehydrochlorination involves a gradual reaction between chlorine and hydrogen atom neighbors in the polymer chain, comprising a double bond between the carbon atoms in positions in which the two atoms were originally connected, forming a structure of allylic chlorine with another chlorine atom of the polymer chain[11,33,35,39]. Complete combustion leads to the formation of hydrochloric acid, carbon dioxide and water.

The addition of PVC residues to polyurethane does not act as a flame retardant on the composite. Therefore, to meet the requirements of the NBR 9178:2015[48], a flame-retardant additive should be included in the formulation[11,51].

According to the NBR 9178:2015[48], flame-retardant additives must be added to coatings, such as polyurethane, for greater security of the environment and emergency evacuation.

The use of polyurethane in the construction industry is regulated by national norms, such as ABNT NBR 9178:2015[48], ABNT NBR ISO 15366-2:2006[59], ABNT NBR ISO 31000:2009[60], ABNT NBR ISO 14001:2004[61], and by international regulations, among which the Polyurethane Products in Fires: Acute Toxicity of Smoke and Fire Gases[62], Fire Safety Guidance: Working with Polyurethane Foam Products During New Construction Retrofit and Repair (AX-426)[63] and Model Building Code Fire Performance Requirements (AX-265)[64].

4. Conclusions This work presented an alternative to the use of conduit residue from the construction industry to produce a new material for thermal insulation. The composites were prepared by simple mixture in polyol and then polymerized, leading to materials that had thermal insulation behavior similar to that of PU. SEM images of the material exhibited a closed-cell structure, which revealed the thermal insulation property exhibited by the composites and explained water absorption with little variation between the samples. The composites presented elastic and plastic characteristics regardless of the percentage of conduit residues added to the PU matrix, but with enough mechanical strength to be used as thermal insulation. In the DSC thermograms, PU characteristics were predominant, whereas the Tg and Tm scores were similar for the PU/PVC composites. Findings from this study revealed that the addition of PVC (conduit) residues to polyurethane did not provide flame retardant properties to the composites as assessed by NBR 9178: 2015. Therefore, a flame-retardant additive must be included in the formulation to meet the requirements of the NBR 9178: 2015[48].

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5. Acknowledgements The authors wish to thank UNISUL for the financial and technical support to carry out this work. This work was awarded a scholarship from the Support Fund for Maintenance and Development of Higher Education (FUMDES) of Santa Catarina, Brazil, with funds provided under Article 171 of the State Constitution.

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comparative burning characteristics of solid plastics in a vertical position. West Conshohocken: ASTM. 55. Xu, Y., Tang, M., Chen, X., Chen, M., Yu, J., Ma, Y., Sun, Z., Zhang, Z., & Lv, J. (2015). Effect of phosphorus masterbatch on flame retardancy and thermal stability of polypropylene/ thermoplastic polyurethane blends. Polymers & Polymer Composites, 23(2), 113-119.

56. American Chemistry Council. (2016, March 11). Polyurethanes.

Retrieved in 13 May 2016, from http://polyurethane. americanchemistry.com/Health-Safety-and-Product-Stewardship 57. Vilar, W. D. (2004). Química e tecnologia dos poliuretanos. Rio de Janeiro: Reagentes & Fundamentos. Retrieved in 13 May 2016, from http://www.poliuretanos.com.br/ Cap1/11mercado. htm 58. Wang, J.-S., Wang, G.-H., Liu, Y., Jiao, Y.-H., & Liu, D. (2014).

Thermal stability, combustion behavior, and toxic gases in fire effluents of an intumescent flame-retarded polypropylene system. Industrial & Engineering Chemistry Research, 53(17), 6978-6984. http://dx.doi.org/10.1021/ie500262w.

59. Associação Brasileira de Normas Técnicas – ABNT. (2006).

ABNT NBR 15366-2: painéis industrializados com espuma rígida de poliuretano. Parte 2: classificação quanto à reação ao fogo. Rio de Janeiro: ABNT. 60. Associação Brasileira de Normas Técnicas – ABNT. (2009).

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ABNT NBR ISO 14001: sistemas de gestão ambiental: requisitos com orientações para uso. Rio de Janeiro: ABNT. 62. Landry, T. D., Pauluhn, J., Daems, D., & Reimann, K. A. (2002).

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Washington. Retrieved in 13 May 2016, from https://polyurethane. americanchemistry.com/Resources-and-Document-Library/11361. pdf Received: May 19, 2016 Revised: Sept. 13, 2016 Accepted: Sept. 16, 2016

Polímeros, 27(2), 141-150, 2017

http://dx.doi.org/10.1590/0104-1428.04216

Extraction and properties of starches from the non‑traditional vegetables Yam and Taro Luan Alberto Andrade1*, Natália Alves Barbosa1 and Joelma Pereira1 Department of Food Science, Universidade Federal de Lavras – UFLA, Lavras, MG, Brazil

1

*luandrade87@yahoo.com.br

Abstract The objective of this study was to assess the chemical, physical, morphological, crystalline and thermal properties of starch from two non-traditional vegetables, yam and taro. The analyses included proximate composition percent, amylose and mineral content, water absorption capacity, absolute density, morphological properties, X-ray diffractometry, thermal properties, pasting properties and infrared spectrum. The extracted starch exhibited a high purity level with low lipid, fiber and ash contents. The electron micrographs suggested that the taro starch granules were smaller than the yam starch granules. The results for the experimental conditions used in this study indicated that the studied starches differed, especially the amylose content, granule size and crystallinity degree and the pattern of the starches. Due to the high amylose content of yam starch, this type of starch can be used for film preparation, whereas the taro starch can be used as a fat substitute due to its small granule size.

Keywords: amylose, colocasia esculenta, dioscorea sp., FTIR, gelatinization.

1. Introduction Starch is an amylaceous product that is extracted from the edible parts of plants, especially cereals, tubercles, roots and rhizomes. In the food industry, this carbohydrate can be used to improve texture and as a thickener, colloidal stabilizer, gelling agent or volume and water retention agent[1].

The main sources of commercial starch worldwide are corn, potatoes and manioc. The composition and structure of starch granules vary considerably according to the botanical source, which affects its properties and functions[2]. Thus, there is a need for extensive research, especially on alternative starch sources, to assess the biochemical, functional, physical and chemical characteristics of the starches so that possible applications may be developed using substitutes for traditional starch sources and chemically modified starches.

Tropical countries, such as Brazil, have a great variety of endemic roots and tubercles, such as yam (Dioscorea sp.) and taro (Colocasia esculenta). These plants contain an average of 70% to 80% water, 16% to 24% starch and less than 4% lipids and proteins[3]. Because of the high starch content of yam and taro tubercles, which can reach 80% or more of the dry mass[4,5], these vegetables could represent alternative sources of starch.

The yam starches have an average granule size of 25 μm, with a type-C crystalline pattern and a relative crystallinity estimated at 34%[6,7]. Due to their high amylose content, yam starches can be used for film preparation[8]. In turn, according to Almeida et al.[9], native taro starch exhibits highly agglomerated round- and polyhedral-shaped granules.

This type of starch can be used as a fat substitute because of its small granule size, and it also has potential application in baby food formulation[8].

The purpose of this study was to assess the chemical, physical, morphological, crystalline and thermal properties of alternative starches extracted from two non-traditional

Polímeros, 27(2), 151-157, 2017

vegetables, yam and taro, that were cultivated in Brazil to determine their capacity for use as an alternative to traditional starches.

2. Materials and Methods 2.1 Starch extraction Approximately 4.07 g of anhydrous sodium metabisulfite was dissolved in 2.5 L of distilled water to preserve the material. One kilogram of previously chopped yam and taro was added to this solution in separate containers. The mixture was maintained under conditions of light and continuous agitation for subsequent grinding of the sample in 3 L of distilled water, using an industrial blender.

After grinding, the fluid paste was sieved using an 80-mesh sieve. The residue volume was measured, twice the amount of water was added, grinding was performed again, and the material was sieved.

The resulting liquid was filtered using a 200-mesh sieve and then allowed to rest for 24 hr for decantation before the supernatant was removed. Water was added to the starch layer, and then the mixture was centrifuged at 3,000 rpm.

After centrifugation, the supernatant was separated from the precipitate (starch), which was dried in a forced-air dryer at 40 °C.

2.2 Chemical and physical analysis 2.2.1 Determination of the composition percentage and amylose and mineral content of the yam (YS) and taro (TS) starches The starch moisture content was determined through a fast method using an infrared moisture analyzer (MOC-120H, Shimadzu, Brazil).

151

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

Andrade, L. A., Barbosa, N. A., & Pereira, J. The ether extract, crude protein, ash and carbohydrate fraction contents, which were low, were determined according to the Association of Analytical Communities[10].

The crude fiber content was determined using the method of Van de Kamer and Van Ginkel[11].

The amylose content was determined using a colorimetric method with iodine as described by McGrance et al.[12].

The mineral levels were determined according to the techniques described by Malavolta et al.[13].

thermogravimetric analyzer (Shimadzu) at temperatures ranging from 30 °C to 600 °C and scanning rates of 10 °C min-1; the analyses were also run under N2 atmosphere with a flow rate of 50 mL min-1.

2.6 Pasting properties

2.2.2 Water absorption capacity (WAC)

To determine the apparent paste viscosity, a Rapid Visco Analyzer (RVA - Newport Scientific, Sydney, Warriewood, Australia) was used according to the International Association for Cereal Science and Technology (ICC) standard method no. 162[17].

Beuchat’s[14] method was used to determine the water absorption capacity of the extracted starches.

2.7 Infrared spectrum

2.2.3 Absolute density The absolute density of the YS and TS was determined by xylene displacement according to the methodology recommended by Schoch and Leach[15].

2.3 Morphological properties After drying and grinding, the powder samples of YS and TS were deposited on double-sided carbon tape, placed on racks covered with aluminum foil and sputter-coated with gold (Balzers Sputter Coater SCD 050). At the end of this procedure, the samples were examined under a scanning electron microscope. The generated images were scanned at varying magnifications at 20.00 kV and work distance between 8.0 and 9.0 mm.

After the samples were dried in a forced-air oven, the YS and TS were analyzed by Fourier transform infrared spectroscopy (FTIR) in a Digilab Excalibur FTS 3000 spectrometer (USA) with a deuterated triglycine sulfate (DTGS) detector using a spectral range between 4000 cm-1 and 400 cm-1 and a resolution of 4 cm-1. The samples were analyzed using the transmission in potassium bromide (KBr) pellets that were 7 mm in diameter.

2.8 Statistical analysis A descriptive statistical analysis was performed to obtain the mean and standard deviation of the three previously analyzed replicates.

3. Results and Discussion

2.4 X-Ray diffraction (XRD)

3.1 Chemical composition

The starches were characterized using the XRD powder method with PANalytical equipment, model X’Pert Pro, a 2θ angular variation of 4° to 70°, CoKα radiation and a scanning rate of 5° min-1. The spacing d was calculated using the Bragg equations (nλ = 2dsinθ, where d is the spacing, n = 1, and λ = 1.7889 Å). From the recorded diffractogram, the crystallinity index was calculated.

Table 1 presents the chemical composition of the isolated yam (YS) and taro (TS) starches.

2.5 Thermal properties The thermal characteristics of the isolated starches were studied using a differential scanning calorimeter (DSC - Auto Differential Scanning Calorimeter DSC-60A, Shimadzu) equipped with a thermal analysis station. Initially, a starch suspension was prepared with 70% distilled water and 30% of each studied starch. The samples were hermetically sealed in pots and allowed to rest for 1 h at room temperature before the samples were heated in the DSC. The temperature of the DSC analyzer was calibrated using tin and indium standards, and an empty aluminum pot was used as a reference. The samples were heated at a rate of 10 °C min-1 from 25 °C to 120 °C. The onset temperature (To), peak temperature (Tp), conclusion temperature (Tc) and enthalpy of gelatinization (ΔHgel) were automatically calculated. The gelatinization temperature range (R) was calculated as (Tc - To). The peak height index (PHI) was calculated using the ratio ΔH (Tp - To)-1, as described by Krueger et al.[16].

Thermogravimetric analysis (TGA) and differential thermal analysis (DTA) were performed in a DTG-60H 152

According to Peroni et al.[2], the purity of a starch is related to its chemical composition, and pure starches generally have low levels of ash, protein and lipids. Therefore, the ether extract, ash and crude fiber contents should have been lower than 1.00 g 100 g-1 according to the literature[2,18-21].

The crude protein content was higher than the values in some studies[2,20,21], although the differences were not significant.

The high protein content may be due to the strong bond of the starch with the mucilage present in this vegetable[8].

Because of the low observed levels, especially for the ether extract, ash and crude fiber contents, and the moderate levels observed for protein, the extracted starch was determined to have a high purity of more than 96 g 100 g-1.

The moisture contents were lower than 14 g 100 g-1, which is generally acceptable for dry products and a desirable shelf life.

The amylose content in starch has an important effect on its functional properties. According to Zuluaga et al.[6], the high amylose content in yam starch (28.8 g 100 g-1) was one of the factors associated with the high level of retrograded starch.

The amylose content of YS was higher than that of TS; therefore, these starches can exhibit different functional properties. According to the literature, YS contains between 21.12 g 100 g-1 and 22.48 g 100 g-1[18] and between 13.58 g 100 g-1 and 20.05 g 100 g-1[19] of amylose, and these values are Polímeros, 27(2), 151-157, 2017

Extraction and properties of starches from the non-traditional vegetables Yam and Taro Table 1. Moisture, ether extract, crude protein, ash, crude fiber, amylose, macro-and micro-mineral content and carbohydrate fraction of yam (YS) and taro (TS) starches.

Starch YS TS Starch YS TS

M

EEa

CPa

10.9 ± 0.10 10.5 ± 0.20 P

0.44 ± 0.08 0.43 ± 0.09 K

0.04 0.05

0.00 0.00

2.44 ± 0.31 1.46 ± 0.10 Ca g 100 g-1 0.00 0.09

Aa g 100 g-1 0.26 ± 0.00 0.98 ± 0.05 Mg

CFa

CFra

0.65 ± 0.16 0.39 ± 0.01 S

96.22 ± 0.29 96.75 ± 0.09 Cu

0.00 0.00

0.00 0.00

14.6 9.5

Amylosea 37.46 ± 0.48 19.37 ± 0.93 Mn Zn mg kg-1 0.00 0.02 0.00 10.4

Fe 67.00 70.80

M: moisture; EE: ether extract; CP: crude protein; A: ash; CF: crude fiber; CFr: carbohydrate fraction; YS: yam starch; TS: taro starch; adry basis.

The results reported are the means of triplicate samples.

lower than those found in the present study. The composition of TS is also similar to the composition reported in most studies of Colocasia esculenta starch, with amylose values of 18.4 g 100 g-1[22] or from 13.5 to 27.66 g 100 g-1[23].

The differences found for YS may result from the different methods used for amylose extraction and measurement.

The contents of the main minerals present in YS and TS are listed in Table 1, which shows the absence of K, Mg, Mn and S for both starches and Ca for YS. According to Tester et al.[24], starches contain small amounts of minerals such as Ca, Mg, P, K and Na; however, only P is functionally significant. The P from the starch of tubercles, such as potatoes, usually occurs in the form of the negatively charged phosphomonoester. The ionic repulsion generated by these groups weakens the intermolecular association forces, thus increasing its water-binding capacity, swelling power and paste clarity. This factor can also result in higher starch viscosity, thus leading to greater gel strength[25]. The P contents reported in the literature for yam and taro starches are approximately 0.022 g 100 g-1 and 0.01 g 100 g-1, respectively[2,20], which are lower than the values found in this study.

Because of the different isolation methods used to determine the starch content, the chemical and mineral composition depends on the botanical source and types of fertilization as well as on the extraction method[20].

3.3 Morphological properties

3.2 Physical properties

3.4 X-Ray diffraction

3.2.1 Water absorption capacity (WAC) and absolute density The WAC values found for YS and TS (323.18 ± 0.99% and 333.28 ± 1.10%, respectively) are consistent with the values reported in the literature[17,23]. According to Hoover and Sosulski, 1986 (cited by Shujun et al.[26]), the lower water absorption capacity of yam starch, compared with taro starch, shows that there is a higher proportion of hydroxyl groups in the formation of hydrogen and covalent bonds between the starch chains than with water.

Starch density is fundamentally important for technological purposes. This parameter represents the content of material per unit of real volume occupied by the material[15]. The density of YS (1.46 ± 0.02 g cm-3) was higher than the TS density (1.31 ± 0.03 g cm-3), and these density values are similar to the values reported in the study by Deepika et al.[18] as well as in the study by Deepika et al.[23] The whole granule content of YS is most likely higher than that of TS, as indicated by its higher absolute density.

Polímeros, 27(2), 151-157, 2017

Figure 1 displays the scanning electron micrographs of the starches extracted from yam and taro.

The YS granules are significantly larger than the TS granules, with the YS granules varying between 15.51 and 30.47 µm and the TS granules varying between 2.273 and 3.986 µm.

Such differences between the starch granules are responsible for their different properties.

Electron micrographs B and D confirm the main elliptical shape of the YS granules. Dioscorea starch, which was studied by Deepika et al.[18], contained oval- and ellipsoidal-shaped granules with diameters between 5 and 10 μm; these values are lower than the diameters found in the present study. The morphology and size of the starch granules depend on the starch biosynthesis and on the plant physiology, which affects the light transmittance, amylose content, water absorption capacity and swelling power[27].

The TS granules are widely agglomerated, and the main shape is polyhedral with both circular and irregular shapes (as shown in Figures 1A and 1C). In a study on Mexican taro[28], the starch granules clearly exhibited a mixture of shapes, such as oval, spherical polygonal and irregular shapes, with granule sizes between 1 mm and 5 mm, which is consistent with the present study.

The XRD patterns of YS and TS are presented in Figure 2. The corresponding calculated crystallinity levels were 32.88% and 44.66% for YS and TS, respectively.

The crystallinity degree of the starch is close to that found in certain studies[19,22]. In accordance with the literature, the TS had a lower amylose percentage (19.37 g 100 g-1) and a higher amylopectin content and crystallinity index[19]; this finding indicates that the crystallinity is primarily responsible for the amylopectin content.

The TS exhibited a type-A crystallinity pattern (Zobel 1988) that is similar to that of the starches studied by Sit et al.[22], Zeng et al.[21], Sukhija et al.[8] and Agama‑Acevedo et al.[28] In the type-A pattern, double helices are packed in a monoclinic unit cell and form a densely packed structure with only four water molecules per cell[29].

Unlike TS, YS exhibited a type-C crystallinity pattern, which Gernat et al.[30] reported to be a mixture of patterns A and B; thus, the crystalline regions of this starch contain a fraction of each pattern. According to Jiang et al.[19], out of the five 153

Andrade, L. A., Barbosa, N. A., & Pereira, J.

Figure 1. Scanning electron micrographs of native taro and yam starches. (A) and (C) taro starch granules at lower and higher magnifications, respectively; (B) and (D) yam starch granules at lower and higher magnifications, respectively.

32.88% crystallinity was found for YS, 44.66% crystallinity was found for TS, and higher enthalpy was found for YS compared to TS (68.28 J g-1 versus 25.21 J g-1).

The differences in the gelatinization temperature may be associated with differences in the amylose content, shape, distribution and water-binding capacity of the starch granules. When a majority of the hydroxyl groups within a granule are involved in the formation of hydrogen and covalent bonds between starch chains instead of with water molecules, a decrease in the WAC can occur, which was observed. Furthermore, an increase in gelatinization enthalpy may also occur[31], which would explain the high value for the gelatinization enthalpy.

Figure 2. X-Ray diffractograms of yam (YS) and taro (TS) starches.

studied Dioscorea varieties, four corresponded to type C, and only one corresponded to type A. Lan et al.[7] also showed a type C crystallinity pattern for Chinese yam.

3.5 Thermal properties The thermal properties of YS and TS are displayed in Table 2. The To, Tp and Tc of the studied starches were very different, with TS gelatinization starting at a temperature that was 12.7 °C lower than the temperature observed for YS. Therefore, when gelatinization must be started and completed at lower temperatures, the use of TS is recommended over YS. Compared with the values found in the literature, the To, Tp and Tc temperatures for TS and YS were lower, and the enthalpy was much higher[19,28].

Zuluaga et al.[6] studied yam starch (potato and corn starch were the standards) and observed that the gelatinization enthalpy increased with a decrease in crystallinity degree, which corroborated the findings of the present study, in which 154

PHI measures the uniformity of gelatinization, which was much higher for YS and indicated differences in the internal organization of the two starch granules. The R values (gelatinization range) were similar and high for both starches but lower than the value reported for the D. opposita Thunb.

YS in the study by Deepika et al.[18].

Figure 3 shows the TGA and DTA curves. For both starches, until temperatures close to 100 °C were reached, only 11.10% and 8.01% water was lost, respectively, which was indicated by the endothermic peak (descending).

Until approximately 250 °C, the starch did not depolymerize with increasing temperature; therefore, the starch was stable and did not exhibit alterations in its structure and properties.

At the second temperature drop, the entire process became endothermic according to the DTA curve. The final masses at 600 °C were 14.29 g 100 g-1 and 14.40 g 100 g-1 for YS and TS, respectively.

3.6 Pasting properties When a starch suspension is heated at a constant speed, the viscosity gradually increases until a maximum value is reached[28]. The pasting properties of the starches are summarized in Table 2.

Polímeros, 27(2), 151-157, 2017

Extraction and properties of starches from the non-traditional vegetables Yam and Taro Table 2. Thermal and pasting properties of the starch extracted from yam (YS) and taro (TS).

Starch

To (°C)

Tp (°C)

Tc (°C)

R

PHI

ΔHgel (J g-1)

YS TS Starch

62.54 49.84 A

71.63 56.42 Tip (°C)

78.36 66.64 B

15.82 16.80 C

7.51 3.83 Breakdown

68.28 25.21 Setback

FVa

TS YS

157.5 165.0

77 73

4042.0 3242.5

2651.0 1723.5

1391.0 1519.0

1248.0 741.5

3899 2465

To: onset temperature; Tp, peak temperature; Tc: conclusion temperature; R: gelatinization range, (Tc–To); PHI: peak height index, ΔHgel (Tp–To)-1; ΔHgel: enthalpy of gelatinization; YS: yam starch; TS: taro starch. aValues in cP; Tip: initial paste temperature. Breakdown = viscosity peak - B; Setback = final viscosity - B. A: maximum viscosity at 25 °C; B: maximum viscosity at 95 °C; C: minimum viscosity at cooling. FV: Final viscosity. The results reported are the means of duplicate samples.

Figure 3. TGA and DTA curves for yam (YS) and taro (TS) starches, where endo stands for endothermic.

The starch viscosity increased when the temperature reached approximately 73 °C and 77 °C for YS and TS, respectively; this temperature is known as the initial pasting temperature. The observed values were lower than those observed in the study of Mexican taro[28] and the study of four out of the five different Dioscorea L. species by Jiang et al.[19] Starch granules that have low amylose content can swell at a higher temperature[32]. However, because the amylose contents in the present study were relatively high, especially for YS, the suspension temperature required for granule swelling was lower.

For both starches, the peak viscosity was high (3242.5 cP and 4042.0 cP) and occurred before the constant temperature of 95 °C was reached. This peak reflects the maximum capacity of the starch to swell freely before physical collapse.

The pasting properties, such as the setback and breakdown, can be influenced by the amylose and lipid contents, and the high amylose content of YS and TS resulted in high values of these two properties[19,33]. Such values demonstrate the high retrogradation tendency in these starches, especially in TS.

3.7 Infrared spectrum (FTIR) The isolated starches were assessed using FTIR to identify functional groups in the starches, with the goal of determining the starch structure (Figure 4). The wide band at 3,380 cm-1 for YS and TS can be attributed to the OH Polímeros, 27(2), 151-157, 2017

Figure 4. Infrared spectra (FTIR) of yam (YS) and taro (TS) starches.

bond, and its width is related to the formation of inter- and intramolecular hydrogen bonds[34]. Bands at 2,931 cm-1 can be attributed to the axial deformation of the C-H bond, which is found in the region between 3,000 and 2,800 cm-1.

Changes in the intensity of the band at 2,931 cm-1 can be caused by variations in the amylose and amylopectin contents of the starches[35]. In the present study, the intensity was higher for TS, which contains less amylose and has a higher crystallinity degree.

155

Andrade, L. A., Barbosa, N. A., & Pereira, J. The band at 1,163 cm-1 corresponds to the coupling of the C-O and C-C stretching mode, and the band at 1,094 cm-1 can be attributed to C-O-H bending.

The vibrational bands (bending and deformation) related to C and H atoms can be observed in the region from 1,500 cm-1 to 1,300 cm-1.

According to Zeng et al.[21], who studied taro starch cultivated in China, the bands at 1,081 cm-1 and 1,019 cm-1 are characteristic of the O-C stretching associated with anhydrous glucose, which can be observed at 1,090 cm-1 and 994 cm-1 for YS and at 1,083 cm-1 and 1,021 cm-1 for TS. According to Kizil et al.[35], the band observed at 1,642 cm-1 can be associated with the water adsorbed in the amorphous region of the starch granules. Because this band is related to the starch crystallinity, the variations in the crystallinity of different starches can affect this band. In the present study, the bands at 1,641 cm-1 and 1,651 cm-1 have different intensities and displacements, which can indicate the difference in the crystallinity patterns of the yam and taro starches, as shown by the XRD. The prominent band at 931 cm-1 corresponds to the water sensitivity and indicates the presence of this molecule in the starch structure[18].

3.8 Suggestions for application of YS and TS Due to the small starch granule size observed for taro (2.273 and 3.986 µm), this type of starch can be used in products that require smaller particle sizes, including as a fat substitute or as flavor or substance carriers[36].

According to Liporacci et al.[37], yam starch is a promising polymer for film preparation due to its high content of amylose (30%), which is a polymer responsible for the film-forming property of starch. In the present study, the amylose content was even higher (37.46 ± 0.48 g 100 g-1), indicating that this yam starch has the potential to be used in the preparation of starch-based biofilms.

4. Conclusions The studied starches were found to exhibit great differences, especially in their amylose content, granule size and crystallinity degree and pattern. TS exhibited a higher retrogradation tendency, whereas YS displayed a high enthalpy of gelatinization and exhibited a higher degree of interaction between its chains. Because of the differences in the measured properties of the studied starches, the two starches can have different applications in the field of food science and technology. YS can be used in film preparation, whereas TS can be used in products that require smaller particle sizes, for example, as a fat substitute.

5. Acknowledgements The authors would like to thank the Brazilian Federal Agency for Support and Evaluation of Graduate Education (Coordenação de Aperfeiçoamento de Pessoal de Nível Superior – CAPES), the National Counsel of Technological and Scientific Development (Conselho Nacional de 156

Desenvolvimento Científico e Tecnológico – CNPq) and the Minas Gerais Research Foundation (Fundação de Amparo à Pesquisa do Estado de Minas Gerais - Fapemig) for financial support. We would also like to thank the Laboratory of Electronic Microscopy and Ultrastructural Analysis (Laboratório de Microscopia Eletrônica e Análise Ultraestrutural – LME) of the Federal University of Lavras (Universidade Federal de Lavras – UFLA).

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Retrieved in 04 April 2016, from http://www.cabdirect.org/ abstracts/19531400026.html;jsessionid=A4788EDF3D603D 51AA1A7575FDBC3DCA# 12. McGrance, S. J., Cornell, H. J., & Rix, C. J. (1998). A simple and rapid colorimetric method for the determination of amylose in starch products. Stärke, 50(4), 158-163. http:// dx.doi.org/10.1002/(SICI)1521-379X(199804)50:4<158::AIDSTAR158>3.0.CO;2-7.

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15. Schoch, T., & Leach, W. (1964). Determination of absolute density; liquid displacement. In R. Whistler & M. Wolfrom (Eds.), Methods in carbohydrates chemistry. Zürich: Academic Press. 101 p.

16. Krueger, B. R., Knutson, C. A., Inglett, G. E., & Walker, C. (1987). A differential scanning calorimetry study on the effect of annealing on gelatinization behavior of corn starch.

Journal of Food Science, 52(3), 715-718. http://dx.doi. org/10.1111/j.1365-2621.1987.tb06709.x.

17. International Association for Cereal Science and Technolog – ICC. (1995). Rapid pasting method using the Newport rapid Visco analyser (Standard Methods, no. 162, pp. 123-456).

Vienna: ICC. Retrieved in 04 April 2016, from https://www. icc.or.at/standard_methods/162 18. Deepika, V., Jayaram Kumar, K., & Anima, P. (2013). Isolation and physicochemical characterization of sustained releasing starches from Dioscorea of Jharkhand. International Journal of Biological Macromolecules, 55, 193-200. PMid:23201777. http://dx.doi.org/10.1016/j.ijbiomac.2012.11.027.

19. Jiang, Q., Gao, W., Li, X., Xia, Y., Wang, H., Wu, S., Huang, L., Liu, C., & Xiao, P. (2012). Characterizations of starches isolated from five different Dioscorea L. Species. Food Hydrocolloids, 29(1), 35-41. http://dx.doi.org/10.1016/j.foodhyd.2012.01.011.

20. Pérez, E., Schultz, F. S., & de Delahaye, E. P. (2005).

Characterization of some properties of starches isolated from Xanthosoma sagittifolium (Tannia) and Colocassia esculenta (Taro). Carbohydrate Polymers, 60(2), 139-145. http://dx.doi. org/10.1016/j.carbpol.2004.11.033.

21. Zeng, F.-K., Liu, H., & Liu, G. (2014). Physicochemical properties of starch extracted from Colocasia esculenta (L.) Schott (Bun-long taro) grown in Hunan, China. Stärke, 66(12), 142-148. http://dx.doi.org/10.1002/star.201300039.

22. Sit, N., Misra, S., Baruah, D., Badwaik, L. S., & Deka, S. (2014). Physicochemical properties of Taro and maize starch and their effect on texture, colour and sensory quality of tomato ketchup. Stärke, 66(3-4), 294-302. http://dx.doi.org/10.1002/ star.201300120.

23. Deepika, V., Jayaram Kumar, K., & Anima, P. (2013). Isolation and partial characterization of delayed releasing starches of Colocasia species from Jharkhand, India. Carbohydrate Polymers, 96(1), 253-258. PMid:23688478. http://dx.doi. org/10.1016/j.carbpol.2013.04.002.

24. Tester, R. F., Karkalas, J., & Qi, X. (2004). Starch: composition, fine structure and architecture. Journal of Cereal Science, 39(2), 151-165. http://dx.doi.org/10.1016/j.jcs.2003.12.001.

25. Ellis, R. P., Cochrane, M. P., Dale, M. F. B., Duffus, C. M., Lynn, A., Morrison, I. M., Prentice, R. D. M., Swanston, J. S., & Tiller, S. A. (1998). Starch production and industrial use. Journal of the Science of Food and Agriculture, 77(3), 289-311. http:// dx.doi.org/10.1002/(SICI)1097-0010(199807)77:3<289::AIDJSFA38>3.0.CO;2-D.

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26. Shujun, W., Hongyan, L., Wenyuan, G., Haixia, C., Jiugao, Y., & Peigen, X. (2006). Characterization of new starches separated from different Chinese yam (Dioscorea opposita Thunb.) cultivars. Food Chemistry, 99(1), 30-37. http://dx.doi. org/10.1016/j.foodchem.2005.07.008.

27. Badenhuizen, N. P. (1969). The biogenesis of starch granules in higher plants. New York: Appleton-Century-Crofts.

28. Agama-Acevedo, E., Garcia-Suarez, F. J., Gutierrez-Meraz, F., Sanchez-Rivera, M. M., San Martin, E., & Bello-Pérez, L. A. (2011). Isolation and partial characterization of Mexican Taro (Colocasia esculenta L.) starch. Stärke, 63(3), 139-146. http://dx.doi.org/10.1002/star.201000113.

29. Imberty, A., Buléon, A., Tran, V., & Péerez, S. (1991). Recent advances in knowledge of starch structure. Stärke, 43(10), 375-384. http://dx.doi.org/10.1002/star.19910431002.

30. Gernat, C., Radosta, S., Damaschun, G., & Schierbaum, F. (1990). Supramolecular structure of legume starches revealed

by X-ray scattering. Stärke, 42(5), 175-178. http://dx.doi. org/10.1002/star.19900420504.

31. Shujun, W., Wenyuan, G., Hongyan, L., Haixia, C., Jiugao, Y., & Peigen, X. (2006). Studies on the physicochemical, morphological, thermal and crystalline properties of starches separated from different Dioscorea opposita cultivars.

Food Chemistry, 99(1), 38-44. http://dx.doi.org/10.1016/j. foodchem.2005.07.007.

32. Singh, N., Singh, J., Kaur, L., Sodhi, N. S., & Gill, B. S. (2003).

Morphological, thermal and rheological properties of starches from different botanical sources. Food Chemistry, 81(2), 219231. http://dx.doi.org/10.1016/S0308-8146(02)00416-8.

33. Chung, H.-J., Liu, Q., Lee, L., & Wei, D. (2011). Relationship between the structure, physicochemical properties and in vitro digestibility of rice starches with different amylose contents.

Food Hydrocolloids, 25(5), 968-975. http://dx.doi.org/10.1016/j. foodhyd.2010.09.011.

34. Saikia, J. P., & Konwar, B. K. (2012). Physicochemical properties of starch from aroids of North East India. International Journal of Food Properties, 15(6), 1247-1261. http://dx.doi.org/10.10 80/10942912.2010.491929.

35. Kizil, R., Irudayaraj, J., & Seetharaman, K. (2002).

Characterization of irradiated starches by using FT-Raman and FTIR spectroscopy. Journal of Agricultural and Food Chemistry, 50(14), 3912-3918. PMid:12083858. http://dx.doi. org/10.1021/jf011652p.

36. Zhao, J., & Whistler, R. L. (1994). Spherical aggregates of starch granules as flavor carriers. Food Technology, 48(7), 104-105. Retrieved in 04 April 2016, from http://agris.fao. org/agris-search/search.do?recordID=US9501867 37. Liporacci, J. S. N., Mali, S., & Grossmann, M. (2005). Efeito do método de extração na composição química e nas propriedades funcionais do amido de inhame (Dioscorea alata) [Effects of extraction method on chemical composition and functional properties of yam starch (Dioscorea alata)]. Semina: Ciências Agrárias, 26(3), 345-352. Retrieved in 04 April 2016, from http://www.uel.br/revistas/uel/index.php/semagrarias/article/ view/2308 Received: Apr. 04, 2016 Revised: Aug. 24, 2016 Accepted: Sept. 28, 2016

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http://dx.doi.org/10.1590/0104-1428.02016

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

Whey protein-based films incorporated with oregano essential oil Sandra Prestes Lessa Fernandes Oliveira1, Larissa Canhadas Bertan2, Christiane Maciel Vasconcellos Barros De Rensis1, Ana Paula Bilck1 and Priscila Cristina Bizam Vianna3* Programa de Pós-graduação em Ciência e Tecnologia de Leite e Derivados, Universidade Norte do Paraná – UNOPAR, Londrina, PR, Brazil 2 Universidade Federal da Fronteira Sul – UFFS, Laranjeiras do Sul, PR, Brazil 3 Departamento de Engenharia de Alimentos, Instituto de Ciências Tecnológicas e Exatas, Universidade Federal do Triângulo Mineiro – UFTM, Uberaba, MG, Brazil

1

*priscila.vianna@uftm.edu.br

Abstract This study aimed to prepare whey protein-based films incorporated with oregano essential oil at different concentrations, and evaluate their properties and antimicrobial activity. Films were more flexible with increasing the concentration of oregano oil and water vapor permeability was higher in the films with oregano oil. Increasing the concentration of essential oil decreased the water solubility. The solubility of control film and film with 1.5% oregano oil was 20.2 and 14.0%, respectively. The addition of 1% of oregano oil improved the resistance of the films. The tensile strength for the control film was 66.0 MPa, while for the film with 1% of oregano oil was 108.7 MPa. Films containing 1.5% oregano oil showed higher antimicrobial activity. The zone of inhibition ranged from 0 to 1.7 cm. The results showed that the whey protein-based films incorporated with oregano essential oil has potential application as active packaging.

Keywords: active film, whey protein, natural antimicrobial, packaging.

1. Introduction Edible coatings and films of various thickness consist of edible materials (proteins, lipids, carbohydrates) applied to the surface of food to improve the appearance and food preservation[1]. In recent years, there has been considerable interest in its use as food packaging, instead of the traditional plastic films. Edible coatings are defined as a thin layer of edible material applied directly to the surface of food, usually by immersion or aspersion, while edible films are pre-formed in molds, dried and applied to food[1-3]. The purpose of the use of films and coatings is to inhibit the migration of moisture, oxygen, carbon dioxide, flavors, and lipids, and carry food ingredients (antioxidants, antimicrobial agents, flavorings), and/or improve the mechanical and handling characteristics of the product. In some cases, the films and edible/biodegradable coatings can replace synthetic packaging[2].

Whey is a by-product of the cheese industry with high nutritional and functional value. When properly processed either as protein isolate or as concentrate, it consists of an excellent ingredient for the production of various processed foods. For a long time, whey had been considered as a waste by-product with no commercial value, which was discharged into watercourses causing serious environmental problems[4,5] or even incorporated into animal feed. This approach, however, was disregarded due to the excellent nutritional and functional properties of whey, and the production of whey protein-based films has been studied for application as alternative packaging[6].

The gelation phenomenon required for the formation of a film or coating, is achieved by physical and chemical

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interactions between the whey protein molecules. Despite the destabilization of whey protein can be achieved by various processes heating is the most used form. These processes induce the breakdown of the structure of the whey proteins, resulting in aggregation and gel formation. The heating modifies the three dimensional structure of proteins by exposure of the -SH groups and hydrophobic groups promoting intramolecular bonds and S-S hydrophobic interactions during drying of the films/coatings. In the absence of heat treatment, the films become brittle and fragile after drying[7,8].

Milk protein based films have good mechanical properties, and act as excellent barriers to lipid, oxygen and aromas.

However, they do not have an ideal water vapor barrier due to its hydrophilic nature[9-11]. When properly processed, whey proteins produced flexible, transparent, and odorless films[12]. The application of these films can be improved by adding antimicrobial agents to ensure a more effective conservation of the packaged or coated food[13-15].

Antimicrobial agents from natural or synthetic sources have been explored with satisfactory results regarding the growth inhibition of spoilage and pathogenic microorganisms.

Among the most studied antimicrobial agents are the organic acids and their salts (propionic, benzoic, sorbic), bacteriocins (nisin), enzymes (lysozyme, glucose oxidase), polysaccharides (chitosan), and natural plant extracts (essential oil from oregano, pepper, rosemary)[15-17]. Antimicrobial substances can be added to the polymers by melting or solubilization of the compound within the matrix. Once these substances

Polímeros, 27(2), 158-164, 2017

Whey protein-based films incorporated with oregano essential oil are heat sensitive, the solubilization method is the most suitable to incorporate them into the biopolymer matrix[13].

The oregano essential oil has emerged as the most studied essential oil with applications in food, and the Origanum vulgare sp has been one of the most studied for its antioxidant and antimicrobial activities [18-20].

This essential oil has high content of phenolic compounds responsible for its antimicrobial and antioxidant activity.

Among the components, both carvacrol and thymol stand out as active ingredients[21]. The hydroxyl groups in this phenolic compounds are the cause of inhibitory action as these groups can interact with the cell membrane of bacteria to disrupt membrane structures and cause the leakage of cellular components. There are suggestions that the active groups of carvacrol and thymol promotes the delocalization of electrons which then act as proton exchangers and reduce the gradient across the cytoplasmic membrane of bacterial cells. This will cause the collapse of the proton motive force and depletion of the ATP pool and ultimately leading to cell death. These phenolic groups can further bind the active site of enzymes by altering the cell metabolism of microorganisms[22]. Sahin et al.[23] reported the antimicrobial activity of oregano essential oil against 10 bacteria strains and 15 fungi and yeast species.

Oregano essential oil has been used in active films with different composition for use in foods. Pelissari et al.[24] observed inhibition of Bacillus cereus, Escherichia coli, Salmonella enteritidis and Staphylococcus aureus in starch‑chitosan based films incorporated with 0.1%, 0.5% and 1.0% of oregano essential oil. Furthermore, the oil addition has improved the water vapor permeability and led to formation of more flexible films. Oussalah et al.[25] produced milk protein based films incorporated with oregano and pepper oil for beef preservation. The films incorporated with oregano oil stabilized lipid oxidation in beef during 7 days of storage and proved effective against E. coli O157: H7 and Pseudomonas spp strains.

Seydim and Sarikus[10] evaluated the antimicrobial properties of whey protein-based films incorporated with different concentrations of essential oil extracted from oregano, rosemary and garlic. Films containing 3 and 4% oregano and garlic essential oils were effective against S. aureus, S. enteretidis, L. monocytogenes, E. coli and L. plantarum.

In contrast, films incorporated with rosemary oil did not exhibit antimicrobial activity. The authors highlighted the potential application of the active films in packaging of cheese slices.

Thus, the whey protein-based films incorporated with essential oils with antimicrobial properties can be used as an active packaging technology applied to dairy products such as processed and sliced cheese to inhibit the growth of pathogenic and spoilage microorganisms, and simultaneously increase the shelf life of these products. The aim of this study was to produce and evaluate the effects of the incorporation of oregano essential oil on the water vapor permeability, solubility, tensile strength and antimicrobial activity of whey protein-based films.

Polímeros, 27(2), 158-164, 2017

2. Materials and Methods 2.1 Materials The films were produced with whey protein isolate (WPI 95% Bipro, Davisco, USA). Glycerol (Synth, Brazil) was added to the formulation as plasticizer. Oregano essential oil (Origanum vulgare) was acquired from Ferquima Indústria e Comércio Ltda (São Paulo, Brazil). The essential oil met all the required specifications in relation to its quality control as the certificate of analysis provided by the supplier.

2.2 Films manufacture Four types of films were made: control film, composed of whey protein isolate and glycerol; and whey protein‑based films incorporated with oregano essential oil at three different concentrations: 0.5%, 1.0% and 1.5% (v/v). The concentrations of essential oil were defined according to literature[10].

The films were produced by casting according to the methodology described by Yoshida and Antunes[26]. The whey protein isolate (7.0%, dry weight basis) was dispersed in distilled water, followed by agitation on a magnetic stirrer until complete solubilization. Then, 3.0% glycerol (w/w) was added and the film forming solution was heated to 90 °C/30 minutes in a water bath to denature the whey proteins.

After cooling to room temperature (25 °C), the pH of the solution was adjusted to 7.0 using NaOH 0.1 M. An aliquot of 10 mL was poured into 9 cm Petri dishes slightly greased with sunflower oil, and left to dry at 24 °C/24 hours, on the laboratory bench. After drying, the films were removed from the dishes with a spatula, stored in desiccators with 53% relative humidity (magnesium nitrate saturated solution), and separated by sheets of tissue paper.

The films containing the antimicrobial agent were produced following the same procedure used for the control film, differing only by the addition of oregano essential oil (OEO) at concentrations of 0.5%, 1.0% and 1.5% (v/v) to their respective film forming solution prior to heat treatment at 90 °C/30 minutes, to facilitate their incorporation into the protein matrix. The drying and conditioning of the films was the same of the control film.

2.3 Films characterization Before the characterization analyses, all dried films were stored in desiccators at 25 °C and 53% relative humidity (RH) for 48 hours. The films were characterized for visual appearance, thickness, water vapor barrier properties, water solubility and tensile strength.

The films were evaluated for visual and tactile aspect aimed to select only the homogeneous films (without the presence of insoluble particles, bubbles and uniform color), with no breaks or frangible areas to the analytical determinations.

The thickness of the films was determined using a digital micrometer with a resolution of 0.001 mm (Mitutoyo MDC-25M, Kanagawa, Japan) and was calculated as the mean of ten measurements of the film area.

The water vapor permeability (WVP) was determined gravimetrically at 25 °C according to standard method E-96-95 ASTM[27]. The best films previously selected by visual 159

Oliveira, S. P. L. F., Bertan, L. C., De Rensis, C. M. V. B., Bilck, A. P., & Vianna, P. C. B. and tactile aspect were cut into discs with 9 cm in diameter and set out in the circular opening of the capsule to ensure moisture migration exclusively through the film. The interior of the capsule was filled with 15 g calcium chloride, and the system was placed into a desiccator containing a saturated solution of sodium chloride, corresponding to RH 75%.

The test was carried out in triplicate for each formulation with RH gradient (0-75%) at 25 °C. Water vapor transport was determined from the weight gain of the calcium chloride measured every 24 hours for 7 days. The relative water vapor permeability (RWVP) was calculated as follows:

(m / t )

RWVP =

× (1/ A )

(1)

where A is the permeation area of the film (m2), m is the mass of vapor permeated (g) and t is the time (days).

The water vapor permeability (WVP) was calculated by the equation: = WVP

( RWVP

× e ) / ( ps x ( RH1 – RH 2 ) )

(2)

where e is the average film thickness (m), pS is the saturated vapor pressure at the test temperature (Pa), RH1 is the relative humidity inside the desiccator and RH2 is the relative humidity inside the capsule.

The water solubility of the films was determined according to the method proposed by Gontard et al.[28], in triplicate.

The samples were cut into discs of 2.5 cm in diameter and dried in an oven at 105 °C for 24 hours to obtain the initial dry mass. The dried samples were immersed in a beaker containing 50 mL distilled water and kept under slow stirring (100 rpm) for 24 hours at 25 °C using a magnetic stirrer.

After this period, the samples were subjected again to drying at 105 °C for 24 hours to obtain final dry mass. The water solubility was calculated using the following equation:

(

SOL =− Mi M f

)

/ M i × 100

(3)

where: SOL is the solubilized mass as a function of the initial dry mass (%), MI is the initial dry mass (g) and MF is the final dry mass after solubilization (g).

The tensile strength (TS) was measured using a TA.XT2 Texture Analyzer (Stable Micro System, Hasleme, England) according to the method of the American Society for Testing and Materials[29]. The measurements were carried out at 25 °C. The samples were cut into 80 mm long × 25 mm wide strips, and was adjusted to the grip device with a distance of 30 mm. The test speed was 1 mm/s. Ten replicates of each formulation were performed. The ultimate tensile strength was calculated according to the equation: TS = Fm / A

(4)

where TS is the ultimate tensile strength at break (MPa), FM is the maximum force at break of the film (N), and A is the cross sectional area of the film (m2).

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2.4 Antimicrobial effectiveness of the films by agar diffusion method To determine the antimicrobial effectiveness of the whey protein-based films, a lyophilized culture of Penicillium commune (CCT 4685) donated by André Tosello Foundation (Campinas, São Paulo, Brazil) was used. The reactivation of lyophilized cells of Penicillium commune was performed according to the following procedure: the rehydration was carried out for 30 minutes in sterile distilled water, and the entire contents of the vial were poured onto a plate containing Potato Dextrose Agar (PDA, Difco Laboratories), which was incubated at 25 °C/5 days. After growth, the colonies were transferred to tubes containing PDA agar, which were incubated at 25 °C/5 days and then kept under refrigeration (7 °C). These tubes contained the inoculum to be used in the diffusion tests.

For the diffusion test, the colonies were collected, with a handle, from the tubes containing the inoculum and transferred to an erlenmeyer having 50 mL 0.1% peptone water with 0.1% Tween 80. The suspension was stirred for 1 min and filtered with the aid of sterile gauze. From this initial inoculum, decimal dilutions were performed, and the suspension for the diffusion test was selected according to the growth characteristics of the target microorganism.

The agar diffusion test was used to qualitatively evaluate the antimicrobial activity of the films according to Pelissari et al.[24] with modifications. The active films were placed on the surface of the solidified PDA, together with 0.1 mL suspension of Penicillium commune seeded with Drigalski handle. Circular cuts (2.5 cm in diameter) of the control films and films incorporated with 0.5%, 1.0% and 1.5% oregano essential oil were exposed to UV light for 15 min and placed aseptically onto the surface of solidified medium inoculated with the target microorganism. The plates were incubated at 25 °C/5 days. The efficiency of the antimicrobial agent was evaluated by the formation of halo around the disks and by quantification of colony-forming units (CFU mL-1). All experiments were performed in triplicate.

2.5 Statistical analysis The effect of the addition of oregano essential oil on the characteristics and antimicrobial activity of edible films was assessed by analysis of variance and Tukey’s test at 5% significance level. Statistical analysis was performed using the Statistica software (version 8.0, 2007; Stat-Soft Inc., Tulsa, OK).

3. Results and Discussions 3.1 Films characterization The whey protein-based films with and without addition of oregano essential oil (OEO) at different concentrations were, in general, homogeneous, transparent, and yellowish.

The films incorporated with higher concentrations of essential oil (1.0% and 1.5%) were visually more malleable than the others were. Ramos et al.[8] studied films produced from whey protein isolate or protein concentrate, and observed similar visual characteristics to those found in this study.

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Whey protein-based films incorporated with oregano essential oil Table 1 shows results for thickness, water vapor permeability, water solubility and tensile strength of active whey protein films.

Film thickness ranged from 0.013 mm to 0.015 mm (Table 1) and the control film showed lower thickness than the films with 0.5% and 1.0% essential oil. The difference between the thicknesses of the control as compared with the films incorporated with 0.5% and 1.0% essential oil could have been caused by lack of accuracy in measuring the volumes of filmogenic solution added to the Petri dishes.

Bertan et al.[30] observed that the addition of hydrophobic substances promoted an increase of the biofilm thickness, being necessary to use different ratios for each formulation aimed at controlling thickness for the repeatability of the measurements and validity of comparisons between properties. In the case of this study, it can be considered that the percentage of the hydrophobic substance (oregano essential oil) was too low to cause such variation.

The water vapor permeability (WVP) was significantly higher in the films containing the oregano essential oil (Table 1), contrary to what is expected for films made with hydrophobic substances. The incorporation of lipids in the protein film matrix reduces the interaction for water molecules, reducing the water vapor permeability of the films[26]. Probably the oregano essential oil could not bind chemically to the protein matrix, leaving empty spaces through which water could permeate, causing a higher WVP in active films.

Gontard et al.[28], who evaluated the incorporation of hydrophobic substances to wheat gluten based films, found similar results. The authors observed that the hydrophobic molecules having substantially spherical large size, when used in formulations of composite films, can lead to the occurrence of cracks and preferred channels through which water can diffuse, resulting in higher water vapor permeability.

Despite the low concentrations of oregano essential oil used in this study, the increase of WVP could have been caused by the incompatibility of oil with the whey protein isolate, resulting in a disruption of the film matrix.

In general, the films were very poorly permeable to water vapor, regardless of the addition of oregano oil. The values found in the present study were lower than those observed by other authors. Yoshida and Antunes[26] found mean WVP values of 0.3023 and 0.2680 g mm h-1 m-2 kPa-1 for control whey protein-based films and whey protein-based films emulsified with stearic acid, respectively. WVP values of 15.013 and 13.526 g mm h-1 m-2 kPa-1 were found by Mei and Zhao[31] for control whey protein-based films and films incorporated with 0.2% α-tocopherol acetate.

The control film and those containing 1% essential oil had significantly higher solubility than the film containing

1.5% oil. No differences among the other films were observed (Table 1). The lower solubility of the film incorporated with 1.5% essential oil can be explained by its higher hydrophobic character.

All films remained intact after the solubility test, indicating the formation of a stable network. McHugh and Krochta[32], Fairley et al.[33] and Galietta et al.[34] stated that whey protein-based films are partly water insoluble due to the presence of intermolecular disulfide bonds, making them resistant to solubilization in aqueous buffers. The presence of this type of binding in the matrix pores may be trapped the oregano oil, thus decreasing water solubility, which was intensified by both the higher oil concentration and the hydrophobic nature of the compound. The water solubility of the protein appears to be related to its hydrophilicity and the structure of the resulting protein matrix. The incorporation of a hydrophobic compound in the formulation of protein films reduces the ability of the film-forming matrix for bonding with water molecules[32]. According to Bertan[35], the water solubility of the films is directly related to its components, i.e., the hydrophilicity/hydrophobicity, and structure.

The control film showed less resistance when compared to the others (Table 1). Oregano oil has improved the resistance of the films. Among the different oil concentrations, the film incorporated with 1% essential oil had significantly higher resistance than the other films. The heat used for the preparation of the films caused denaturation of whey proteins and exposure of sulfhydryl groups, resulting in an increase of S-S covalent bonds in the films, featuring the formation of stable films capable to extend[7,33].

Yoshida[6] found that the tensile strength was higher for films containing high concentrations of whey protein (5.5% to 7.5%), resulting in an increase in tensile strength of about 0.4493 MPa, which is related to the increased number of S-S covalent bonds in the film matrix, due to the larger amount of sulfhydryl groups on the protein surface.

The incorporation of oregano oil in the film forming solution caused the destabilization of the matrix, resulting in increased WVP, and the bonds formed due to protein denaturation have trapped the oil, thus increasing hydrophobicity of the matrix.

The “spacing” caused by the oil may have promoted most interactions between the side groups of the protein chain, resulting in greater interactions and consequently greater tensile strength. The increase in concentration from 1% to 1.5% oil resulted in a decrease in the tensile strength from 108.7 MPa to 91.5 MPa, probably due to the plasticizing effect caused by the “excess” of oil or the lubricating effect of the matrix. Pelissari et al.[24] also observed a reduction in TS in starch-chitosan based films incorporated with 0.1%, 0.5% and 1.0% oregano essential oil, and attributed the results to the plasticizing capacity of oregano oil.

Table 1. Mean values ​​of thickness, water vapor permeability, water solubility and tensile strength (TS) of the whey protein-based films.

Treatments Control 0.5% OEO1 1.0% OEO 1.5% OEO

Thickness

WVP

Solubility

TS

(mm) 0.01342 ± 0.002b 0.01542 ± 0.001a 0.01500 ± 0.001a 0.01475 ± 0.001ab

(g mm m-2 day-1 kPa-1) 0.00070 ± 0.00007a 0.0010 ± 0.0004b 0.0011 ± 0.0001b 0.00100 ± 0.00007b

(%) 20 ± 2a 17 ± 2ab 17 ± 1a 14 ± 3b

(MPa) 66 ± 5c 96 ± 3b 109 ± 7a 92 ± 8b

OEO: oregano essential oil. Averages with different letters in the same column differ significantly (P < 0.05).

1

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Oliveira, S. P. L. F., Bertan, L. C., De Rensis, C. M. V. B., Bilck, A. P., & Vianna, P. C. B. Table 2. Microbiological efficiency of the active films by diffusion test against Penicillium commune.

Treatments Control 0.5% OEO1 1.0% OEO 1.5% OEO

CFU mL-1 5.13 x 105a 3.17 x 105a 3.14 x 105a 1.98 x 105a

Diameter of zone of inhibition (cm) 0.00 ± 0b 0.07 ± 0.10b 1.3 ± 0.8a 1.7 ± 0.5a

OEO: oregano essential oil. Averages with different letters in the same column differ significantly (P < 0.05).

1

Figure 1. Diffusion test for whey protein-based films incorporated with different concentrations of oregano essential oil (OEO) against Penicillium commune.

3.2 Diffusion test The results of the diffusion test are presented in Table 2.

The average concentration of the Penicillium commune suspension was 5.5 x 105 CFU mL-1. The addition of oregano essential oil caused no significant difference in the number of colony-forming units, which was in the order of 105 CFU mL-1 for all treatments.

The zone of inhibition was measured as the shortest distance between the film and the nearest colony-forming unit. The halo diameter varied from 0 to 1.7 cm. A significant increase in the halo diameter with increasing concentration of oregano (Figure 1) oil was observed. Films containing 1.0% and 1.5% oregano oil showed significantly higher halos than both the control film and that containing 0.5% oil. No significant difference was observed between the halos of the films with 1.0 and 1.5% oregano oil.

Oussalah et al.[25] showed that 1% oregano essential oil incorporated into a calcium caseinate and whey protein isolate edible films containing carboxymethyl cellulose showed inhibitory effect against E. coli O157: H7 and Pseudomonas spp. on the surface of beefsteaks. This study suggests that 2% oregano essential oil in WPI based films was the minimum concentration with inhibitory effect against S. aureus, S. enteritidis, L. monocytogenes, L. plantarum e E. coli O157: H7.

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Seydim and Sarikus[10] performed the halo test in whey protein isolate edible films containing oregano essential oil against different microorganisms. Unlike the results observed in this study, the authors found no antimicrobial effect in the films with 1% oregano oil for the target microorganisms.

The minimum concentration required for inhibition was 2%.

The largest zone of inhibition was observed at concentration of 4% oregano oil against S. aureus, Salmonella enteritidis and Listeria monocytogenes.

Pelissari et al.[24] studied starch-chitosan based films incorporated with oregano essential oil at different concentrations (0%, 0.1%, 0.5%, and 1.0%), and found an increase of the zones of inhibition with increasing oil concentrations for the different microorganisms. The largest zones of inhibition were observed for Gram-positive bacteria (B. cereus and S. aureus) when compared with Gram-negative bacteria (S. enteritidis e E. coli).

4. Conclusions Our results indicate that whey protein isolate is a good matrix for the production of biodegradable flexible films, and the addition of oregano essential oil, even in small quantities positively affected the characteristics of the films. Low vapor permeability values and low water Polímeros, 27(2), 158-164, 2017

Whey protein-based films incorporated with oregano essential oil solubility of films with oregano essential oil makes it can be used in both dry foods and in products with a greater water activity. The addition of oregano oil also allowed the production of more flexible films and more resistant at break, particularly at a concentration of 1.0%. In addition, the active films exhibited antimicrobial activity against Penicillium commune at concentrations of 1.0% and 1.5%.

Based on these results it can be concluded that the films added with 1.0% of oregano essential oil presented the best characteristics. Thus, the whey protein isolate-based films incorporated with oregano essential oil has potential application as active food packaging, which may extend the shelf life of food products. Due to their characteristic aroma, it is recommended to use these active films as an alternative to commercial coatings for cheeses.

5. Acknowledgements The authors are grateful to the Fundação Nacional de Desenvolvimento do Ensino Superior Particular (FUNADESP, Brasília, DF, Brazil) for the financial support and Fundação Tropical de Pesquisas e Tecnologia André Tosello for donating the strain of Penicillium commune.

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21. Sivropoulou, A., Papanikolaou, E., Nikolaou, C., Kokkini, S., Lanaras, T., & Arsenakis, M. (1996). Antimicrobial and cytotoxic activities of origanum essential oils. Journal of Agricultural and Food Chemistry, 44(5), 1202-1205. http:// dx.doi.org/10.1021/jf950540t.

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23. Sahin, F., Güllüce, M., Daferera, D., Sökmen, A., Sökmen, M., Polissiou, M., Agar, G., & Özer, H. (2004). Biological activities of the essential oil and methanol extract of Origanum vulgare ssp. vulgare in the Eastern Anatolia region of Turkey.

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25. Oussalah, M., Caillet, S., Salmiéri, S., Saucier, L., & Lacroix, M. (2004). Antimicrobial and antioxidant effects of milk protein based film containing essential oils for the preservation of whole beef muscle. Journal of Agricultural and Food Chemistry, 52(18), 5598-5605. PMid:15373399. http://dx.doi. org/10.1021/jf049389q.

26. Yoshida, C. M. P., & Antunes, A. J. (2009). Aplicação de filmes proteicos à base de soro de leite. Ciência e Tecnologia de Alimentos, 29(2), 420-430. http://dx.doi.org/10.1590/ S0101-20612009000200030.

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Edible composite films of wheat and lipids: water vapour permeability and other physical properties. International Journal of Food Science & Technology, 29(1), 39-50. http:// dx.doi.org/10.1111/j.1365-2621.1994.tb02045.x.

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31. Mei, Y., & Zhao, Y. (2003). Barrier and mechanical properties of milk protein-based edible films containing nutraceuticals.

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33. Fairley, P., Monahan, F. J., German, B. B., & Krochta, J. M. (1996). Mechanical properties and water vapor permeability of edible films from whey proteins isolate and N-ethylamaleimidie or cysteine. Journal of Agricultural and Food Chemistry, 44(12), 3789-3792. http://dx.doi.org/10.1021/jf9601731.

34. Galietta, G., Di Gioia, L., Guilbert, S., & Cuq, B. (1998).

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35. Bertan, L. C. (2003). Desenvolvimento e caracterização de filmes simples e compostos a base de gelatina, ácidos graxos e breu branco (Master’s dissertation). Universidade Estadual de Campinas, Campinas.

Received: Mar. 07, 2016 Revised: July 26, 2016 Accepted: Dec. 13, 2016

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http://dx.doi.org/10.1590/0104-1428.05916

Gamma radiation effect on sisal / polyurethane composites without coupling agents Marina Cardoso Vasco1*, Salvador Claro Neto2, Eduardo Mauro Nascimento3 and Elaine Azevedo3 Department of Mechanical Engineering and Aerospace, University of Patras, Patras, Greece 2 Instituto Qumica de São Carlos, Universidade de São Paulo – USP, São Carlos, SP, Brazil 3 Programa de Pós-graduação em Engenharia Mecânica e de Materiais, Universidade Tecnologica Federal do Paraná – UTFPR, Curitiba, PR, Brazil 1

*marina.mcv@gmail.com

Abstract Natural fibers and polyurethane based composites may present chemical bonding between the components of the polymer and the lignin of the fiber. The incidence of radiation can cause degradation of the polymeric material and alter its mechanical properties. The objective of this study was to obtain and characterize cold pressed composites from polyurethane derived from castor oil and sisal fibers, without coupling agents, through thermogravimetric and mechanical tests, before and after the incidence of 25 kGy dose of gamma radiation. Woven composites that were not irradiated had maximum values of 4.40 GPa for flexural elastic modulus on three point flexural test and dispersed fiber composite that were not irradiated had maximum values of 2.25 GPa. These materials are adequate for use in non‑structural applications in radiotherapy and radiodiagnostic rooms.

Keywords: gamma radiation, sisal, polyurethane, mechanical properties, green composite.

1. Introduction The use of natural fibers as load or reinforcement in polymeric composites is increasing in the last years.

Characteristics such as processing flexibility, high specific stiffness and low cost make them attractive to industrial applications[1-5]. Polymeric matrices are preferred because of their low cost and ease of processing, and their main functions are joining the fibers, distributing stresses and stopping crack propagation in the material[6]. Conventional polymers used in radiotherapy furniture and radio diagnostic rooms, such as adhesives and agglomerates, have their mechanical properties considerably altered when exposed to high energy radiation[7].

The effects of gamma radiation in polyurethane derived from castor oil adhesive specimens were studied

by Azevedo et al.[8], and they concluded that there were no significant alterations of the mechanical properties when materials were irradiated with doses up to 100 kGy, since the main effect of radiation in this polymer is the formation of cross-linked bonds within polymeric chains.

The lack of adhesion between the fiber and the matrix is one of the greatest concerns regarding the production of bio composites. Mechanical loads are held by the fibers, and the matrix/fiber interface should transfer the applied load from the matrix to the fibers in order to achieve the desired mechanical strength[3,9]. In general, natural fibers are hydrophilic and do not show good adhesion with hydrophobic polymers normally used as composite matrices[9].

Many different approaches described elsewhere[9-16] have been explored in order to improve the fiber/matrix adhesion.

They include the chemical modification of the fiber before the composites processing, either through esterification[17-22], eterification[17,21,22], silane[11,12,21] or isocyanate[9,23] treatments,

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or through physical means as plasma[11,13,24,25] or corona[24,25] treatments, as well as the modification of the polymeric matrix[26,27].

During the process of polymerization of the polyurethane/sisal composite utilized in this work, the active hydrogen present in the sisal lignin, either in its hydroxyl or carboxyl groups, may chemically interact with free isocyanate from polyurethane and form urethane bonds or carbamic anhydride[28]. Chemical interaction between isocyanate and lignin may reinforce the polymer-fiber interface, leading to a better load distribution throughout the composite structure when it’s subjected to mechanical stresses. This may enhance its mechanical performance with no need of adding coupling agents or conducting pre-treatments on the fibers.

The aging resistance of sisal fibers is good[7], and they are able to decrease the effect of gamma radiation on the mechanical properties of the polyurethane matrix, as well as to increase the fiber/matrix cohesion. Huang et al.[29] observed the influence of gamma, UVA, and UVC radiations on the flexural strength of composites with different volumetric proportions of fiber and resin, and concluded that the incidence of radiation degraded the mechanical properties of the studied materials, without affecting the necessary characteristics for their use as accessories in radiotherapy and radio diagnostic rooms. However, the possible changes in chemical affinity between the components and the mass loss with temperature as a function of the incident radiation were not studied.

The aim of this work is to evaluate the mechanical properties of composites made with castor oil derived polyurethane and dispersed and woven sisal fibers, with no compatibility treatments, before and after incidence of

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

Vasco, M. C., Claro, S., No., Nascimento, E. M., & Azevedo, E. a 25kGy dose of gamma radiation. This dose is the same used sterilization of medical materials. Specimens were characterized and evaluated by means of thermogravimetric analysis (TGA), cryogenic fracture tests, nanoindentation and nanoscratch tests, Izod impact test and flexural test, and fracture surfaces were analyzed by scanning electron microscopy (SEM).

2. Materials and Methods The polyurethane used in this work was donated by the company Cequil, from Araraquara - SP, Brazil. It was supplied as a bi-component, a poliol and a pre-polymer, identified as 442 and 253, respectively. The preparation of the polymer was made according to the supplier instructions, and the ratio poliol:pre-polymer was 1:1 for all specimens.

The setting time of the polymer was 24h.

Sisal fibers were acquired from local stores in Curitiba-PR, Brazil, and there was no concern regarding its precedence.

Two different geometries were used: (1) dispersed fibers of approximately 5mm in length, and (2) bidirectional woven with thickness of about 2mm composed of sisal threads with approximate diameter of 1mm. Fibers was previously kept in an oven at 70 °C for six hours in order to eliminate moisture, and this was the only processing that they subjected to before mixing with the polymer. Two different sources of fibers were used, since it was not possible to make dispersed fibers form the woven, because it resulted in agglomerates, making the homogenization of the composite very difficult.

Fibers and polymer were manually mixed with a fiber:polymer ratio of 80/20 wt%, and rectangular sheets 3-4mm thick were produced. The mixture was pressed at room temperature in a hydraulic press, and the applied pressure was 266.6 Pa. No vacuum was made during pressing. Specimens were then cut from pressed sheets for characterizations, and the dimensions of the specimens used for destructive tests are in accordance with ASTM 790-03 and ASTM D 256 standards (flexure tests and impact tests, respectively).

Gamma irradiation tests, with dose of 25 kGy, were performed by the company Embrarad - Cotia/SP with the aid of industrial source Cobalt 60 MDS Nordion JS-9600.

D256 standard. All woven composites specimens were cut in a parallel direction to the length of the fibers.

Scanning Electron Microscopy was performed in a Zeiss microscope, model EVO MA 15, and it was used to evaluate the surface of fibers after drying and the fracture surface of the composites after flexural test.

3. Results and Discussion Table 1 shows the flexural properties of polyurethane derived from castor oil before and after the gamma irradiation with dose of 25 kGy. A slight increase on flexural stress and flexural elastic modulus after irradiation is noticed, which, according to Azevedo et al.[8,30], can be attributed to the formation of crosslink bonds in the polymeric chains caused by the gamma radiation incidence.

Figure 1 shows the obtained TGA curves of the dispersed fibers composite without irradiation. A thermal event starting at 50 °C can be observed, and this can be attributed to the presence of water[31] and other substances that are present onto the surface of the fibers. This thermal event finishes at 150 °C, and the next one occurs at 520 °C. It can be assumed that the thermal behavior of the composite is combination of the thermal behavior of its individual components, and the phase with higher volumetric fraction gives the major contribution. The 20% of the mass remaining at the end of the test can be identified as burning residues.

The TGA curves of the dispersed fibers composite after gamma irradiation with dose of 25kGy are shown in Figure 2.

Comparing to the non-irradiated material, it can be noticed that the temperatures of the thermal events did not change, but the peaks intensities decreased in the derivative curve, which is a sign of degradation of the fibers caused by the Table 1. Stress rupture values in flexion and elastic modulus in bending the polyurethane derived from castor oil before and after incidence of gamma radiation (25 kGy dose).

PU Before irradiation After irradiation

Flexural stress (MPa) 42 ± 3 47 ± 12

Elastic modulus (GPa) 1.64 ± 0.12 1.66 ± 0.50

The thermogravimetric analyses (TGA) were carried out in a Perkin Elmer STA 6000 thermo scale, and the following parameters were used: • Specimen mass: 4.0 to 5.0 mg; • Temperature range: 50 °C to 800 °C; • Heating rate: 10 °C/min; • Flow of N2 atmosphere: 20 mL/min.

Flexural tests were performed according to the ASTM D790-03 standard in an EMIC DL10000 universal tester, with load cell of 5 kN and test speed of 1mm/min. To determine impact resistance, test specimens of 55 × 10 × 4 mm (length × width × thickness) were submitted to Izod impact test in Ceast equipment, model Resil 25, notched in Ceast chisel with depth of 2.54 ± 0.1 mm, according to the ASTM 166

Figure 1. Thermogravimetric Analysis curves obtained for the non-irradiated dispersed sisal fibers composite.

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Gamma radiation effect on sisal / polyurethane composites without coupling agents incident radiation. The amount of burning residues after the test showed no significant variation.

The curves obtained from thermogravimetric analysis of the non-irradiated woven composite are shown in Figure 3.

A thermal event starts at 50 °C and finishes at 150 °C, and this event is also attributed to the presence of water, similar to the previous cases. Following that, another thermal event starts at 150 °C and finishes at 520 °C, and the same amount (20%) of burning residues remains at the end of the test.

Figure 4 shows the TGA curves of the woven composite after gamma irradiation. Again, when comparing to the non-irradiates specimens, no variation of the initial and final temperatures of the thermal events was observed.

The decreasing intensity of the peaks in the derivative curve indicates the influnece of the gamma radiation on the degradation of the sisal woven, and the burning residues mass was the same as observed previously.

It was observed that the intensity of the thermal events for the sisal woven composites occurring in the temperature range between 250 °C and 450 °C is larger than its correspondents for dispersed fibers composites. That could be explained by the difference in morphology of the reinforcements, implying in different modes of thermal energy transference inside the materials.

Figure 5 shows a SEM micrograph of the fracture surface after cryogenic fracture of dispersed fibers composite without incidence of gamma radiation.

Figure 2. Thermogravimetric Analysis curves obtained for the gamma irradiated dispersed sisal fibers composite.

Figure 4. Thermogravimetric Analysis curves obtained for gamma irradiated woven sisal fibers composite.

Figure 3. Thermogravimetric Analysis curves obtained for the non-irradiated woven sisal fibers composite.

Figure 5. SEM micrograph of cryogenic fracture of non-irradiated dispersed sisal fibers composite.

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From Figure 5 one can see that there are no voids or gaps on the fiber-matrix interface, suggesting that, besides the mechanical interlocking, fibers may present chemical affinity with this type of polyurethane, increasing surface adhesion and providing better mechanical properties, since load transfer becomes more efficient. More than one failure mode was observed in the fracture surfaces of the composites, which is expected for this type of material. Fiber breakage and tearing, as well as polymer matrix fracture were observed.

The polymer fracture was of brittle in nature, evidenced by mirrored areas starting at stress concentration points such as bubbles, which are inherent to the manufacturing process of the polymer, impurities at the fibers surface and fibers edges. Figure 6 shows a SEM micrograph of the dispersed fibers composite after gamma irradiation. It is observed that there are no gaps around the fiber-polymer interface, suggesting that even after the incidence of the applied dose of gamma radiation the chemical affinity between the fibers and the polyurethane matrix is maintained. It is also observed that the amount of empty regions in the fibrils interior has enhanced, which can indicate embrittlement due to radiation,

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Vasco, M. C., Claro, S., No., Nascimento, E. M., & Azevedo, E. behavior also evidenced by detachment of micro fibrils in other regions of the material.

Figures 7a and 7b show the surface fractures after cryogenic fracture of the non-irradiated and irradiated sisal woven composites, respectively. In both cases, the fracture aspect was the same observed for the dispersed fiber composite, with embrittlement of micro fibrils after irradiation and no presence of interfacial gaps between the fibers and the polymer matrix.

Figure 6. SEM micrograph of cryogenic fracture of gamma irradiated dispersed sisal fibers composite, dose of 25 kGy.

The absence of fiber-matrix interfacial gaps was reported by other authors[32,33] only for fibers that were treated with coupling agents, witchy suggests that the chemical affinity between the materials investigated in this work is suitable even without the use of coupling agents.

The results obtained for flexural stress and elastic modules before and after the incidence of gamma radiation are shown in Figures 8a and 8b. Specimens were identified as follows: DF - dispersed fiber composite; W - woven fiber composite; NI - non-irradiated.

Both flexural stress and elastic modules of dispersed fibers composites were lower when compared to woven composites and non-reinforced polyurethane. Such result could be attributed to the fact that the fibers extremities act as stress risers and crack initiation sites, decreasing the material flexural strength[6]. In this case, the higher fiber-to-matrix volume ratio contributed to this decrease in flexural strength, since there are more stress risers points[34,35].

It was also observed that the incidence of gamma radiation decreased the mechanical strength of the materials, because of the degradation on fibers caused by radiation[7], which was evidenced by the surface fractures observation. The highest flexural strength was obtained by the non-irradiated sisal woven composite, with an increase of 190% in flexural strength and 268% in flexural elastic modules when compared to the non-irradiated and non‑reinforced polyurethane.

Figure 7. SEM micrograph of fracture surfaces after cryogenic fracture of woven sisal fibers composite: (a) before gamma radiation, and (b) after gamma irradiation with dose of 25 kGy.

Figure 8. Average values and standard deviations for (a) flexion stress, and (b) flexion elastic modulus for each composite, before and after gamma radiation incidence with dose of 25 kGy. 168

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Gamma radiation effect on sisal / polyurethane composites without coupling agents Table 2. Impact resistance values of the dispersed fibers and woven composites before and after incidence of gamma radiation (dose of 25 kGy).

Composite Dispersed fibers without radiation Dispersed fibers 25 kGy Woven without radiation Woven 25 kGy

Impact resistance (kJ/m2) 10.4 ± 3.4 13.1 ± 3.2 8.7 ± 1.3 9.1 ± 1.3

The comparison between impact resistance of the composites before and after irradiation is shown in Table 2.

Higher values were measured for disperse fibers composite, probably because the fibers extremities acts as crack arresters, impeding the crack propagation and consequently increasing impact resistance. It was also noticed that standard deviation values for dispersed fibers composites are higher than for woven composites, due to lack of uniformity of fibers dispersion among different specimens. The manufacturing process of the specimens was made manually, which may have originated heterogeneous regions with different fibers concentration within the polymer matrix.

The impact resistance of both composites increased after the incidence of gamma radiation, as a result of the higher density of cross-linked bonds in the polyurethane matrix.

Silva et al.[1] reported impact resistance values of 23kJ/m2 for short untreated sisal fibers with fiber content of 20% in volume and 25kJ/m2 with 27% with fiber content. Alkali treatment performed on the fibers in[1] resulted in decreased impact resistance of the composites - 11kJ/m2 with 20% fiber volume and 10kJ/m2 with 27% of fibers. Since the fibers used in this work were commercially available ones, with no information regarding any possible pretreatment, it is possible that they were previously processed in a similar manner as reported by Silva et al.[1].

4. Conclusions The objective of this study was to obtain and characterize cold-pressed composites made of polyurethane derived from castor oil and sisal fibers, without the use of coupling agents, through termogravimetric analysis, surface observation and mechanical tests, before and after the incidence of gamma irradiation, with dose of 25 kGy. Through thermogravimetric analysis it is possible to conclude that all studied composites showed thermal stability up to 150 °C, even after incidence of gamma radiation, and that none of them had solvents in their composition. SEM cryogenic fractures observation indicated that there was no gap between sisal fibers and polyurethane matrix in any of the materials, even after gamma irradiation. Flexural tests results showed that dispersed fibers composites had lower flexural strength and flexural elastic modulus values. This happened because dispersed fibers extremities act as stress concentration points, diminishing flexural mechanical strength. Non-irradiated woven composites had the highest results in flexural tests, since its interlaced structure provides better load distribution. After radiation incidence, both composites had their mechanical properties decreased. Dispersed fibers composites had better impact resistance when compared to woven composites, which is also explained by the highest amount of inner stresses Polímeros, 27(2), 165-170, 2017

caused by fibers extremities. Gamma radiation incidence has increased the impact resistance in both materials. Based on these observations it is possible to conclude that these composites are suitable for applications in radio diagnosis and radiotherapy rooms.

5. Acknowledgements The authors gratefully acknowledge CNPq for financial support to this research, and also gratefully acknowledge Cequil for the materials donated to this project.

The Multi‑User Center for Materials Characterization – CMCM of UTFPR-CT is also greatly acknowledged.

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Degradable polymers: principles and applications. London: Chapman & Hall.

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Received: June 03, 2016 Revised: Aug. 22, 2016 Accepted: Dec. 15, 2016

Polímeros, 27(2), 165-170, 2017

http://dx.doi.org/10.1590/0104-1428.05516

Understanding the water uptake in F-161 glass-epoxy composites using the techniques of luminescence spectroscopy and FT-NIR Rita Sales1,2*, Gilmar Thim2 and Deborah Brunelli2 Faculdade de Tecnologia – FATEC, São José dos Campos, SP, Brazil Chemistry Department, Instituto Tecnológico de Aeronáutica – ITA, São José dos Campos, SP, Brazil 1

2

*rita.sales@fatec.sp.gov.br

Abstract This paper investigates the application of the luminescence spectroscopy technique in steady–state to study the moisture influence in glass fiber/epoxy prepreg and their laminates. The studies were monitored by intrinsic luminescence comparing the results with gravimetric analysis and near infrared with Fourier transform. Samples are cured and submitted to humidity controlled at 60 and 80 °C until 90 days. It is verified that the decrease in the maximum emission of the samples is directly related to the material moisture content. However, for very short periods, there is an increase in the relative intensity and blue shift of the emission band for all samples treated at 60 °C, which is related to an increase of the rigidity of the polymeric matrix. The results in this paper have a great significance because it brings a wide discussion on the interaction of water in epoxy composites materials.

Keywords: F-161 laminates, FT-NIR, luminescence spectroscopy, water uptake.

1. Introduction The composite materials present in the aircraft structures are exposed to environmental conditions with elevated levels of moisture. The composite materials can absorb water molecules and this chemical interaction can result in a delamination process [1]. The delamination process consists on separating the layers resulted from interlaminar tensions.

The study of the water interaction with the polymer matrix at the molecular level is very important since the water uptake will determine the mechanical strength of the material at the macroscopic level. These materials are used in primary structural parts of aircraft under great mechanical stress. The knowledge of the mechanisms which drive moisture sorption, as well as of the influence of sample dimensions, temperature and relative humidity, becomes crucial when the long-term properties of the material are needed[2].

Many researches have been done to understand the behavior of water in epoxy resin and composites, since this cured polymeric matrix contains strongly polar groups, such as hydroxyl groups and amine groups[3-6]. Adamson[6] has done experiments of thermal expansion of two graphite/epoxy composite systems and concluded that water is initially absorbed and diffuses into the resin; then it begins to occupy polymer free volume, causing no swelling, while some of the water disrupts intermolecular hydrogen bonds, causing swelling by hydrogen bonding with the resin. In the final stage of water absorption, the swelling efficiency drops far below that observed in the initial stage. Apicella et al.[2] have analyzed epoxy resin cured with amine compounds using DSC. These authors proposed three sorption mechanisms: (1) bulk dissolution of the water in the polymer network; (2) adsorption on hydrophilic sites and (3) adsorption on the surfaces of free volume elements.

Polímeros, 27(2), 171-182, 2017

Li et al.[7] used time-resolved attenuated total reflection Fourier transform infrared spectroscopy (ATR-IR), near-infrared (NIR) spectroscopy and positron annihilation lifetime spectroscopy (PALS) to investigate water diffusion processes and the state of water molecules in six different epoxy resins. These authors concluded that water diffusion is controlled by local chain reorientation and bond dissociation of water molecules from epoxy networks. Four types of water molecules were detected: nonbonded, single bonded and two types of double hydrogen bonded.

Cotugno et al.[8] e Roy et al.[9] showed that water molecules can affect the epoxy matrices by (a) plasticization, (b) changes in physical properties and (c) hygrothermal degradation. The plasticization effects have been previously discussed. The changes in physical properties induce a decrease of yield strength and change of yield/deformation mechanisms. The hygrothermal degradation is characterized by micro-cracking, aging, breakdown of the polymer chain and the degradation of the fiber-matrix interface in polymer composites.

Usually, water absorption by the epoxy resin decreases the glass transition temperature (Tg) due to the plasticization of the polymer network. This effect can lead to a decrease of the mechanical performance of the composite material.

Those properties related to the compressive strength and intralaminar/interlaminar shear is significantly affected by the moisture absorbed, especially at high temperatures[10-12].

According to Asp[13], the water absorption can cause interlaminar delamination or damage in carbon-epoxy composites.

One of the most important hygrothermal aging processes of the epoxy composites involves molecular interactions between water and epoxy resin, such as the plasticization

171

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

Sales, R., Thim, G., & Brunelli, D. phenomenon of the molecular network. This process may be reversible, but can lead to irreversible loss of mechanical properties[14]. According to Majerus et al.[15], this phenomenon is related to the presence of water molecules in the matrix microcavities, by the interaction between water molecules with specific segments of the polymeric matrix. Epoxy resins can absorb water through three different ways: (a) bulk dissolution of water in polymer network, (b) adsorption on the hydrophilic sites and (c) adsorption in the voids. The transport of moisture below the Tg consists of a three-step process.

First, the water occupies the voids of the polymer matrix.

Next, the water molecules bind to the polymer network by hydrogen bonds, causing swelling. Finally, water diffuses into the densely crosslinked resin. Polarity and topology of the structure of crosslinked epoxy resin influence the adsorption of water[16].

The effects of water absorption on the epoxy resin have been widely investigated by several techniques. However, none of them is sensitive and non-destructive for monitoring the moisture aging effects on composite materials like luminescence spectroscopy[17-21]. Sung and Sung[17,21] used luminescence spectroscopy, intrinsic and extrinsic methods, to characterize the water uptake and the curing extension in composite materials. The intrinsic method involved the study of the emission of the resin matrix itself, while the extrinsic one consisted in adding a luminescent probe in the resin. These researchers noticed that there was some spectral shift in the wavenumber and/or some decrease in intensity of the emission spectra when water molecules were uptaken.

In this work, it was used glass fiber/epoxy prepreg F161 (Hexcel), a composite system that is widely used in aeronautics industry in the manufacture of structural parts. The moisture effect on glass fiber/epoxy prepreg and glass fiber/epoxy laminates was studied by luminescence spectroscopy in steady-state condition using the intrinsic fluorescence method. The results were compared to Fourier Transform Near Infrared Spectroscopy (FT-NIR) and the sample weighting was performed according to ASTM D 5229/D 5229M-12[22]. The changes in the maximum emission of the samples is directly related to the moisture content of the material. Other authors had not previously observed this photophysical behavior.

2. Materials and Methods 2.1 Sample preparation

2.2 Water absorption analysis The water absorption (% H2Oabs) was determined using Equation 1, according to ASTM D 5229/D 5229M-12[22]: %H 2O =

( M f − Mi ) Mi

x100 (1)

where MF and MI are, respectively, the mass of the samples after and before being added into the desiccators.

2.3 Water diffusion The diffusion coefficients are easily calculated from the moisture-uptake profiles with Fickian diffusion[23] (Equation 2): 1/ 2

Mt  Dt  4 2  = M∞ h 

These experiments used a plain wave glass fiber/epoxy resin prepreg, F161 (Hexcel Composites Ltd., Livermor, CA), with size of 30 mm × 20 mm. Some prepreg samples were analyzed as received to understand their photophysical behavior and they were named as btt-pre (btt= before thermal treatment and “pre”= prepreg). The curing process was made according to the manufacturer’s instructions. The prepreg samples were heated from room temperature to 177 °C at a heating rate of 2 °C/min, and then, they were maintained in this temperature for 120 min. After this time, the samples were removed from the oven and maintained at room temperature for 20 min. Samples submitted to this thermal treatment were named as tt-pre (tt=thermal treatment).

172

A set of desiccators was maintained at 60 and at 80 °C. Beckers with saturated aqueous solution of NaBr, NaCl and KCl (Merck) were added to these desiccators to simulate a relative humidity of 58, 75 and 84%, respectively. Pieces of samples already thermally treated were put onto glass plates next to the salt solution inside of each desiccator.

The samples were maintained for 1, 7, 15, 30, 60 and 90 days in each desiccator. The samples obtained after exposure to different environmental conditions were named pre-relative humidity-T-t, where “pre”, relative humidity, T and t are, respectively, prepreg, relative humidity value, temperature and time, for example, pre-75-60-30 stands for a cured F-161 prepreg sample exposed to a relative humidity of 75% at 60 °C for 30 days.

To manufacturing the laminate samples, the prepreg layers were laid down in a plain mold forming a rectangular panel (160 mm × 160 mm) with 14 layers. The curing process was running using a hydraulic press (SOLAB) at a pressure of 50 kgf/cm2 and 177 °C using a heating rate of 2 °C/min and then maintained at this temperature for 120 min. After that, samples were cut in pieces of 25mm × 15mm and added to desiccators for 1, 7, 15 and 30 days. The laminated samples were exposed to the same environmental conditions as the prepreg samples into desiccators. The samples obtained after the exposure to different environmental conditions were named lam-relative humidity-T-t, where lam stands for a cured laminate sample constituted of 14 prepregs.

The remaining labels are similar to pre- relative humidity -T-t that was previously described.

 1 ∞ nh  n   (2) ∑ ( −1) ierfc 1/ 2 1/ 2  π + 2 n =1 2 ( Dt ) 

where MT is the moisture uptake at time t, M∞ is the equilibrium moisture uptake, h is the sample thickness, D is the diffusion coefficient, ierfc is the integral error function of each element and n is the concentration starting at 0 at the surface and spreading infinitely deep in the material. This solution of Fick’s second law holds true for the conditions of an infinite sheet with constant penetrant activity on both sides of the sheet and a concentration independent of D. At very short times, where MT/M∞ is less than 0.5, Equation 2 can be defined approximated by the following (Equation 3): Mt 4 Dt = (3) M∞ h π Polímeros, 27(2), 171-182, 2017

Understanding the water uptake in F-161 glass-epoxy composites using the techniques of luminescence spectroscopy and FT-NIR Equation 3 is easily rearranged to reveal that the initial slope in a plot of Mt/M versus =t/h is related to the diffusion coefficient D through the following (Equation 4): 2

  π  Mt / M∞    (4) D= 16   t   h  

2.4 Infrared spectroscopy in the near region The infrared spectroscopy was also used to analyze the water content in the samples. The samples were analyzed in the near region using a Spectrum One Perkin Elmer spectrometer with a universal attenuated total reflectance (UATR) accessory; the acquisition parameters were: 8000-4000 cm-1, 4 cm-1 resolution, 1 gain and 40 accumulations. The relation between the epoxy group intensity band (4523 cm-1) and reference band (4623 cm-1). The reference band selected is the combination of the two stretching modes, C–H and C=C, of the phenyl group of the epoxy group. They were analyzed in order to determine the water absorption by the samples. The water content (wT) was calculated by Equation 5: = wt 0,096 + 0,587

Aw,t − Aw,0 Aref

(5)

where AW,0 is the integrated absorbance related to the non-exposed sample to moisture, AW,T is the integrated absorbance of sample exposed to moisture for time equals to t and AREF is the integrated absorbance of the reference band. Gonzales-Benito and co-workers described this procedure in details as well the meaning of the constants of the Equation 5[24].

2.5 Emission and excitation spectra Fluorescence spectra were measured with a steady-state luminescence spectrometer (FS920-Edinburgh Analytical Instruments) in the photocounting mode with a xenon arc lamp 450 W (Osram Co.) and double holographic grating monochromators of excitation and emission (Czerny- Turner configurations). The fluorescence spectra of the solid samples were obtained through a front-face illumination.

The range of wavelength of excitation (290-380 nm) was chosen taking into account that the curing agent groups shows emission band after this region. The first moment of luminescence (<ν>), defined by the weighted average wavenumber (Equation 6)[25], was used in order to analyze the water absorption into F-161 prepregs and laminates: ∑ ν i I (ν i ) ν =i

∑ I (ν i )

(6)

i

where I(νI) is the relative emission intensity at wavenumber νI.

3. Results and Discussions 3.1 F-161 prepreg samples under different humidities at 60 and 80 °C Figure 1A shows the spectra of electronic emission of the samples pre-84-80-t. These samples presented two superimposed emission bands: one centered at range 451-455 nm and another centered at 488 nm. The emission bands in the region from 400-600 nm of pre-relative humidity-60-t samples can be attributed to the curing agent, which probably contains amine groups[26]. The band profile is very similar to the emission band of the cured prepreg sample. Sales et al.[26] have discussed that the spectra of the

Figure 1. Electronic emission spectra of the samples (λexc = 370 nm): (A) pre-84-60-t, (B) pre-75-60-t and (C) pre-58-60-t in the following periods of time (t in days): (a) 1, (b) 7, (c) 15, (d) 30, (e) 60 and (f) 90.

Polímeros, 27(2), 171-182, 2017

173

Sales, R., Thim, G., & Brunelli, D. samples non-totally cured showed a red shift of maximum emission due to the conversion of primary amines into secondary and tertiary amines. As might be expected, the sample pre-84-80-1 (Figure 1A.a) showed no significant influence of water absorption in the polymeric matrix.

The band centered at 488 nm is the most intense one for samples with t=1, 7, 60 and 90 days, while the band centered in the range 451-455 nm is the most intense for samples with t=15 and 30 days. Another remarkable observation is that the displacements were dependent on the exposure time. Therefore, one can observe a displacement towards to the blue region for exposure time from 7 to 15 days and another displacement towards to the red region for periods of time of 30 to 60 days.

The emission spectrum of pre-84-60-7 (Figure 1A.b) does not present a significant change when compared to the pre‑84-60-1. However, the spectrum of the sample pre-84-60-15 (Figure 1A.c) when compared to that sample pre-84-60-7 shows a considerable increase in the relative intensity of the maximum emission and a blue shift to 455 nm. The relative intensity of the emission band of sample pre-84-60-30 (Figure 1A.d) is lower than samples pre-84-60-15 and its maximum is centered at 455 nm. Sample pre-84-60-60 and pre-84-60-90 show emission bands whose profiles are very similar to the band of the pre-84-60-1, but there is a decreasing of the relative intensities. Emission spectra of samples pre-75-60-t (Figure 1B) and pre-58-60-t (Figure 1C) are very similar with spectra of the pre-84-60-t samples; the only difference is related to exposure time where the blue shift is observed. The emission spectrum of sample pre-84-60-t shows a shift towards the blue region with 15 and 30 days, while those related to samples pre-75-60-t and pre-58-60-t show the same shift after a period of time of 7, 15 and 30 days.

These two effects on the photophysical behavior of samples subjected to moisture at intermediate times can be related.

The relative emission intensity increases when there is an increase in the radiative rate constant of the fluorophores, which can be attributed to the stiffness of the polymeric matrix. On the other hand, the blue shift can be related to an important change of the polarity of the medium, which could influence the fluorophore photophysical behavior.

According to Lakowicz[27], a decrease in the medium polarity can lead to a blue shift of maximum emission. On other hand, the interaction with a more polar solvent could lead to the stabilization of the excited electronic state, shifting the emission maximum to the red region. Since the emission band was attributed to secondary and tertiary amine compounds, one can assume that water molecules diffuse through the polymeric matrix and form hydrogen bonds with the amine and hydroxyl groups. The effect of water diffusion through the polymeric matrix in the photophysical behavior of the fluorophores is the increase in energy of its excited electronic state and consequent shift of the maximum emission to the blue region. The shift to the blue region could indicate that the formation of hydrogen bonds of water with polar groups reduces its polarity. In this case, water would act as a pseudo curing agent and would contribute to an increase in the stiffness of the polymeric matrix, which is related to an increase in the rate constants of radiative processes and 174

consequent increase in emission intensity, as observed in the spectra of samples.

Novolac epoxy matrix of F-161 prepreg type is a very polar material, since it contains a considerable number of hydroxyl[25]. These groups result from the opening of the epoxy rings by crosslinking reaction. The water molecule could act as a pseudo-curing agent by its interaction with hydroxyl groups of the different polymer chains. This interaction increases the polymer matrix stiffness and increases the radiative decay rate constant. Consequently, an increase in the relative intensity of the luminescence emission is observed for very short times of exposure to moisture.

According to Soles et al.[28], resin polarity is of primary importance in determining the equilibrium moisture uptake.

Less polar resins, such as the non-amine series, absorb very little water compared to the amine-containing materials.

Accompanying polarity, topology has a secondary effect on uptake. A greater intrinsic hole volume fraction (V0) leads to an increase in the equilibrium moisture content.

V0 is isothermally elevated by increasing crosslink density, the consequence of which is enhancement of the effective polarity. Steric hindrances at the crosslink junctions introduce unoccupied volume elements coincident with the polar hydroxyls and amines. V0 is localized to regions immediately adjacent to crosslink junctions. Calculations indicate there will be approximately 7×1020 nanopores per cm3 for epoxy matrix and 2×1021 polar groups per cm3. If a typical equilibrium uptake (for the bifunctional resins) of 3 wt% is assumed, approximately 1×1021 water molecules per cm3 of resin will be absorbed[28,29].

The decrease in the relative intensity of samples treated for longer periods of time (30 and 60 days) denotes an increasing of the rate constant of non-radiative decay that can be related to the plasticization effects of the polymer matrix due to water uptake. This process is associated to the amount of unbound water that occupies the free volume of the polymeric matrix. On the other hand, higher water content in the bulk increases the medium polarity, resulting in the red shift of the emission band. Sung and Sung et al.[19] verified that water absorption induces changes in the medium polarity and a decrease in the relative intensity of the emission band luminescence. According to Zhou and Lucas[30], there are two ways that water molecules can bind to the epoxy matrix. In the first one, the absorbed water can diffuse into the polymer chain breaking the van der Waals bonds, resulting in an increase of the chain segment mobility and swelling. Water molecule interacts strongly with hydrophilic functional groups such as hydroxyl or amine in epoxy resin.

The second one occurs when the absorbed water molecule connects to the polymer chain forming pseudo or secondary crosslinking. This process turns the secondary amine groups into tertiary amines, increasing the conversion degree.

The results of this work strongly suggest that water first acts as a pseudo curing agent, since it was observed an increasing of the emission intensity. Then occurs plasticization of the polymeric matrix related to decreasing of the emission intensity. It must be clear that this photophysical behavior was observed for all samples submitted to different relative humidities treated at 60 °C. Polímeros, 27(2), 171-182, 2017

Understanding the water uptake in F-161 glass-epoxy composites using the techniques of luminescence spectroscopy and FT-NIR Prepreg samples submitted to relative humidities of 84 (Figure 2A), 75 (Figure 2B) and 58% (Figure 2C) at 80 °C showed emission bands with the same profiles and a wavelength maximum at 488 nm, a band overlapped around 450 nm and there are no shift of the maximum emission bands. Comparing these spectra with those obtained for samples submitted to the same relative humidity at 60 °C, it is observed that the bands have the same emission band profile than those submitted for longer times, in which the water content is higher.

First moment of fluorescence (Equation 6) allows the comparison of the emission spectra of the samples exposed to three conditions of relative humidity for different periods.

This equation has been used for many authors not only to study the influence of water in polymeric matrix[24], but also to study polymerization reaction[21,31] and investigate the structure of elastomeric network[32]. Figure 3A shows the first moment (<ν>) versus exposure time for samples pre-84-60-t, pre-75-60-t and pre-58-60-t. For a better evaluation of water effects in epoxy matrix, all scales are in the same range. The sample pre-58-60-t showed the highest value for the luminescence band first moment (<ν>) and the lowest one for water content. The variation of first moment of samples pre-58-60-1 and pre-58-60-7 has a significant difference of 857 cm-1, indicating that the absorbed water has an influence in polymeric matrix.

On the other hand, the values of water content and luminescence band first moment (<ν>) for the samples pre-84-60-t and pre-75-60-t are very close (Figure 3A).

The effect of increased emission intensity is much more pronounced in samples pre-58-60-7, pre-58-60-15 and

pre-58-60-30 compared with the other samples (Figure 3A).

However, the results of gravimetric analysis showed water contents between 0.1 and 0.3% m/m, which are half the water content of the samples exposed to relative humidity of 84% for the same periods of time (Figure 3B). The photophysical behavior of the samples pre-58-60-t can indicate that the influence of bound water is significant only for lower water contents. For higher water contents, one can conclude that the <ν> variation is inversely proportional to the water content.

Figures 3A and 3B show the first moment of luminescence (<ν>) and water content of the samples pre-58-80-t, pre-75-80-t and pre-84-80-t, respectively. It can be observed that <ν> does not change significantly for samples submitted at 80 °C, in relation to the same parameter for samples treated at 60 °C. According to Sales[33], the increase in the relative luminescence emission intensities and in the rate constant of radiative decay is due to the increased rigidity of the polymer matrix. Probably, the water forms hydrogen bonds with polar groups of the polymer matrix, acting as pseudo-curing agent by exposure time periods for each specific humidity, but as the rate of diffusion of water in the polymer matrix should be greater than 80 °C (Table 1), it is assumed that the matrix plasticization occurs in a shorter time period (between 7 to 30 days) when the samples are submitted to the higher relative humidity and temperature.

Table 1 summarizes the wavelengths of the emission bands, water contents, first moments of luminescence and diffusion coefficients of prepreg F161 samples at 60 and 80 °C. Sample pre-58-60 treated for 15 days presented the higher first moment of luminescence – 21986 cm-1. The emission maximum of this sample was shifted to the blue region at

Figure 2. Electronic emission spectra of the samples (λexc = 370 nm): (A) pre-84-80-t, (B) pre-75-80-t and (C) pre-58-80-t in the following periods of time (t in days): (a) 1, (b) 7, (c) 15, (d) 30, (e) 60 and (f) 90.

Polímeros, 27(2), 171-182, 2017

175

Sales, R., Thim, G., & Brunelli, D.

Figure 3. F-161 prepreg samples at 60 and 80 °C: (A) First moment of luminescence <v> versus exposure time and (B) Dependence of the water content with exposure time.

Table 1. Summary of wavelengths of the emission bands, water content, first moment of luminescence and diffusion coefficients of prepreg F161 samples at 60 and 80 °C. Samples

λmax (nm) 452

pre-84-60-ta

488 1 7

15b 30 60 90 1

pre-75-60-t

7b 15 30 60 90 1

pre-58-60-t

7b 15 30

pre-84-80-ta

pre-75-80-t

pre-58-80-t

60 90 1 7b 15 30 60 90 1 7 15b 30 60 90 1 7 15 30b 60 90

Water content (ΔMrel %)

First moment of fluorescence (cm-1)

0.00 0.55 0.54 0.49 0.52 0.57 0.08 0.35 0.44 0.49 0.50 0.59 0.08 0.18 0.24 0.27 0.38 0.39 0.56 0.28 0.28 0.03 0.40 0.51 0.14 0.22 0.24 0.31 0.43 0.65 0.12 0.07 0 0.28 0.07 0.26

20969 20926 21556 21240 20927 20975 20775 21469 21413 21383 20811 20757 21100 21957 21986 21922 21005 21066 20909 20816 21115 21049 21023 21043 20957 20900 20946 21020 21003 20913 21203 21134 21081 21035 21134 21089

Diffusion coefficient (cm2/min)

1.22475x10-7

2.01633x10-8

2.0742x10-8

5.1126x10-8

1.61816x10-9

1.8963x10-8

t = exposure time; bmost intense band.

a

176

Polímeros, 27(2), 171-182, 2017

Understanding the water uptake in F-161 glass-epoxy composites using the techniques of luminescence spectroscopy and FT-NIR 452 nm and was the more intense one. The water content and diffusion coefficient were 0.24% and 2.0742x10-8 cm2/min, respectively. Samples treated at 75 and 84% of relative moisture at 60 °C for 15 days presented maximum first moments of luminescence of 21413 and 21556 cm-1, both lower than that of the sample pre-58-60-15. The water contents for samples pre-75-60-15 and pre-84-60-15 were of 0.44 and 0.54%.

The water content increases as the increase of the relative humidity, as it was expected, but decreases the effect of water uptake in the first moment of luminescence. Samples pre-58-60-t and pre-75-60-t presented diffusion coefficients very close, but the first sample had the effect of increasing the emission intensity more pronounced. Results indicate that the sample contains a significant amount of bound water, since there is an increased relative intensity of the emission band related to the increased rigidity of the polymer matrix and shift of the emission band to the blue region, due to hydrogen bonding of water with amine and hydroxyl groups.

All the samples treated at 80 °C present emission maxima at 488 nm and values of first moment of luminescence very close. Emission spectra of the samples pre-84-80-7 and pre-84-80-15 show an increase of the relative intensity of the band that could indicate an increase of the stiffness of

the polymeric matrix, but there is no noticeable change in the polarity of the medium, since it was not a blue shift.

Values of diffusion coefficient were to close too. It should be emphasized that the gravimetric measurement of the sample pre-58-80-15 is unreliable since there is no mass change.

It must be assumed that the samples treated at 80 °C do not contain detectable amounts of bound water. Soles et al.[29] have discussed that this behavior is partially understood in terms of the interactions between water and a polar group.

Increasing the temperature increases the external heat and shifts the above reaction toward the formation of free water.

The analysis of the luminescence spectra of the samples treated at 80 °C confirms the conclusions of those authors.

There is a great amount of free water in the samples that maintains a higher polarity of polymeric matrix.

3.2 F-161 laminated samples under different humidities at 60 °C e 80 °C Figure 4A shows the intrinsic luminescence spectra of samples lam-84-60-t. The spectra related to samples lam-75-60-t (Figure 4C) and lam-58-60-t (Figure 4E) are quite similar to those of pre-84-60-t (Figure 1A). All spectra

Figure 4. Electronic emission spectra of the samples (A) lam-84-60-t, (B) lam-84-80-t, (C) lam-75-60-t, (D) lam-75-80-t, (E) lam-58-60-t and (F) lam-58-80-t in the following periods of time (in days): (a) 1, (b) 7, (c) 15 and (d) 30.

Polímeros, 27(2), 171-182, 2017

177

Sales, R., Thim, G., & Brunelli, D. show the emission bands with the same profile and the wavelength of the maximum emission around 450 nm, as well as the samples pre-84-60-t treated for 1, 7, 60 and 90 days.

These bands are attributed to the emission of secondary and tertiary amine groups of the curing agent, which are covalently bound to the epoxy matrix. The emission band at 450 nm can indicate a reducing of the medium polarity related to hydrogen bonding of water with polar groups.

All intrinsic luminescence spectra related to samples lam-84-80-t, lam-75-80-t and lam-58-80-t are quite similar and show emission bands with the same profile and the maximum emission of the sample lam-84-60-t do not showing significant changes.

Figures 5A and 5B show the luminescence first moment and the water content variations with the exposure time of the samples lam-84-60-t, lam-75-60-t, lam-58-60-t, lam-58-80-t, lam-75-80-t and lam-84-80-t, respectively. The dependence the values of first moment luminescence (<ν>) of the three samples on the root square of the exposure time is very small. The biggest difference between the higher value and the lower one was equal to 142 cm-1, while the curve related to the prepreg samples showed the major difference equal to 982 cm-1. This behavior was expected, since the amount of the free volume decreased during the curing and lamination process. This increase in the laminate bulk density decreases the capacity of water absorption.

However, the matrix polarity does not change during the experiments, but the emission maximum is located at 448 nm and is attributed to bound water. This same photophysical behavior was observed mainly for the samples pre-58-60-7, pre-58-60-15, pre-58-60-30 and pre-75-60-7 and indicated the presence of bound water.

Table 2 summarizes wavelengths of the emission bands, water content, first moment of luminescence and diffusion coefficients of laminated F161 samples at 60 and 80 °C. All spectra showed a maximum emission at 448 nm that indicate a less polar medium due to hydrogen bonding of water to amine and hydroxyl groups. However, samples of laminates not showed a significant increase in the relative emission intensity as observed for samples of prepreg. It can be inferred that the amount of bound water in the laminate is not enough to increase the rigidity of the polymer matrix.

Water contents confirm this conclusion since they were much

smaller compared with the samples of prepreg, as it was expected since laminated samples have a lower free volume.

The values of the first moment of fluorescence showed no significant variations too. The values of the diffusion coefficients ranged from 2.021x10-5 to 4.28x10-5 cm2/min for samples lam-84-60, lam-75-60 and lam-58-60. Samples treated at 80 °C had diffusion coefficients varying from 4.9537x10-7 to 1.233x10-5 cm2/min.

3.3 FT-NIR spectroscopy Figure 6 shows FT-NIR spectrum of samples btt-pre, pre-84-60-90 and pre-84-80-90. Table 3 shows the main absorption bands, including epoxy, methylene, and phenyl groups. These assignments agreed with those from literature[34] related to epoxy resin. Peaks at 6064 cm-1 and at 4524 cm-1 are related to the epoxy ring[35] and they can be attributed to[35,36]: i) the first overtone of the terminal CH2 stretching mode; ii) to a combination band of the second overtone of the epoxy ring stretching at 916 cm-1 with the fundamental C-H stretching at about 2725 cm-1 . Table 3 shows the attribution of the vibrational modes of the FT-NIR bands of the prepreg F-161. The band at 4524 cm-1 can be attributed to a combination of the stretching fundamental (3050 cm-1) with the CH2 deformation fundamental (1460 cm-1) of the epoxy ring[37-39].

The band at 5238 cm-1 is related to water absorption process by epoxy resin[40] and can also be attributed to a combination band of groups -CH and -CH2[41]. The intensity of the band at 5238 cm-1 of the sample exposed to moisture at 60 and 80 °C (Figure 6B) did not show an increase of intensity when compared with the btt-pre band; this band did not make the water absorption analysis possible.

The Novolac epoxy resin structure shows an oxirane group attached to each benzene ring, consequently the density of OH group of this resin is higher than other epoxy resins.

Therefore, the band at 4524 cm-1 is the best for monitoring the water uptake in the polymeric matrix, since it is related to the epoxy ring. According Gonzalez-Benito et al.[24], when water molecule enters into the epoxy resin, some molecules may interact with the functional groups of resin, since the resin may not be fully cured[24]. Water molecule is considered a curing agent of the “Lewis base” type[42].

The anion generated by the reaction between the amine and

Figure 5. F-161 laminated samples at 60 and 80 °C: (A) First moment of luminescence <v> versus exposure time and (B) Dependence on the water content with exposure time.

178

Polímeros, 27(2), 171-182, 2017

Understanding the water uptake in F-161 glass-epoxy composites using the techniques of luminescence spectroscopy and FT-NIR Table 2. Summary of wavelengths of the emission bands, water content, first moment of luminescence and diffusion coefficients of laminated F161 samples at 60 and 80 °C. Samples

lam-84-60-t

a

lam-75-60-t

lam-58-60-t

lam-84-80-ta

lam-75-80-t

lam-58-80-t

λmax (nm)

Water content (ΔMrel %)

First moment of fluorescence (cm-1)

448 1 7b 15 30 1 7 15b 30 1b 7 15 30 1 7 15b 30 1 7 15 30b 1 7b 15 30

478 0.044 0.098 0.087 0.131 0.040 0.234 0.192 0.282 0.000 0.120 0.135 0.184 0.091 0.241 0.255 0.628 0.091 0.422 0.390 0.417 0.098 0.325 0.361 0.416

21483 21502 21443 21584 21536 21644 21582 21510 21486 21510 21566 21584 21562 21563 21452 21520 21599 21540 21498 21506 21486 21510 21566 21584

Diffusion coefficient (cm2/min)

3.3803.10-5

2.0220.10-5

4.2864.10-5

4.9537.10-7

1.2333.10-5

5.8829.10-6

t = exposure time; bmost intense band.

a

Figure 6. FT-NIR spectra of samples: (a) btt-pre (b) pre-84-60-90 and (c) pre-84-80-90.

the epoxy ring can react with an active hydrogen present in water, alcohol, phenol or carboxyl to form a new anion.

This anion is able to open a second epoxy ring continuing the curing reaction[10,42-46].

Figure 7A shows the absorbed water content by the samples of F-161 prepreg samples submitted to moisture of 84, 75 and 58% at 60 °C. There is a significant increase in the water content during periods of time1/2 of 2.6, 3.9 and 5.5.

In the following periods of time1/2 there is an abrupt decrease in intensity, because of the stabilization of the quantity of Polímeros, 27(2), 171-182, 2017

absorbed water by samples. This behavior suggests that in the early periods, the water binds to epoxide groups increasing the intensity of the band at 4524 cm-1 and, in the subsequent period of time, water molecule acts as a curing agent, promoting the opening of the epoxide ring supporting the curing process. Comparing the results obtained by the weight measurements and water content measured by FT-NIR, one can observe the same behavior: an increase in the earlier periods of time due to water absorption and, in the subsequent periods a stabilization of water content due to the crosslink process promoted by water.

179

Sales, R., Thim, G., & Brunelli, D. Table 3. Attribution of the normal mode vibrations of the NIR absorption spectra of prepreg samples.

Wavenumber (cm-1) 6071 5961 5886 5661 5247 4672 4625 4524 4480 4362 4216 4155 4058

Band assignments First overtone of terminal (methylene)-CH fundamental stretching vibration (epoxy ring).

Overtone of the CH axial deformation of the phenyl group.

First overtones of the fundamental –CH2 and –CH stretching vibration.

-OH combination band.

-CH, -CH2 combination band.

Combination band of the aromatic conjugated C=C stretching (=1625 cm-1) with the aromatic –CH fundamental stretching (=3050 cm-1).

Combination band of the aromatic conjugated C=C stretching (=1625 cm-1) with the aromatic –CH fundamental stretching (=3050 cm-1).

Conjugated epoxy CH2 deformation band (=1460 cm-1) with the aromatic –CH fundamental stretch (=3050 cm-1), -NH2 absorption band.

-CH, -CH2 combination band.

-CH2 combination band.

-CH, -CH2 combination band.

Aromatic –CH combination band.

Aromatic combination band.

Figure 7. Water content versus time measured by FT-NIR: (A) F-161 prepreg samples at 60 °C for the following periods of time (in days1/2): 1.0, 2.6, 3.9, 5.5, 7.7 and 9.5 and (B) F-161 laminates samples at 60 °C, for the following periods of time (in days1/2): 1.0, 2.6, 3.9 and 5.5.

Figure 7B shows the amount of water absorbed by F-161 laminates samples submitted to relative humidity of 84, 75 and 58% at 60 °C measured by FT-NIR, respectively.

It can be claimed that no significant changes in the water content of the laminates (Figure 5). This behavior was expected because laminate samples were already cured and this process decreases free volume of polymeric matrix.

Therefore, there is a greater compression of the polymer chains and less capacity to absorb water.

4. Conclusions All samples of prepreg F-161, independent on relative humidity and exposure time, showed very similar photophysical behavior: (a) a blue shift of the maximum emission and (b) an increase in relative intensity of luminescence.

This photophysical behavior can be attributed to a change of the medium polarity.

The decrease in the relative intensity of luminescence in the spectra of samples treated for longer times (30, 60 and 90 days) may be related to the plasticization process of the polymer matrix performed by water molecules. The results of this 180

work strongly suggest that water first acts as an pseudo curing agent, since it was observed an increasing of the emission intensity and then occurs plasticization of the polymeric matrix related to decreasing of the emission intensity. It must be clear that this photophysical behavior was observed for all samples submitted to different relative humidities. This photophysical behavior had not been previously observed by other authors The first moment of luminescence (ν) variation is inversely proportional to the quantity of water absorbed by the polymer matrix. Prepreg samples results indicated that the polarity of prepreg samples change during the experiments, characterized by the red shift of the emission spectra. On the other hand, it is not observed either a red shift emission or an increasing of water content in laminated samples, due to the curing process conditions of temperature and pressure.

The results provided by FT-NIR of samples treated for short periods of time showed an increase in the intensity of the band 4524 cm-1. This increase is related to hydrogen bonding between the absorbed water molecules and epoxy groups and confirms the presence of bound water observed through luminescence spectroscopy.

Polímeros, 27(2), 171-182, 2017

Understanding the water uptake in F-161 glass-epoxy composites using the techniques of luminescence spectroscopy and FT-NIR

5. Acknowledgements The authors thank Instituto de Aeronáutica e Espaço (IAE), Fundação de Apoio à Pesquisa do Estado de São Paulo (FAPESP) and Coordenação de Aperfeiçoamento de Nível Superior (CAPES) (Brazil).

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Received: May 23, 2016 Revised: Aug. 19, 2016 Accepted: Dec. 20, 2016

Polímeros, 27(2), 171-182, 2017

waters.com

40.00

35.00

APC vs. GPC com padrão de poliestireno Mp = 510 30.00

µRIU

25.00 40.00 20.00 35.00 15.00 30.00

µRIU

10.00 25.00 5.00 20.00 0.00 15.00 1.60

1.80

2.00

2.20

2.40

2.60

10.00

APC 2.80

3.00

3.20

3.40

3.60

3.80

4.00

3.00

3.20

3.40

3.60

3.80

4.00

Minutes

5.00

0.00 1.60

1.80

2.00

2.20

2.40

2.60

2.80

Minutes 22.0

gPC

20.0 18.0 16.0

µRIU

14.0 22.0 12.0 20.0 10.0 18.0 8.0 16.0

µRIU

6.0 14.0 4.0 12.0 2.0 10.0 0.0 8.0 4.00

4.20

4.40

4.60

4.80

5.00

5.20

5.40

5.60

5.80

6.00

6.20

6.40

6.0

Minutes

4.0 2.0 0.0

O CAMINHO PARA A INOVAÇÃO. 4.00

4.20

4.40

4.60

4.80

5.00

5.20

5.40

5.60

5.80

6.00

6.20

6.40

Minutes

COM ADVANCED POLYMER CHROMATOGRAPHY, É MUITO FÁCIL.

Mais informação sobre seus polímeros em menos tempo. Sempre. É o que se precisa para inovar no mercado químico. E é exatamente o que você pode esperar desse equipamento verdadeiramente único e das novas químicas de colunas do novo sistema Waters ACQUITY ®

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