Polímeros: Ciência e Tecnologia (Polimeros) 3rd. issue, vol 27, 2017

PolĂ­meros

Water polymer

VOLUME XXVII - Issue III - JULY/SEP - 2017

Superabsorbents as soil conditioner

Biosorbents as soil and water recovery

Polymers and Agriculture

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2.0 0.0

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http://dx.doi.org/10.1590/0104-1428.2703

Editorial August 17th, 2017 Dear readers, authors, reviewers and contributors.

In the last years, our journal Polímeros undertook a series of changes and adaptations to new events. With the lack of financial support, Polímeros discontinued the printed version but also altered the web submission platform and introduced new policies to attract authors. The latter actions placed our journal on a new level in relation to the international scientific community. Because of this effort, we have received submissions of high-level manuscripts from Brazilian authors and from different countries; a competent group of referees reviewed these manuscripts. This effort resulted in the recent increase in our impact factor (see below). We expect to maintain this trend in the coming years with the continuous contribution of all the polymer science and technology community.

Marco Aurelio De Paoli President of the Editorial Council of Polímeros

Polímeros, 27(3) , 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

ISSN 0104-1428 (printed)

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

ISSN 1678-5169 (online)

P o l í m e r o s - I ss u e I 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: September 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 3 (Jul./Ago./Set. 2017) ISSN 0104-1428 ISSN 1678-5169 (electronic version)

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1. Polímeros. l. Associação Brasileira de Polímeros.

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Polímeros, 27(3), 2017

Editorial Section Editorial................................................................................................................................................................................................E1 News....................................................................................................................................................................................................E4 Agenda.................................................................................................................................................................................................E5 Funding Institutions.............................................................................................................................................................................E6

O r i g in a l A r t ic l e Elastic modulus of PVDF with bentonite or LiNbO3 using deformation energy

Carlos Alberto Fonzar Pintão and Celso Xavier Cardoso.............................................................................................................................. 183

Sulfonation degree effect on ion-conducting SPEEK-titanium oxide membranes properties Jacqueline Costa Marrero, Ailton de Souza Gomes, Wang Shu Hui, José Carlos Dutra Filho and Vivianna Silva de Oliveira..................... 189

Effect of solvents on the morphology of PMMA films fabricated by spin-coating Giovana da Silva Padilha, Virginia Mansanares Giacon and Julio Roberto Bartoli...................................................................................... 195

Evaluation of the mechanical and thermal properties of PHB/canola oil films Cláudia Daniela Melo Giaquinto, Grasielly Karine Martins de Souza, Viviane Fonseca Caetano and Glória Maria Vinhas....................... 201

Epoxidized natural rubber and hydrotalcite compounds: rheological and thermal characterization Vanessa Macedo da Silva, Regina Célia Reis Nunes and Ana Maria Furtado de Sousa................................................................................ 208

Influence of addition of silanized nanosilica and glycerol on hydrophobicity of starch using a factorial design Fernando Luis Panin Lopes, Vicente Lira Kupfer, Júlio César Dainezi de Oliveira, Eduardo Radovanovic, Andrelson Wellington Rinaldi, Murilo Pereira Moisés and Silvia Luciana Favaro......................................................................................... 213

Cowper-Symonds parameters estimation for ABS material using design of experiments with finite element simulation Alexandre Luis Marangoni and Ernesto Massaroppi Junior........................................................................................................................... 220

Chemical resistance of core-shell particles (PS/PMMA) polymerized by seeded suspension Luiz Fernando Belchior Ribeiro, Odinei Hess Gonçalves, Cintia Marangoni, Günter Motz and Ricardo Antonio Francisco Machado...... 225

Evaluation of degree of conversion, microtensile bond strength and mechanical properties of three etch-and-rinse dental adhesives Samantha Ariadne Alves de Freitas, Marco Daniel Septimo Lanza, Karina Kato Carneiro, Alessandro Dourado Loguercio and José Bauer....................................................................................................................................................................................................... 230

Thermal and mechanical properties of bio-based plasticizers mixtures on poly (vinyl chloride) Boussaha Bouchoul, Mohamed Tahar Benaniba and Valérie Massardier....................................................................................................... 237

R e vi e w A r t ic l e Carbon nanotube buckypaper reinforced polymer composites: a review Bruno Ribeiro, Edson Cocchieri Botelho, Michelle Leali Costa and Cirlene Fourquet Bandeira.................................................................. 247

Polymers and its applications in agriculture Priscila Milani, Débora França, Aline Gambaro Balieiro and Roselena Faez............................................................................................... 256

Cover: Graphical Abstract Arts by Editora Cubo.

Polímeros, 27(3), 2017

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

Global thermoplastic polyolefin market to register CAGR of 7.1% to reach US$3.7 bln in 2022

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

The global thermoplastic polyolefin market was valued at US$2.4 bln in 2016 and is expected to reach US$3.7 bln in 2022, growing at a Compound Annual Growth Rate CAGR of 7.1% between 2016 and 2022, as per Zion Market Research. Thermoplastic polyolefin is a light weight material which is a major product in thermoplastic elastomers family.

It has lowest specific gravities compared to all other thermoplastic elastomers. TPO is extensively used in automobile and construction industry. In the automotive industry, it is used in interior and exterior components. In the construction industry, TPO is widely used in the production of roofing membrane systems and door panels. It is also used in various end-users like industrial, medical, footwear and rubber goods. Also, the weight reducing property of TPO enhances the vehicle performance in terms of speed, fuel efficiency, and handling.

Advancement in the TPO industry is a major growth factor for global thermoplastic polyolefin market. However, rapid growth in the automotive industry along with a rise in building and construction activities is driving the demand for TPO in the last few years. The TPO market is likely to grow at a healthy growth rate in coming five to six years owing to the growing sales of tires and other rubber products from automotive industry. However, technological challenges can hamper the TPO market. Fluctuating prices of raw materials can slow down the market of TPO near the future. Emerging applications of TPO is expected to serve as an opportunity for global thermoplastic polyolefin market.

The global market for TPO is segmented on the basis of end – user application and region. According to the application, it is divided into automotive, building & construction, home appliances, medical, industrial, footwear, and others. Automotive is the largest end-use industry for TPOs and growth of the automotive industry is proportional to the global TPO market’s growth. The growth of automotive industry mainly in the emerging economies is likely to drive the future demand for TPOs. North America led the market in terms of volume.

The market size of TPO in North America is expected to show a significant growth over 2017-2022 owing to rapidly growing automotive, construction, packaging and medical industries in the region. U.S is likely to contribute considerably towards the regional growth over the forecast period. Europe’s TPO market is predicted to record a CAGR of 4.9% over 2017 to 2022 due to its growing applications in packaging and healthcare industries. UK, France, Germany, and Italy are likely to be key revenue pockets of the region.

Central & South America TPO market is expected to witness highest gains of 6.9% over the forecast period.

Asia Pacific is estimated to be the largest market for TPO in coming period. Rise in application scope in automotive and construction sectors of the emerging markets in Asia Pacific region, especially China and India is likely to be a growth factor for TPO. China is the largest automobile manufacturing country in the world. Moreover, markets in India and Brazil are considered to highly profitable for automobile manufacturing both in terms of production and consumption. The growing automobile market in E4

these countries is expected to boost the market for TPOs over the coming period. Asia Pacific market is anticipated to exceed 1.2 billion by 2022. Increasing expenditure in primarily in Japan, China and India are likely to fuel the products demand over coming period. The market for TPO is emerging in Latin America and is expected to hold a strong CAGR due to growing construction industry. South America (CSA) is the fastest regional segment holding above 6.8% CAGR. The region is observing massive product application in packaging and medical packaging.

Middle East & Africa are expected to experience a substantial growth in TPO market in coming years owing to a rapid increase in urbanization leading to rising in the construction industry.

Key industry players include Sumitomo Chemical Company, Arkema S.A., LyondellBasell Industries, ExxonMobil Corporation, Dow Chemical Company, S&E Specialty polymers, DuPont, INEOS Chemicals Company, Saudi Aramco Company, A. Schulman, Mitsui & Company Limited, SABIC Chemical manufacturing company, Noble polymers, and Polisystem UK Limited.

Source: Plastemart - www.plastemart.com

Braskem and A. Schulman Enter Into Partnership to Use Green Polyethylene in Rotomolding Processes Braskem, the largest petrochemical company in the Americas, has entered into a partnership with A. Schulman, Inc., a leading global producer of high-performance plastic compounds and resins, to produce and market a new Green Polyethylene application: a solution for the rotomolding process.

A. Schulman will bring this solution to the market by featuring the “I’m greenTM” seal, which indicates its helping to reduce greenhouse gas emissions.

After identifying a market demand for more sustainable solutions in rotomolded products, Braskem started to develop a solution to enable the rotational molding of general-purpose parts, with applications ranging from toys and furniture to agricultural tools that can contain more than 50% of Green Plastic in their composition.

A. Schulman, which contributes to the partnership through its industrial and commercial expertise in serving clients directly with products that meet market needs, will introduce the product at Rotoplas 2017, the largest trade fair of the rotomolding industry, which takes place from September 26-28, 2017 in the United States.

“The partnership with A. Schulman will benefit a market that requires innovative products. The new compound is another step of the petrochemical industry towards reinforcing the commitment of companies to new solutions that help reduces greenhouse gas emissions,” said Gustavo Sergi, Director of Renewable Chemicals, Braskem.

“A. Schulman is honored to have a long-standing collaborative relationship with Braskem and we are equally pleased to play a part in helping drive green innovation the specialty chemical industry and specifically for the rotomolding market,” said Gustavo Perez, Senior Vice President and General Manager - Latin America, A. Schulman.

Source: Braskem - www.braskem.com.br

Polímeros, 27(3), 2017

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 PLASTEC West Date: February 6–8, 2018 Location: Anaheim – USA Website: plastecwest.plasticstoday.com Caribbean Conference on Functional Materials (CARIBMAT 2018) Date: February 7–9, 2018 Location: Cartagena de Indias - Colombia Website: www.caribmat.org 10th International Plastics Exhibition, Conference & Convention (PLASTINDIA 2018) Date: February 7–12, 2018 Location: Gujarat – India Website: www.plastindia.org/plastindia-2018/index.html

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 5th International Conference on Plastics, Rubber and Composites (ICPRC 2018) Date: May 4–5, 2018 Location: Phuket - Thailand Website: www.icprc.org The International Plastic Showcase Date: May 7–11, 2018 Location: Orlando - USA Website: www.npe.org 34th International Conference of the Polymer Processing Society Date: May 21-25, 2018 Location: Taipei - Taiwan Website: www.pps-34.com.

World Congress on Biopolymers and Polymer Chemistry Date: May 28–30, 2018 Location: Osaka - Japan Website: biopolymerscongress.conferenceseries.com

June

6th Asian Symposium on Emulsion Polymerization and Functional Polymeric Microspheres Date: March 7-10, 2018 Location: Fukui - Japan Website: matse.u-fukui.ac.jp/~pm/asepfpm6/index.htm Plastics Regulations Date: March 14–15, 2018 Location: Cologne - Germany Website: www.amiplastics.com/events/event?Code=C874 Conductive Plastics Date: March 20–21, 2018 Location: Pittsburgh - USA Website: www.amiplastics.com/events/event?Code=C892 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

Polymers and Organic Chemistry (POC 2018) Date: June 4–7, 2018 Location: Montpellier - France Website: iupac.org/event/polymers-organic-chemistry-2018poc-2018 4th Functional Polymeric Materials Conference Date: June 5–8, 2018 Location: Nassau - Bahamas Website: https://www.fusion-conferences.com/conference76.php PLASTEC East Date: June 12-14, 2018 Location: New York – USA Website: plastec-east.plasticstoday.com Polymer Gels and Networks Date: June 17-21, 2018 Location: Prague - Czech Republic Website: www.imc.cas.cz/sympo/82pmm_png2018 Polymer Foam Date: June 19-20, 2018 Location: Pittsburgh - USA Website: www.ami.international/events/event?Code=C883 8th World Congress on Biopolymers Date: June 28-30, 2018 Location: Berlin - Germany Website: biopolymers.conferenceseries.com

April

July

Fire Retardants in Plastics Date: April 10–11, 2018 Location: Pittsburgh - USA Website: www.amiplastics.com/events/event?Code=C0881

IUPAC World Polymer Congress (Macro 2018) Date: July 1-5, 2018 Location: Cairns – Australia Website: http://www.macro18.org

March

Polímeros, 27(3), 2017 E5

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(3), 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 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(3), 2017 E7

http://dx.doi.org/10.1590/0104-1428.06016

Elastic modulus of PVDF with bentonite or LiNbO3 using deformation energy Carlos Alberto Fonzar Pintão1* and Celso Xavier Cardoso2 Department of Physics, Faculty of Science, Universidade Estadual Paulista “Julio de Mesquita Filho” – UNESP, Bauru, SP, Brazil 2 Department of Physics, Chemistry and Biology, Faculty of Science and Technology, Universidade Estadual Paulista “Julio de Mesquita Filho” – UNESP, Presidente Prudente, SP, Brazil

1

*fonzar@fc.unesp.br

Abstract Polyvinylidene fluoride (PVDF) is valued for its properties of transparency to light, lightness, flexibility, mechanical strength, chemical stability, ease of processing, and low-cost production. Ceramics have low mechanical strength and poor processability, but have excellent piezo- and pyroelectric characteristics. The deficiencies of ceramics can be minimized by combining them with polymers. Accordingly, PVDF samples with different percentages of bentonite or LiNbO3 were used to obtain composites via “casting,” and the modulus of elasticity (E) of the composites was studied using a specially designed system. The method used to obtain E took into account the strain energy and the strength of the materials. Based on the results, E decreased with an increased percentage of bentonite and, in the case of LiNbO3, for the percentages of 30% and 35% increases.

Keywords: bentonite, deformation energy, LiNbO3, modulus of elasticity, PVDF.

1. Introduction This study uses a specially designed system to determine the modulus of elasticity (E) of the tensile strength of polymer composites. Several methods and techniques can be used to obtain E, including static testing (tensile, torsion, bending), dynamic testing (resonant frequency method), wave propagation methods (ultrasonic echo-pulse method), and nanoindentation testing. Each method has advantages and disadvantages. Moreover, the measured Young’s modulus values obtained from these methods are different, even for the same sample material. Of these methods, the pulse-echo ultrasound method is most commonly used[1,2]. This technique is nondestructive and does not alter the sample’s physical or chemical properties; thus, it has significant practical interest. However, the pulse-echo ultrasound method can only measure E values for a single sample temperature and for well-defined sizes. In such cases, Poisson coefficient information is needed to obtain E, and the sample dimensions may be crucial. Clay and polymeric materials have complementary characteristics regarding the preparation of ferroelectric materials. Such materials are widely used in the electronics, sensors, and transducers industries, which require piezoelectric materials and easy processability. We used polyvinylidene fluoride (PVDF), which is an important transducer material because of its mechanical resistance, chemical inertness[3], and high piezoelectricity[4]. PVDF is used in hydrophones, industrial acoustic materials, and vibration sensors. Depending on the conditions used to process PVDF, it presents at least four crystalline phases, known as α, β, γ, and δ[5-7]. PVDF has a spherulitic crystallization morphology in which the spherulites are formed by lamellar crystalline regions, which grow from the center to the edges in a radial direction, and by amorphous regions, located between the crystalline lamellae

Polímeros, 27(3) , 183-188, 2017

of the spherulites[8,9]. In a dimethylformamide (DMF) solution from which, in the crystallization of PVDF, both the α and β phases are obtained, it has been found[10] that, depending on the crystallization temperature (T), the predominantly α-phase films crystallize at T < 160 °C. The influence of the PVDF phases is scientifically and technologically important because these phases can influence the physical and chemical properties of PVDF. In addition, PVDF is easy to process and its production costs are low[11]. Unfortunately, neat PVDF cannot completely meet the mechanical, thermal, and oxidation resistance property requirements of some harsh environments[12,13]. Many efforts have been made to improve the properties of PVDF. For example, the incorporation of organic polymers or inorganic fillers into the PVDF matrix to produce composites has been extensively studied to further improve its properties[14,15]. This present study aims to introduce and apply a technique to measure the modulus of elasticity (E) of polymeric composites.

2. Materials and Methods The casting method was used to incorporate ceramic LiNbO3 or bentonite into the PVDF, in which the PVDF grains (Florafon F4000 HD; Atochem) were dissolved in DMF under stirring and heating at 100 °C for 30 minutes in a shaker magnetic heater. The ceramics were dispersed in the DMF, and they were mixed with and dissolved in the PVDF. After stirring and heating the mixture for 10 minutes until it was homogenized, the resulting solution was poured into petri dishes and placed in an oven at 100 °C for three hours to dry and evaporate the solvent. Once all the samples were prepared, they were cut and fixed, one at a time, and

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Pintão, C. A. F., & Cardoso, C. X. measurements were taken (Figure 1A). A caliper or micrometer was used to obtain the parameters (a and b) for determining the area of the rectangular cross-section (A) and the length (L) at the instant the tensile force was applied (Figure 1B).

The values for six samples of PVDF/LiNbO3, with increasing percentages, by weight, of LiNbO3, are shown in Table 1, while the values for 10 samples of PVDF/bentonite, with increasing percentages, by weight, of the bentonite, are presented in Table 2.

2.1 Deflection energy method for obtaining E Consider an elastic structure submitted to applied loads and deformed elastically. In this deformation process, the principle of energy conservation[16] is used, expressed as: WE + Q = ∆E (1)

WE is the work carried out by the applied external forces, Q is the heat that the structure exchanges with its surrounding area, and ΔE is the variation in the associated energies of the structure: kinetic energy (K) and internal energy (U).

Considering that the increase in these loads is gradual and that a state of equilibrium is maintained during this process, then the variation of kinetic energy is zero and ΔE is due only to the variation in internal energy U. Under these conditions, Equation 1 is reduced to: WE = ∆U (2)

The work, which is energy stored in the structure due to tensile force in an infinitesimal volume element[6,7], is represented by the strain and stress tensors, σij and eij, respectively. The energy, dU, which is stored in this element when the deformation has reached its final value eij, is:

Figure 1. (A) System for measuring E: (1) Pieces of attachment of samples; (2) Force sensor (FS-PASCO: CI6537); (3) XY table; (4) Samples: PVDF/LiNbO3 or PVDF/Bentonite; (5) Rotational movement sensor (RMS-PASCO: CI6538; [8]); (6) Motor for moving XY table; (7) Key to activate motor; (8) Computer with specific software; (9) Interface (PASCO: CI7650-750); (B) Schematic diagram of the XY table, Pieces of attachment of samples and FS; Cross section of sample of area A; Prismatic bar width a, height b and length L. Table 1. Results of PVDF/LiNbO3.

LiNbO3, wt. %

B0 (N/rad)

L (mm)

a (mm)

b (mm)

A (mm2)

E (MPa)

4 10 15 25 30 35

1.480 ± 0.004 1.347 ± 0.004 0.727 ± 0.002 0.727 ± 0.002 7.34 ± 0.02 4.97 ± 0.02

18.23 ± 0.05 13.66 ± 0.05 12.67 ± 0.05 12.67 ± 0.05 8.48 ± 0.05 5.29 ± 0.05

15.45 ± 0.05 15.50 ± 0.05 13.75 ± 0.05 13.75 ± 0.05 11.6 ± 0.05 10.7 ± 0.05

0.04 ± 0.01 0.03 ± 0.01 0.04 ± 0.01 0.04 ± 0.01 0.07 ± 0.01 0.04 ± 0.01

0.6 ± 0.1 0.5 ± 0.2 0.6 ± 0.1 0.6 ± 0.1 0.8 ± 0.1 0.4 ± 0.1

1510 ± 378 1369 ± 458 579 ± 145 579 ± 145 2651 ± 379 2126 ± 532

L (mm) 12.26 ± 0.05 12.38 ± 0.05 11.96 ± 0.05 10.89 ± 0.05 11.37 ± 0.05 11.37 ± 0.05 10.28 ± 0.05 10.11 ± 0.05 10.66 ± 0.05 13.40 ± 0.05

a (mm) 8.80 ± 0.05 10.50 ± 0.05 12.05 ± 0.05 9.20 ± 0.05 12.20 ± 0.05 11.60 ± 0.05 12.20 ± 0.05 12.80 ± 0.05 13.90 ± 0.05 13.30 ± 0.05

b (mm) 0.15 ± 0.01 0.16 ± 0.01 0.15 ± 0.01 0.21 ± 0.01 0.14 ± 0.01 0.18 ± 0.01 0.20 ± 0.01 0.22 ± 0.01 0.20 ± 0.01 0.23 ± 0.01

A (mm2) 1.32 ± 0.09 1.7 ± 0.1 1.8 ± 0.1 1.93 ± 0.09 1.71 ± 0.01 2.09 ± 0.01 2.4 ± 0.1 2.8 ± 0.1 2.8 ± 0.1 3.06 ± 0.01

E (MPa) 1752 ± 118 1435 ± 91 1613 ± 108 1271 ± 61 1518 ± 109 1228 ± 69 1057 ± 53 737 ± 34 880 ± 45 439 ± 19

Table 2. Results of PVDF/bentonite. bentonite, wt. % 1 3 4 5 10 15 20 25 30 35

184

B0 (N/rad) 5.45 ± 0.01 5.62 ± 0.02 7.05 ± 0.02 6.52 ± 0.01 6.59 ± 0.02 6.52 ± 0.01 7.26 ± 0.02 5.93 ± 0.02 6.63 ± 0.04 2.90 ± 0.01

Polímeros, 27(3) , 183-188, 2017

Elastic modulus of PVDF with bentonite or LiNbO3 using deformation energy eij

∫ sij deij (3)

dU =

0

By integrating the total volume (V) of the structure, we obtain the total internal energy U due to the deflection, which is expressed as: eij

= U ∫ ( ∫ sij deij ) dV (4) 0

In cases where the elastic structure shows linear behavior, it is isotropic, and it is subjected to pure tensile stress. Using Hooke’s law[16], it is established that: = U ∫

1 2 s11dV (5) 2E

where E is the modulus of elasticity of the tensile force.

To measure E, a prismatic bar is used with a uniform cross-section of area A and length L, subjected to applied force F by a force sensor (FS) at one end with another attached to XY table (Figure 1A and B).

With regard to material resistance[17], it is known that the stress state of an internal point of a polymer, xi, based on the experimental conditions shown in Figure 1, is expressed as: F s11 = (6) A

The cross-sectional area is represented by A, which is given by the product of the sides: width (a) and height (b).

Substituting this stress component in Equation 5 yields the internal strain energy in the structure for this specific application: U=

F 2 L (7) 2 EA

The external work conducted by F is: WE= F ∆L (8)

From Equation 7 and Equation 8, we obtain: F=

2E A ∆L = B* ∆L (9) L

If we replace the values of the cross-sectional area (A) and use B* and the calibration factor of the force sensor (f), we obtain the equation for the calculation of E: E=

B* L f (10) 2A

2.2 Measurement of the modulus of elasticity of tensile force A rotation movement sensor (RMS) was used to measure the turning angle, φ (rad), of the screw that moves the XY table when it shifts from ΔL[18]. A pulley, with a diameter of ϕ=28.70 ± 0.05 mm, was attached to the RMS shaft.

With the belt that passes through this pulley, and in the other pulley of the same diameter ϕ affixed to the shaft of the table Polímeros, 27(3) , 183-188, 2017

transmission mechanism, one can register φ. On the same side as the rotation sensor, we used a microwave motor with a constant torque to move the XY table in the two required directions: one for approximation and the other for drawing XY table away from the force sensor. One piece of the XY table is attached to the opposite side of the rotation sensors, and screws are used to fix one side of the sample. The force sensor (FS) has a fixture that is similar to that of the table.

That fixture is attached with a screw to the other end of the sample (Figure 1A and B). The FS is fixed to a steel base that sits on a countertop. Turning the screw moves the XY table at an angle φ (rad), displacing the sample. The motion is measured by ΔL, i.e., φ (0.000160 m).

Thus, tensile force is slowly applied to the sample by force F. FS is zeroed before measuring F. The value of F is automatically recorded for each angle φ with an RMS. The two sensors are connected at an interface to a computer.

Using PASCO software, the experimental points for F (N) and φ (rad) are obtained in real-time, simultaneously. Thus, we can adjust the straight line and obtain the slope of B0 because the points are linear in a regime in which the material shows elastic behavior. The slope B0 has units expressed in N/rad, which is transformed to N/m using the B*=B0 (1/0.000160) ratio, since it was evaluated for the measurement system in question where 1 rad equals 0.000160 m.

To determine these measurements, we had to obtain a calibration factor (f) because the value of F measured (FMEASURED) in FS is different from that of F applied (FAPPLIED). Factor f can be obtained experimentally by applying known forces (FAPPLIED), and then measuring the forces with FS (FMEASURED). For this, we used known masses: 0.020 kg, 0.050 kg, 0.100 kg, 0.200 kg, 0.500 kg, and 1.00 kg. They were fixed at one end of an inextensible cord passing over a pulley at the other end, and attached to FS. Before placing the samples of polymer/LiNbO3 or polymer/bentonite into clamps, their dimensions (a and b) were measured with a caliper and a micrometer.

At the beginning of the experiment, the samples were fixed on the XY table and FS by locking the components without an applied tensile force. L0 was measured with a caliper. The length, L, was obtained from the curve of force F (N) as a function of the angle φ (rad). For each of the samples studied, three measures of F as a function of the angle φ were obtained. The first measurement was very important for providing information about how many rotations (φ [ rad ]) were needed to turn the screw of the XY table to exert tensile force on the sample. L is reported as the sum of (L0 + φ [0.00016 m]). For the second and third measures, the screw was turned counterclockwise to apply new tensile force again, taking care to drive the XY table into a position where the sample would not be submitted to initial tensile stress. The value of L, which was the same as that measured in the first experimental curve (F as a function of φ), was used in Equation 10. At the beginning of every measurement, the force sensor was reset.

185

Pintão, C. A. F., & Cardoso, C. X.

3. Results and Discussion The results presented in Figure 2 show the calibration curve of the FS. The value obtained was f = 11.0697 ± 0.0003.

This value was used in Equation 10 to obtain E. A typical curve, F, as a function of φ for a PVDF/LiNbO3 (65/35 wt. %) composites, is shown in Figure 3. The slope B0 of the three curves shown in Figure 3 was obtained by fitting a straight line to each of the curves when they showed the linear behavior that is characteristic of an elastic regime.

For the proposed method to be applied, it is important that a region on the curve of force F (N) as a function of the angle φ (rad) represents linear behavior. In our experimental procedure, the limits of elasticity of these samples were not exceeded. Thus, we achieved good reproducibility for the E value. On the other hand, a different procedure can be adopted that uses various samples cut from the same sample source. Then, the complete experimental curve (F as a function of φ) can be determined for each sample.

In that case, we can determine E using the strain energy method. It is possible to obtain the values of the forces for the limit of elasticity and for the rupture of the sample.

Toward that end, we must consider the maximum load limits of FS (± 20N). Because we had a number of reduced samples, we chose the method described in this paper.

An examination of the samples under a microscope shows the presence of isotropic and anisotropic structures, which could be due to the method in which these compounds were prepared[5]. This may explain the greater or lesser rigidity of the samples when adding increasing mass percentages of LiNbO3. The procedure to obtain E shows that these samples behaved as an elastic structure, but that does not mean that they were completely isotropic, and it does not invalidate the proposed method. As the crystallization temperature is decreased to < 160 °C, it is expected that there would be a predominance of the α phase[8-10]. For each PVDF sample with different percentages of LiNbO3, we calculated the average value of B0, measured by the sensors FS and RMS, relative to the three measurements of the curves as a function of the angular position. Equation 10 was then used to calculate the value of E. The results are shown in Table 1.

Figure 2. Calibration curve of the force sensor. The calibration factor f is 11.0697 ± 0.0003.

Figure 3. Typical curves of force F (N) as a function of the angular position φ (rad). Sample of PVDF/LiNbO 3 : L0 = 4.65 ± 0.05 mm; L = 4.65 + 4.0 (0.16) = 5.29 mm ± 0.05 mm; a = 10.70 ± 0.05 mm; b = 0.04 ± 0.01 mm; A = 0.4 ± 0.1 mm2; B0 = 4.97 ± 0.02 N/rad; B*= 31083 ± 148 N/m. E is calculated using Equation 10: E = 2126 ± 532 MPa.

Figure 4 shows the final results of E, which is expressed in MPa for all the studied samples with increasing percentages of LiNbO3. Using our method, the E value of the PVDF was close to 2000 MPa, which is close to the expected value[19].

A typical curve, F, as a function of φ for a composite based on PVDF/bentonite (90/10 wt. %), is shown in Figure 5.

The same procedure for PVDF/LiNbO3 was performed for PVDF/bentonite. For each PVDF sample with different percentages of bentonite, we calculated the average value of B0, measured by the sensors FS and RMS, relative to the three measurements of the curves as a function of the angular position. Equation 10 was then used to calculate the value of E. The results are shown in Table 2.

Figure 6 shows the final results of E, which is expressed in MPa for all the studied samples with increasing percentages of bentonite. In this case, the rigidity of the composite decreases with the increasing percentage, by weight, of bentonite.

186

Figure 4. Experimental values of E obtained in the samples (PVDF/LiNbO3) at different percentages by weight. The value of E for the sample PVDF/LINbO3 (100/0, wt. %) was obtained according to Wallner et al.[19], that is, 1771 ± 46 MPa. Polímeros, 27(3) , 183-188, 2017

Elastic modulus of PVDF with bentonite or LiNbO3 using deformation energy

6. References

Figure 5. Typical curves of force F (N) as a function of the angular position φ (rad). Sample of PVDF with bentonite: L0 = 11.00 ± 0.05 mm; L = 11.00 +2.3 (0.16) = 11.37 mm ± 0.05 mm; a = 12.20 ± 0.05 mm; b = 0.14 ± 0.01 mm; A = 1.71 ± 0.01 mm2; B0 = 6.59 ± 0.02 N/rad; B*= 41188 ± 119 N/m. E is calculated by Equation 10: E = 1518 ± 109 MPa.

Figure 6. Experimental values of E obtained from the samples (PVDF/bentonite) at different percentages by weight. The value of E for the sample PVDF/bentonite (100/0, wt. %) was obtained according to Wallner et al.[19], that is, 1771 ± 46 MPa.

4. Conclusions The system described in this study is an alternative method for obtaining E in polymer composites. This method took into account the tensile strain energy and the strength of the materials. The method was applied to PVDF samples with different percentages, by weight, of LiNbO3 or bentonite.

It can be used with other types of materials in cases where the elastic structure shows linear behavior. Based on our results, the studied samples showed variations in the values​​ of E when the percentage of LiNbO3 or bentonite changed.

The values of E decreased as the percentage of bentonite increased. The value of E increased when the percentages, by weight, for LiNbO3 were 30% and 35%.

5. Acknowledgements The authors thank FAPESP, process number: 2007/04094-9, and CAPES, process number: BEX 6571/14-0.

Polímeros, 27(3) , 183-188, 2017

1. Truel, R., Elbaum, C., & Chic, B. B. (1969). Ultrasonic methods in solid state physics. New York: Academic Press.

2. Nowick, A. S., & Berry, B. S. (1972). Anelastic relaxation in crystalline solids. New York: Academic Press.

3. Roh, Y., Varadan, V. V., & Varadan, V. K. (2002). Characterization of all the elastic, dielectric, and piezoelectric constants of uniaxially oriented poled PVDF films. IEEE Transactions on Ultrasonics, Ferroelectrics, and Frequency Control, 49(6), 836-847. PMid:12075977. http://dx.doi.org/10.1109/ TUFFC.2002.1009344.

4. Motyl, E. (2001). Comparison between step and pulsed electroacoustic techniques using both PVDF and LiNbO3 transducers. Journal of Electrostatics, 51-52, 530-537. http:// dx.doi.org/10.1016/S0304-3886(01)00064-X. 5. Costa, C. M., Mendes, S. F., Sencadas, V., Ferreira, A., Gregorio, R. Jr, Ribelles, J. L. G., & Méndez, S. L. (2010). Influence of processing parameters on the polymer phase, microstructure and macroscopic properties of poly (vinilidene fluoride)/ Pb(Zr0.53Ti0.47)O3 composites. Journal of Non-Crystalline Solids, 356(41-42), 2127-2133. http://dx.doi.org/10.1016/j. jnoncrysol.2010.07.037.

6. Mendes, S. F., Costa, C. M., Caparros, C. V., Sencadas, V. S., & Lanceros-Méndez, S. (2012). Effect of filler size and concentration on the structure and properties of poly(vinylidene fluoride)/BaTiO3 nanocomposites. Journal of Materials Science, 47(3), 1378-1388. http://dx.doi.org/10.1007/s10853-011-59167.

7. Mendes, S. F., Costa, C. M., Sencadas, V., Nunes, J. S., Costa, P., Gregorio, R., Jr., & Méndez, S. L. (2009). Effect of the ceramic grain size and concentration on the dynamical mechanical and dielectic behavior of poly (vinilidene fluoride)/Pb(Zr0.53Ti0.47)O3 composites. Applied Physics A: Materials Science & Processing, 96(4), 899-908. http://dx.doi.org/10.1007/s00339-009-5141-2.

8. Lovinger, A. J. (1982). Developments in crystalline polymers.

Journal of Polymer Science Part C: Polymer Letters, 20(10), 559-560. http://dx.doi.org/10.1002/pol.1982.130201011.

9. Broadhurst, M. G., Davis, G. T., Mckinney, J. E., & Collins, R. E. (1978). Piezoelectricity and pyroelectricity in polyvinylidene fluoride: a model. Journal of Applied Physics, 49(10), 49924997. http://dx.doi.org/10.1063/1.324445.

10. Gregorio, R., Jr., & Cestari, M. (1994). Effect of crystallization temperature on the crystalline phase content and morphology of poly(vinylidene fluoride). Journal of Polymer Science. Part B, Polymer Physics, 32(5), 859-870. http://dx.doi.org/10.1002/ polb.1994.090320509.

11. Inderherbergh, J. (1990). Polyvinylidene fluoride (PVDF) appearance, general properties and processing. Ferroelectrics, 115(4), 295-302. http://dx.doi.org/10.1080/00150193.1991.1 1876614.

12. Dargaville, T. R., Celina, M., & Chaplya, P. M. (2005). Evaluation of piezoelectric Poly (vinylidene fluoride) polymers for use in space environments. I. Temperature limitations. Journal of Polymer Science. Part B, Polymer Physics, 43(11), 1310-1320. http://dx.doi.org/10.1002/polb.20436.

13. Dargaville, T. R., Celina, M., Martin, J. W., & Banks, B. A. (2005). Evaluation of piezoelectric PVDF polymers for use in space environments. II. Effects of atomic oxygen and vacuum UV exposure. Journal of Polymer Science. Part B, Polymer Physics, 43(18), 2503-2513. http://dx.doi.org/10.1002/ polb.20549.

14. Malmonge, L. F., Langiano, S. C., Cordeiro, J. M. M., Mattoso, L. H. C., & Malmonge, J. A. (2010). Thermal and Mechanical Properties of PVDF/PANI Blends. Materials Research, 13(4), 465-470. http://dx.doi.org/10.1590/S1516-14392010000400007.

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Pintão, C. A. F., & Cardoso, C. X. 15. Xu, H. P., Dang, Z. M., Jiang, M. J., Yao, S. H., & Bai, J. (2008). Enhanced dielectric properties and positive temperature coefficient effect in the binary polymer composites with surface modified carbon black. Journal of Materials Chemistry, 18(2), 229-234. http://dx.doi.org/10.1039/B713857A. 16. Tauchert, T. R. (1974). Energy principles in structural mechanics.

New York: McGraw-Hill.

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3rd ed. Rio de Janeiro: Guanabara Dois.

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Fracture analysis of annealed PVDF films. Polymer Testing, 27(3), 392-402. http://dx.doi.org/10.1016/j.polymertesting.2008.01.006.

Received: May 31, 2016 Revised: Oct. 20, 2016 Accepted: Nov. 10, 2016

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http://dx.doi.org/10.1590/0104-1428.07216

Sulfonation degree effect on ion-conducting SPEEK-titanium oxide membranes properties Jacqueline Costa Marrero1*, Ailton de Souza Gomes1, Wang Shu Hui2, José Carlos Dutra Filho1 and Vivianna Silva de Oliveira3 Instituto de Macromoléculas Professora Eloisa Mano – IMA, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brazil 2 Departamento de Engenharia Metalúrgica e de Materiais, Escola Politécnica, Universidade de São Paulo – USP, São Paulo, SP, Brazil 3 Escola Técnica Rezende Rammel – ETRR, Rio de Janeiro, RJ, Brazil 1

*jacquecosta@gmail.com

Abstract Polymeric membranes were developed using a SPEEK (sulfonated poly(ether ether ketone)) polymer matrix, containing titanium oxide (TiO2) (incorporated by sol-gel method). SPEEK with different sulfonation degrees (SD): 63% and 50% were used. The influence of sulfonation degree on membrane properties was investigated. The thermal analysis (TGA and DTGA) and X-ray diffraction (XRD) were carried out to characterize the membranes and electrochemical impedance spectroscopy (EIS) was carried out to evaluate the proton conductivity of the membranes. The proton conductivities in water were of 3.25 to 37.08 mS.cm-1. Experimental data of impedance spectroscopy were analyzed with equivalent circuits using the Zview software, and the results showed that, the best fitted was at 80 °C. Keywords: membrane, sulfonation degree, SPEEK.

1. Introduction Polymer electrolyte membrane fuel cells (PEMFCs) are considered as the most promising fuel cell technology for a wide range of applications due to the stable operation, the high energy generation yield and the simplicity of the system[1].

The polymer electrolyte membranes (PEMs) has gain considerable attention due to its applications in many energy related fields such as water electrolyzers, fuel cells, battery systems, etc. The function of the membrane in PEM fuel cells is two fold: proton conduction from the anode to the cathode, and effective separation of the anode and cathode gases[2]. PEMs exhibit several advantages over liquid or solid inorganic electrolytes. The polymeric membrane should satisfy several requirements: high proton conductivity, good chemical, thermal and mechanical stability, and low permeability to reactants. Low cost and ready availability are important economical requirements. Furthermore, the membrane should work at an operative temperature around 120 °C for long time[1,3].

The poly(ether ether ketone) (PEEK) is one polymer with good characteristics: excellent chemical resistance, high thermo-oxidative stability and low cost[4]. Sulfonated

poly(ether ether ketone is obtained by sulfonation of PEEK by concentrated sulfuric acid. The SPEEK offers the advantages of lower cost, easier preparation, controllable proton conductivity and excellent chemical and thermal stability[5]. The chemical structure of PEEK and SPEEK are shown in Figure 1.

Sulfonation is a powerful method for the modification of polymers to introduce sulfonic acid functionality, which aids in the transport of protons. It makes the polymer hydrophilic and hence increases the solubility of polymers[7].

Higher sulfonation degree can enhance the density of acid sites and facilitate the proton transport. Therefore, there exists a conflict between high SD and poor mechanical strength of SPEEK membrane[8]. In order to enhance its competitive superiority in PEMFC application, varied efforts were made to modify the structure-property relationship of SPEEK. Incorporation of nanosized binary oxide materials (SiO2, TiO2, ZrO2) in SPEEK membranes has several attributes of interest, including decreased membrane swelling, reduced permeability towards methanol and improved morphological stability without compromising proton conductivity at high degree of sulfonation[9].

Figure 1. Chemical structures of PEEK and SPEEK[6].

Polímeros, 27(3) , 189-194, 2017

189

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

Marrero, J. C., Gomes, A. S., Hui, W. S., Dutra, J. C., Fo., & Oliveira, V. S. In this study, composite membranes using SPEEK (with different sulfonation degrees (63 and 50%)) and titanium oxide were prepared and characterized by TGA and DTGA, XRD and EIS to evaluating the influence of sulfonation degree in the structure, thermal properties and ionic conduction of the membranes.

2. Materials and Methods 2.1 Sulfonation of PEEK 10 g of PEEK (VictrexTM USA Inc) was dissolved in 200 ml of sulfuric acid (H2SO4) (VETEC Química Fina LTDA) under nitrogen atmosphere. The mixture was stirred at 30 °C for 4 h. To follow the reaction, the mixture was stirred at 50 °C for 5 h. The sulfonation reaction was terminated by precipitating the acidic polymer solution into a large excess of ice-cold water under continuous mechanical stirring. The polymer precipitate was washed several times with deionized water until the pH was neutral. The polymer was dried under vacuum for 24 h at 70 °C. The same procedure was done with the difference that the mixture was stirred at 30 °C for 4 h and after that the mixture was stirred at 50 °C for 4 h to obtain different sulfonation degree.

2.2 Determination of the ion-exchange capacity (IEC) and SD The IEC and SD of SPEEK was determined by titration method with NaOH solution and phenolphthalein, using the following equations: IEC =

SD =

Volume of NaOH x molar concentration of NaOH Weight of SPEEK

M p x IEC 1000 − ( M SO3H x IEC )

(1)

(2)

where MP = 288 g mol and MSO3H = 80 g mol are the molecular weights of the monomer unit of the polymer and of a sulfonic group, respectively.

The measured IEC average value (after three measurements) for SPEEK was: 1.87 meq g−1 corresponding to an SD of 63% (SPEEK 63%) and 1.52 meq g−1 corresponding to an SD of 50% (SPEEK 50%).

-1

-1

stored in a vacuum oven for 72 h at 80 °C. The membranes from SPEEK 50% were prepared by the same technique described.

2.4 X-ray diffraction (XRD) XRD patterns of dry membranes were recorded in the 2θ range 2 – 80° at a scan rate of 0.02 2θ/s with Rigaku Miniflex II X-ray diffractometer, using CuKα (λ =1.54Å) radiation.

2.5 Thermal analysis (TGA and DTGA) Thermal gravimetric studies (TGA/DTGA) were performed in a TA TGA Q500 in the temperature range between 30 and 700 °C at a heating rate of 10 °C/min and carried out under nitrogen flow 60 mL/min.

2.6 Electrochemical impedance spectroscopy (EIS) The proton conductivity of membranes was measured using electrochemical impedance spectroscopy with an Autolab PGSTAT-30 potentiostat/galvanostat instrument, in the frequency range of 1 MHz to 10 Hz and amplitude of 5 mV, using two stainless steel blocking electrodes and the contact area of 0.1 cm2. The measurements were done in the normal direction to the plane of the membranes, with samples immersed in deionized water at different temperatures (30, 50 and 80 °C).

The resistance value associated to membrane conductivity was determined from the high frequency intercept of the impedance with the real axis (Z’) and was called RS(exp.).

The proton conductivity (σ) was calculated from the impedance data according to Equation 3: s=

L (3) RS

where, σ is the proton conductivity (S.cm–1), L is the membrane thickness (cm), R is the measured membrane resistivity (Ohm) and S is the electrode area perpendicular to current flow (cm2).

2.3 Membranes preparation

Experimental data have been fitted with the Randles equivalent circuit (Figure 2) using the Zview software.

In this software an equivalent circuit is assembled which generates a theoretical curve, adjusted to experimental measurements, thus, we obtain the values of the circuit elements. A Randles circuit is an equivalent electrical circuit

Membranes were prepared using the sol-gel process.

For membrane preparation, the SPEEK 63% polymer was dissolved in NMP (1-Methyl-2-pyrrolidone) (VETEC Química Fina LTDA). After complete polymer dissolution, the solution was placed to cool. After that, acetylacetone (ACAC) (Aldrich Chemical Co) and titanium oxide (4 wt%) were added to the polymer solution and was stirred for 1 h. The amount of titanium alkoxide was calculated to produced samples containing 96:4 wt% SPEEK/TiO2 rations using Titanium tetrabutoxide ((Ti(C4H9O)4) or (Ti(OBu)4)) (MERCK-Schuchardt). ACAC was used as a chelating agent to avoid precipitation of the inorganic compound.

The mixtures were cast on glass plates heated to 70 °C for solvent evaporation. After casting, the membranes were

Figure 2. The Randles equivalent circuit [10]. The Randles parameters are RS (active electrolyte resistance); CDL (double-layer capacitance); RCT (active charge transfer resistance) and ZW (Warburg impedance).

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Polímeros, 27(3) , 189-194, 2017

Sulfonation degree effect on ion-conducting SPEEK-titanium oxide membranes properties that consists of an active electrolyte resistance RS in series with the parallel combination of the double-layer capacitance CDL and an impedance of a faradic reaction. The impedance of a faradic reaction consists of an active charge transfer resistance RCT and a specific electrochemical element of diffusion W, where ZW is the Warburg impedance (the first distributed element (DE) introduced into electrochemistry)[10].

The electrolyte resistance, the double-layer capacitance, the charge transfer resistance and the Warburg impedance were obtained by fitting the impedance spectra using the equivalent circuit and software Zview. The electrolyte resistance obtained by fitting the experimental data was called RS(sim.) and was compared with RS(exp.). When use the Circuit of Figure 2 and the software Zview to simulate the experimental impedance spectrum, appears ZW(1-R), ZW(1-T) and ZW(1-P), where 1 means the distributed element type (of 11 available) and R, T, P are free parameters which have different meanings for each DE type; in general they use the following convention: R - Usually resistance, T - A time constant or capacitance and P - an exponent[11].

3. Results and Discussions X-ray diffraction patterns of membranes are presented in Figure 3. The membranes are fully amorphous: a broad diffraction peak at 2θ around of 20° is indicative of the

lack of crystallinity. The amorphous nature of membranes indicates SPEEK with medium and high SD, in this case 63 and 50%[7]. The introduction of SO3H groups into the PEEK alters the chain conformation and packing, and thus causes loss of crystallinity. It is know that the amorphous nature of membranes imparts easier mobility of polymer chains which accounts for faster proton exchange resulting in higher proton conductivity of membranes.

The results of thermogravimetric analysis and the differential thermogravimetric analysis of prepared membranes are shown in Figure 4. The first weight loss around 100 °C can be attributed to water molecules adsorbed by hydrophilic groups and lost until the dry state of the sample is reached[9]. The second weight loss between 100 °C and 400 °C is attributed to loss of residual solvent NMP and the decomposition of the sulfonic acid groups of SPEEK. The decomposition of the polymer matrix was from 400 °C to 700 °C. The exothermic peaks, which corresponds to the water loss and the decomposition of the sulfonic groups, are more intense for SPEEK 63%- TiO2, with higher SD, because the water amount is associated to the sulfonic groups, the peak intensity increases with the increase of sulfonation degree[12].

AC impedance spectroscopy was performed to determine the conductivity of these membranes. The variation of conductivity with temperature for the different composite

Figure 3. XRD patterns of membranes: (A) SPEEK 63%-TiO2 and (B) SPEEK 50%- TiO2.

Figure 4. Thermal analysis (TGA and DTGA) of SPEEK 63%- TiO2 and SPEEK 50%- TiO2 membranes.

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191

Marrero, J. C., Gomes, A. S., Hui, W. S., Dutra, J. C., Fo., & Oliveira, V. S. membranes is shown in Figure 5. The Arrhenius plots of the temperature dependency of the conductivity exhibit that the conductivity increases with temperature. Proton conductivity data for studied membranes are presented in Table 1. In the hydrated composite membranes the proton conduction is likely to happen by vehicular mechanism, rather than by Grotthuss mechanism, resulting in a long-range conductivity, with protons percolating through the sample[2].

Di Vona et al.[3] have reported the preparation of hybrid membranes based on highly sulfonated poly(ether ether ketone) (SPEEK, SD = 0.9) where titanium oxide network was dispersed by in situ sol-gel reactions. They obtained membranes with enhanced thermal stability, reduced water uptake and good proton conductivity (σ = 58.00 mS.cm-1) up to 120 °C and concluded that the samples were suitable for application as polymeric electrolytes at intermediate temperature. Dutra et al.[13] also reported the proton-conducting of hybrid membranes consisting of SPEEK and titanium oxide, the proton conductivity in ethanol solution was of the order of 10-3 S.cm-1 when 4 or 8 wt% TiO2 were added, and generally increased with addition of TiO2.

In this study, the membranes showed conductivity values in the range 3.25-37.08 mS.cm-1. A maximum conductivity of 37.08 mS.cm-1 was obtained for the SPEEK 63%- TiO2 membranes at 80 °C. For all studied temperatures, the SPEEK 63%- TiO2 membranes with SD = 63% showed higher conductivity than the SPEEK 50%- TiO2 membranes with SD = 50%, as expected. Proton transfer enhances by increasing the number of acid sites enhances the proton transfer.

The impedance characteristics of the SPEEK 63%- TiO2 and SPEEK 50%- TiO2 membranes under different operating temperatures (30, 50 and 80 °C) were also investigated.

The result was plotted as a Cole-Cole plot to show the real/imaginary parts of the impedance at various frequencies.

A typical Nyquist plot of a composite membrane sandwiched stainless steel electrodes is shown in Figure 6 for SPEEK 50%- TiO2 membranes (similar result were obtained for SPEEK 63%- TiO2 membranes). The profile shows that the impedance decreases with increasing frequency.

It indicates that the interfacial impedance decreases with increasing frequency, which can be attributed to double layer formation and charge transfer reaction.

Figure 5. Arrhenius plots for the proton conductivity of the membranes as a function of temperature.

Table 1. Proton conductivity values of the membranes.

T (°C) 30 50 80

192

Conductivity (mS.cm-1) SPEEK 63%-TiO2 SPEEK 50%-TiO2 15.55 23.42 37.08

3.25 4.68 13.14

Figure 6. Nyquist plots of the SPEEK 50%- TiO2 membrane for experimental impedance spectrum and simulated impedance spectrum (using the Randles equivalent circuit and the Zview software) for: (A) T = 30 °C; (B) T = 50 °C; (C) T = 80 °C. Polímeros, 27(3) , 189-194, 2017

Sulfonation degree effect on ion-conducting SPEEK-titanium oxide membranes properties Table 2. RS(exp.) obtained from the experimental impedance spectrum and circuit parameters used to simulated the impedance spectrum.

The Randles parameters are RS(exp.) (active electrolyte resistance experimental), RS(sim.) (electrolyte resistance obtained by fitting the experimental), CDL (double-layer capacitance), RCT (active charge transfer resistance) and ZW (Warburg impedance), where ZW(1-R), ZW(1-T) and ZW(1-P) and 1 means the distributed element type (of 11 available) and R, T, P are free parameters which have different meanings for each distributed elements type; in general they use the following convention: R - Usually resistance, T - A time constant or capacitance and P - an exponent.

SPEEK 63%-TiO2 T (°C)

SPEEK 50%-TiO2

30

50

80

30

50

80

RS(exp.) (Ω)

1.508

1.121

0.845

4.433

2.504

1.122

RS(sim.) (Ω)

1.281

1.127

0.830

4.618

2.594

1.147

CDL(µF)

1.134

2.137

8.096

1.145

4.529

5.574

RCT(Ω)

0.453

0.519

0.032

1.464

0.899

0.427

ZW(1-R) (Ω)

25.99

23.86

2.708

1770

7050

3119

ZW(1-T) (Ω)

0.061

0.063

0.075

0.416

3.598

18.28

ZW(1-P) (Ω)

0.600

0.601

0.311

0.763

0.761

0.709

A simulated impedance plot for SPEEK 50%- TiO2 membranes for 30 °C, 50 °C and 80 °C is shown in Figure 6 using circuit parameters estimated from the experimental impedance plot; these circuit parameters are given in Table 2.

For all membranes the electrolyte resistances decreased with increasing temperature and with SD. RS(exp.) and RS(sim.) values were too close indicating a good fit, the highest coincidence occurs for higher temperatures. The SPEEK 63%- TiO2 membrane showed lower resistivity than the SPEEK 50%- TiO2 membrane, resulting in higher conductivity.

The level of agreement between experiment and simulation is quite satisfactory in terms of shape and distribution of frequencies, supporting the view that circuit equivalents are a meaningful way of representing the data. The results showed the best fitted were at 80 °C (Figure 6).

4. Conclusions Membranes based on SPEEK (with SD = 63% and 50%) and titanium oxide have been synthesized using the sol-gel process. The influence of sulfonation degree of the polymer on membrane properties was investigated. The ion exchange capacity increased with increasing of SD. The amorphous nature of composite membranes was confirmed by XRD studies. Water is coordinated to the sulfonic groups. The peak intensity which corresponds to the water loss and the loss of sulfonic groups increases with the increase of SD. For all membranes, the Rexp. and Rsim. values decreased with increasing of temperature and were too close indicating a good fit. The proton conductivity increased with increase of temperature and improves with increasing of sulfonation degree. The highest proton conductivity of 37.08 mS.cm-1 in water was obtained for the SPEEK 63%-TiO2 membrane at 80 °C. Although still not reaching very high conductivity values, the SPEEK 63%-TiO2 membranes shows the best performance in terms of thermal properties and proton conductivity, for this reason, we believe that this particular membranes have good prospects to be tested as electrolyte in PEMFC. The present investigation suggests measuring the conductivity of these composite membranes in ethanol and for higher temperatures.

Polímeros, 27(3) , 189-194, 2017

5. Acknowledgements The authors thank PNPD/CAPES for financial support granted to carry out this work.

6. References 1. Nikolic, V. M., Krkljes, A., Popovic, Z. K., Lausevic, Z. V., & Miljanic, S. S. (2007). On the use of gamma irradiation crosslinked PVA membranes in hydrogen fuel cells.

Electrochemistry Communications, 11(9), 2661-2665. http:// dx.doi.org/10.1016/j.elecom.2007.08.022.

2. Anis, A., Banthia, A. K., & Bandyopadhyay, S. (2008). Synthesis & characterization of PVA/STA composite polymer electrolyte membranes for fuel cell application. Journal of Materials Engineering and Performance, 17(5), 772-779. http://dx.doi. org/10.1007/s11665-008-9200-1.

3. Di Vona, M. L., Ahmed, Z., Bellitto, S., Lenci, A., Traversa, E., & Licoccia, S. (2007). SPEEK-TiO2 nanocomposite hybrid proton conductive membranes via in situ mixed sol-gel process.

Journal of Membrane Science, 296(1-2), 156-161. http://dx.doi. org/10.1016/j.memsci.2007.03.037.

4. Kawaguti, C. A., Dahmouche, K., & Gomes, A. de S. (2012).

Nanostructure and properties of proton-conducting sulfonated poly(ether ether ketone) (SPEEK) and zirconia-SPEEK hybrid membranes for direct alcohol fuel cells: effect of the nature of swelling solvent and incorporation of heteropolyacid. Polymer International, 61(1), 82-92. http://dx.doi.org/10.1002/pi.3151.

5. Yang, T., Xu, Q., Wang, Y., Lu, B., & Zhang, P. (2008). Primary study on double-layer membranes for direct methanol fuel cell.

International Journal of Hydrogen Energy, 33(22), 6766-6771. http://dx.doi.org/10.1016/j.ijhydene.2008.08.011.

6. Kobayashi, T., Rikukawa, M., Sanui, K., & Ogata, N. (1998).

Proton Conducting Polymers Derived from Poly (ether-ether ketone) and Poly (4-phenoxybenzoyl-1,4 phenylene). Solid State Ionics, 106(3-4), 219-225. http://dx.doi.org/10.1016/ S0167-2738(97)00512-2.

7. Zaidi, S. M. J. (2003). Polymer sulfonation: a versatile route to prepare proton-conducting membrane material for advanced technologies. Arabian Journal for Science and Engineering, 28(2B), 183-194.

8. Hou, H., Polini, R., Di Vona, M. L., Liu, X., Sgreccia, E., Chailan, J.-F., & Knauth, P. (2013). Thermal crosslinked and nanodiamond reinforced SPEEK composite membrane for PEMFC. International Journal of Hydrogen Energy, 38(8), 3346-3351. http://dx.doi.org/10.1016/j.ijhydene.2012.12.019.

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Marrero, J. C., Gomes, A. S., Hui, W. S., Dutra, J. C., Fo., & Oliveira, V. S. 9. Di Vona, M. L., Sgreccia, E., Donnadio, A., Casciola, M., Chailan, J. F., Auer, G., & Knauth, P. (2011). Composite polymer electrolytes of sulfonated poly-ether-ether-ketone (SPEEK) with organically functionalized TiO2. Journal of Membrane Science, 369(1-2), 536-544. http://dx.doi.org/10.1016/j. memsci.2010.12.044.

10. Macdonald, J. R. (1987). Impedance spectroscopy-Emphasizing solid materials and systems. New York: Wiley-Interscience.

11. Macdonald, J. R., & Potter, L. D. Jr (1987). A flexible procedure for analyzing impedance spectroscopy results: Description and illustrations. Solid State Ionics, 23(1), 61-79. http://dx.doi. org/10.1016/0167-2738(87)90068-3.

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12. Carbone, A., Pedicini, R., Portale, G., Longo, A., D’Ilario, L., & Passalacqua, E. (2006). Sulphonated poly(ether ether ketone) membranes for fuel cell aplication: thermal and structural characterization. Journal of Power Sources, 163(1), 18-26. http://dx.doi.org/10.1016/j.jpowsour.2005.12.066.

13. Dutra, J. C., Fo., Santos, T. R., & Gomes, A. S. (2014).

Nanostructured Polyelectrolytes Based on SPEEK/TiO2 for Direct Ethanol Fuel Cells (DEFCs). Polímeros: Ciência e Tecnologia, 24, 43-48. http://dx.doi.org/10.4322/polimeros.2014.069.

Received: June 06, 2016 Revised: Dec. 01, 2016 Accepted: Jan. 05, 2017

Polímeros, 27(3) , 189-194, 2017

http://dx.doi.org/10.1590/0104-1428.12516

Effect of solvents on the morphology of PMMA films fabricated by spin-coating Giovana da Silva Padilha1*, Virginia Mansanares Giacon2 and Julio Roberto Bartoli3 Centro de Pesquisa em Manufatura de Materiais Avançados, Faculdade de Ciências Aplicadas – FCA, Universidade Estadual de Campinas – UNICAMP, Limeira, SP, Brazil 2 Departamento de Engenharia de Materiais, Universidade Federal do Amazonas – UFAM, Manaus, AM, Brazil 3 Departamento de Engenharia de Materiais e de Bioprocessos, Faculdade de Engenharia Química – FEQ, Universidade Estadual de Campinas – UNICAMP, Campinas, SP, Brazil 1

*giovana.padilha@fca.unicamp.br

Abstract A method to produce thin layer of PMMA (poly (methyl methacrylate) films by spin coating is reported. PMMA is dissolved in xylene, a mixture of MIBK-xylene and chloroform. Varying the PMMA concentration and spin-coating speeds controls the thickness of the PMMA films. Using chloroform as solvent evidenced thickness around 10 μm.

This thickness is suitable for core of optical polymeric films. These PMMA films with lower roughness are treated with CHF3 plasma to produce the cladding. The morphology of PMMA films is analyzed by atomic force microscopy and scanning electron microscopy. Profilometry, metricon and contact angle analysis are associated with the effective parameters in roughness and their effect before and after plasma treatment. Gel permeation chromatography (GPC) is used for estimating molecular weights of PMMA. Keywords: PMMA, spin-coating, solvent effect, morphology, CHF3 plasma.

1. Introduction Poly (methyl methacrylate) (PMMA) is the most important thermoplastic polymer that has been widely used in various industrial fields due to its transparency, light weight and low cost[1]. It can substitute glass in several hard-use products such as antiriot vehicles, and it can be used in different technological applications, such as optical devices, optical lenses, and in gratings or waveguides[2-4].

In the field of optics, PMMA has been studied due to the intrinsic versatility of its molecular structure, which allows for enhanced refractive index modeling for core material, and also for the ease of manufacturing process and patterning capability[5,6]. In the preparation of PMMA films, spin-coating can be used mainly because it allows fast and reproducible large scale production. Basically, the polymer solution is deposited on a static substrate, which is rotated at a given angular velocity during a given period of time (Figure 1).

The thickness of the resulting films depends on the concentration of the polymer in solution, polymer molecular weight, spinning velocity, and spinning time[8-11]. Solvents have been used to dissolve PMMA to obtain solutions.

Hamdy et al.[4] dissolved PMMA in chloroform to produce transparent blue thin films/sheets using methylene blue dye (MB) as a dopant in PMMA for UV sensors.

In Liu et al.[12] PMMA was dissolved in a binary solvent of N,N‑dimethylacetamide and acetone to investigate the effect of solvents on solution properties and electrospun PMMA membrane morphology. Using different solvents, different PMMA morphologies and thicknesses can be obtained, and

Polímeros, 27(3) , 195-200, 2017

the type of solvent selected plays an important role in the uniformity and thickness of the PMMA films. A solvent with lower vaporizability results in better uniformity and lower thickness. Thus, the solvent should be selected based on the desired thickness and uniformity[13,14].

Previous works showed PMMA films of 14 μm as core of planar polymeric waveguides[3]. Thickeness around 10 μm improving the connect in optical fiber, waveguides and optical interconect[3]. Then, the present work we attempt to compare the influence of spin-coating speeds and the concentration of PMMA dissolved in different solvents. A treatment of plasma on PMMA surface films around 10 μm of thickness was carried out. The morphology and roughness of PMMA films, before or/and after plasma treatment, were analyzed by atomic force microscopy (AFM), scanning electron microscopy (SEM), water contact angle measurement, refractive index and proflilometry.

2. Materials and Methods 2.1 Materials PMMA was purchased from Rohm & Hass (Philadelphia, USA). Xylene, MIBK (methyl isobutyl ketone) and chloroform were supplied by Synth (Diadema, Brazil). All reagents were analytical grade and were used as received without further treatment. (100) Silicon substrates (10-20 ohm.cm, one-side polished, 0.610-0.640 mm thickness, and 75 mm diameter) were supplied from University Wafer (Boston, USA).

195

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

Padilha, G. S., Giacon, V. M., & Bartoli, J. R. same level (Figures 2a and 2b). After polymerization, the samples were wrapped in paraffin and stored in a petri dish for further analysis.

2.2 Methods 2.2.1 Gel Permeation Chromatography Number average molecular weight (Mn), weight average molecular weight (Mw), peak molecular weight (Mp) and poly-dispersity (Mw/Mn) of PMMA and the polymer standards were determined by gel permeation chromatography (GPC) (Waters 410) equipped with three columns arranged in sequence (HR-3, HR-4 and HR-5) with the stationary phase of styrene-divinylbenzene with pore dimensions of 103, 104 and 105 Å, respectively. A constant flow of 1 mL/min of tetrahydrofuran (THF) solution was used. PMMA samples were solubilized at 0.1% (w/w) in THF and injected into the chromatograph (100 μL). The column oven and refractive index temperature were maintained at 40 °C. The molecular weight was determined using the software program Millennium 2010.

2.2.2 Preparation of PMMA solutions for spin-coating and plasma PMMA films PMMA coatings were prepared by spin-coating using a PWM32 spin-coater by Headway Research Inc. The spin‑coating solution consisted of PMMA (Plexiglas – VO 052) dissolved at a concentration in distinct solvents – xylene, a mixture of MIBK-xylene, and chloroform. The ratio of MIBK-xylene was 1:1 (v/v). The concentration of polymer and spin-coating speed were selected to provide a homogenous coating thickness over the entire silicon. The range and level of the independent variables and coded values are presented in Table 1. The presence of solvents in the PMMA films was monitored until mass was kept constant.

PMMA films with thicknesses of 10 μm were exposed to CHF3 gas plasma in a parallel plate reactor (ACG 3PT, Serial 587 - ENIRF Inc.) connected to a radio frequency (RF) power source (13.56 MHz). The plasma chamber is made of stainless steel and the electrode uses water as a coolant. The chamber can accommodate up to 12 samples altogether in one round of deposition. The gas inlet is located at the top of the chamber, while the bottom has circular grooves to fit samples and position them at the Table 1. Factors and their levels used for two-level factorial design experiment.

Test set Concentracion (g/L) Speed (rpm)

Level botton (-1) 10 1,000

Central point (0) 15 1,500

Level up (+1) 20 2,000

2.3 Characterization of PMMA films 2.3.1 Thickness measurements and refractive index Gravimetry (10–5 g) technique was used to estimate thickness of the PMMA films, using density of pure PMMA. The thickness of PMMA films and refractive index of the fluorinated film and PMMA film was confirmed by metricon prism coupler (METRICON 2010, λ= 632.8 nm).

2.3.2 Atomic force microscopic measurements of surfaces The surface topography and roughness of the PMMA films were studied using an atomic force microscopy (AFM) (Auto Probe CPII Park Scientific Instrument). Scans of 100 μm are recorded in noncontact mode with a silicon probe with a curvature radius of less than 10 nm, a constant force of 1.75 N/m and operating at a resonance frequency of 155 kHz. The root mean square roughness was calculated over the surface (5x5 μm).

2.3.3 Surface morphology using scanning electron microscopy The morphologies of silicon and PMMA films surfaces were observed using a scanning electron microscopy (SEM) (440i LEO). Samples were cut (0.7 x 0.4 cm) and covered with gold, operating at an acceleration voltage of 5 kV. 2.3.4 Profilometry measurements of roughness Roughness parameters of silicon and PMMA films before and after plasma treatment were obtained by profilometry (Dektac-Veeco - 6M). A diamond probe traversed all samples (5,000 μm - central region for 30 s, with a vertical resolution of 650 kÅ) in different areas.

2.3.5 Contact angle measurements The water contact angles of all the surfaces were measured using contact angle of surface wetting (Tantec Half-Angle CAM - MICRO), with deionized water. The contact angle was measured at three different points on each sample and the average values of triplicate samples were reported.

3. Results and Discussions The average molecular weight of PMMA studied by GPC is shown in Table 2. The value measured for PMMA is in agreement with the reported in literature for spin‑coated films[15].

PMMA was dissolved in xylene, a mixture of MIBK‑xylene and chloroform. Chloroform was the fastest solvent to dissolve PMMA (2 h at 25 °C ± 1 °C), followed by the MIBK-xylene mixture (8 h at 50 °C ± 2 °C). Xylene was the solvent with Table 2. Mn (number molecular weight), Mp (peak molecular weight), Mw (weight average molecular weight) and Mw/Mn (poly-dispersity) of the PMMA.

M n (g/mol) Μ p (g/mol) Figure 1. Spin-Coating procedure .

[7]

196

PMMA

73.143

120.045

M w (g/mol)

M w/M n

115.116

1.57

Polímeros, 27(3) , 195-200, 2017

Effect of solvents on the morphology of PMMA films fabricated by spin-coating greatest dissolution time (9 h at 50 °C ± 2 °C). Although MIBK solvent can also be used to solubilize PMMA, it was not used in this work because of its solubility parameter.

The miscibility of a polymer can be determined quantitatively by the difference between δP (polymer solubility) and δS (solvent solubility), where solvents may be classified as good, medium and bad for a given polymer[16]. Based on this, chloroform, which has a solubility parameter range around 19.03 (J1/2 cm-3/2) is a good solvent for PMMA that has solubility parameter of 19.0 (J1/2 cm-3/2). MIBK has no a good solubility parameter for PMMA - 17.2 (J1/2 cm-3/2), which provides a difference greater than 1.5. In the case of using the mixture of MIBK-xylene, the solubility is the average of those of the pure solvents[17]. However, it is not just the type of solvent/polymer that influences the process. Other variables inherent to the structure of the polymer significantly contribute to the solubility process. High crystallinity and high molecular weights have a negative effect, which makes it more difficult to choose an appropriate solvent[18].

The response surface methodology for seven factorial design experiments for each solvent is presented in Table 3.

Data analysis was done using the Statistica program (version 6.0). It was found that both the dissolved PMMA into xylene and the MIBK-xylene mixture have induced thinner thicknesses than the chloroform solvent, as expected, considering similar conditions. Variance analysis of the films thicknesses obtained using factorial design is presented in Table 4. It shows that the Fcalculated value for the regression using chloroform was 68.4, while the Fcritical value (p = 0.05, degrees of freedom for regression: 3; degrees of freedom for residual: 3) was 9.28, validating the response surface shown in Figure 3 with a determination coefficient (R2) of 0.98.

The mathematical model generated is shown in Equation 1 (X1: chloroform concentration and X2: spin-coating speed).

Other solvents, by Fcalculated less than the Fcritical value were not shown in the response surface. thickeness =+ 12.1 7.28* X1 − 2.05* X 2 − 1.29* X1 * X 2 (1)

Figure 2. (a) plasma apparatus; (b) internal reactor (bottom electrode).

Table 3. Thickness of PMMA films dissolved in xylene, MIBK-xylene mixture and chloroform.

PMMA concentration (g/L)

Spin-Coating speed (rpm)

10 20 10 20 15 15 15

1,000 1,000 2,000 2,000 1,500 1,500 1,500

Thin films (μm) MIBK- xylene 1.69 ± 0.2 7.57 ± 0.1 1.45 ± 0.2 5.21 ± 0.3 4.42 ± 0.2 4.16 ± 0.5 5.55 ± 0.1

Xylene 1.99 ± 0.3 7.70 ± 0.5 1.48 ± 0.2 4.86 ± 0.3 1.77 ± 0.8 2.56 ± 0.3 2.04 ± 0.5

Chloroform 6.02 ± 0.1 23.17 ± 0.3 4.48 ± 0.2 16.48 ± 0.3 10.58 ± 0.6 11.56 ± 0.5 12.41 ± 0.3

Table 4. Variance analysis of the PMMA films thickness.

Xylene Effect Regression Error Total

SS

DF

MS

Fcalculated

Fcritical

24.82 0.32 31.23

3 2

8.27 0.16

3.86

9.28

SS

DF

MS

Fcalculated

Fcritical

31.85 2.34 34.91

3 2

10.62 1.17

10.41

9.28

SS

DF

MS

Fcalculated

Fcritical

235.99 1.68 239.46

3 2

78.66 0.84

68.4

9.28

MIBK-xylene Effect Regression Error Total

Chloroform Effect Regression Error Total

Polímeros, 27(3) , 195-200, 2017

197

Padilha, G. S., Giacon, V. M., & Bartoli, J. R. However, using a nonlinear first-order model for the film thickness variable (Equation 1), the exact value was determined: 1,900 rpm (spin-coating speed) and 15.36 (g/L) (concentration of PMMA) for films around of 10 μm.

To understand the behavior of PMMA films dissolved in xylene, the mixture of MIBK-xylene and chloroform, atomic force microscopy study has been carried with PMMA films obtained in 20 g/L and 1,000 rpm (xylene and MIBK-xylene) and 15.36 g/L and 1,500 rpm for chloroform. PMMA films dissolved in xylene with thickness of 7.70 μm (Figure 4a) and MIBK-xylene with thickness of 7.57 μm (Figure 4b) showed increasing peaks and values (276 nm and 237 nm, respectively) compared to PMMA films dissolved in chloroform (171 nm) (Figure 4c). The attained thin films of both the xylene and MIBK-xylene mixture showed morphological changes with larger formations of convection cells than the film dissolved in chloroform. This behavior seems to be associated with the physico-chemical properties of the solvents. In case of solvent evaporation, convection cells structures can be formed[19]. Yamamura et al.[19] showed that solvent-drying rates provided quantitative information for the concentration variation in the phase-separating solutions, which lead to the formation of the particular cellular patterns. With decreasing air velocity in polymeric films, the transition from the cellular

Figure 3. Response surface of the effect of thickness of PMMA films (μm) as function of spin-coating speed (rpm) and chloroform concentration (g/L).

patterns to the pattern-free structures occurred. However, these features can be used to control the process conditions, such as rotation and spinner deposition time, as well as the solids concentration in solution.

Surface roughness measurement using AFM showed that average root mean square (rms) of PMMA films in xylene, MIBK-xylene and chloroform was 20, 20 and 15 nm, respectively. However even using the non-contact mode, PMMA films showed stripes probably by film fragility. These stripes were observed in all samples. PMMA films dissolved in chloroform showed reduction of roughness using AFM. The morphological in the silicon substrate and PMMA films dissolved in chloroform over silicon were observed using scanning electron microscopy and the SEM micrographs are showed in Figure 5. Silicon exhibited a plane surface (Figure 5a), but when PMMA was dissolved in chloroform we observed roughness by formation of clusters and clumps, with structures non-uniform in the PMMA films deposited on silicon. Figure 5b shows the micrograph of PMMA films over silicon exhibiting rough surface structure. To discuss the appearance of PMMA film dissolved in chloroform, the film was removed from the silicon and the micrograph is show in Figure 5c. The appearance of the film without the silicon shows great morphological difference indicating clearly the intrinsic characteristic of convection cells structures confirmed by the AFM analysis.

Padilha et al.[20] studied the effect of plasma treatment on the deposition of films fluorinated over PMMA, the results showed that films treated using CHF3 plasma obtained acceptable thickness for optical devices. Under these conditions, a profilemeter analysis has been carried out. Their results are shown in Table 5. These results showed that the time was a factor in films that were smoother after plasma treatment using 0.7 torr - 1423 Å (40 min) to 4402 Å (20 min) on 10 μm PMMA films. This behavior has also been observed in samples at 0.3 torr and 30 min as shown in Table 5. Figure 6 shows a comparison of the silicon substrate, whose value was not detected by the device due to measured roughness above 25 Å (silicon shows a low roughness, close to 3 Å[21]), and the both roughness of PMMA film dissolved in chloroform and plasma deposited on PMMA (0.7 torr and 20 min.).

In fact, it is possible to fabricate optical polymeric films.

Therefore, just varying plasma time over PMMA films, low roughness can be obtained. Giacon et al.[3] showed that thickness for the fluorinated film around 0.57 μm was

Figure 4. Atomic force microscopy (a) PMMA dissolved in xylene; (b) PMMA dissolved in MIBK-xylene; (c) PMMA dissolved in chloroform.

198

Polímeros, 27(3) , 195-200, 2017

Effect of solvents on the morphology of PMMA films fabricated by spin-coating

Figure 6. Roughness analysis of silicon, PMMA film dissolved in chloroform and plasma over PMMA film.

Contact angle measurement for the PMMA dissolved in chloroform (before plasma treatment) was 71.3° ± 3.4, where the typical value is 70.5° ± 3.1[22]. PMMA samples exposed to CHF3 plasma shows a significant increase in the contact angle (the average of all samples after plasma treatment was 105.5° ± 3.4) in relation to the original PMMA and very close to the reported for fluorinated polymer – PTFE (110°)[23], proving that plasma treatment can get a good modification for films. This modification of PMMA film surface can be explained by the presence of fluorine atoms that decrease the surface free energy, increasing the values of contact angle[24,25].

PMMA studies showed low level of attenuation in optical devices. Usually, the cladding has slightly smaller refractive index than pure PMMA. Bartoli et al.[26] observed greater speed and higher quality of transmission in polymeric optical fibers. PMMA was coated with hydrofluorocarbon (HFC) plasma. The reduced gradually of the refractive index in coating layer allowed performance gains.

Figure 5. Scanning electron microscopy (a) silicon; (b) PMMA dissolved in chloroform; (c) PMMA film without the silicon.

Table 5. Plasma films roughness by profilometry.

Plasma Process Pressure (torr) Time (min) 0.3 20 0.7 20 0.5 30 0.3 40 0.7 40

Roughness (Å) After plasma 2059 ± 125 1423 ± 101 3400 ± 98 3701 ± 113 4402 ± 99

suitable for cladding of PMMA with core of 14 μm. While Padilha et al.[20] observed 0.61 μm of fluorinated polymer at 0.7 torr and 20 min. In this present work, PMMA films treated with plasma showed low roughness at this plasma condition. This observation can be explained by decrease of the contact time of CHF3 with the samples.

Polímeros, 27(3) , 195-200, 2017

The refractive index of PMMA solved in chloroform was 1.49. However, after of the plasma treatment the refractive index of 1.41 was attained. In similar cases, when plasma of CF4 + H2 was deposited on PMMA, the refractive index was 1.43[26]. Thus, plasma treatment used in this study to produce polymeric optical films allows polymer waveguides to be manufactured, with a numerical aperture (NA) near 0.5. Polymeric structures generally have been manufactured with an NA between 0.3 and 0.5, and the structure of silica with NA of about 0.14 is attained, as previously reported[27].

The use of chloroform as a solvent can contribute to an acceptable roughness for polymeric optical films.

4. Conclusions PMMA films dissolved in chloroform showed acceptable roughness for thin films when compared with other solvents – xylene and the mixture of MIBK-xylene. PMMA dissolved in chloroform - 15.36 (g/L), produced films with thickness of 10 μm by spin coating. These films were exposed to CHF3 plasma. Using chloroform to dissolve PMMA, tiny and shallow cells were created according to AFM results, 199

Padilha, G. S., Giacon, V. M., & Bartoli, J. R. while deep cells with frequent cavities were observed for different solvents. Both SEM and AFM analyses have revealed typical cell structures. Plasma treatment has favored to decrease the roughness, as detected using profilemeter.

Profilemeter results showed different values of roughness according to plasma process (1423 ± 101 to 4402 ± 99).

The contact angle confirmed modification on the PMMA film surface after plasma treatment. We also investigated the refractive index after plasma treatment. In this case, the refractive index was lower than the PMMA film alone.

The use of plasma for cladding using PMMA core presents satisfactory results and may be an alternative for polymeric optical devices application.

5. Acknowledgments The authors wish to thank Dr. Marcelo Nelson Páez Carreño (LME/PSI/EPUSP), Dr. Marco-Aurélio de Paoli (IQ/UNICAMP), Dr. Mauricio Kleinke (IFGW/UNICAMP) and Dr. José Alexandre Diniz (DFIS/FEEC/UNICAMP).

The research was supported by CAPES and CNPQ (Brazil).

6. References 1. Borges, A. M. G., Benetoli, L. O., Licinio, M. A., Zoldan, V. C., Santos-Silva, M. C., Assreuy, J., Pasa, A. A., Debacher, N. A., & Soldi, V. (2013). Polymers films with surfaces unmodified and modified by non-thermal plasma as new substrates for cell adhesion. Materials Science and Engineering C, 33(3), 1315-1324. PMid:23827577. http://dx.doi.org/10.1016/j. msec.2012.12.031.

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12. Liu, Z., Zhao, J.-H., Liu, P., & He, J.-H. (2016). Tunable surface morphology of electrospun PMMA fiber using binary 200

solvent. Applied Surface Science, 364, 516-521. http://dx.doi. org/10.1016/j.apsusc.2015.12.176.

13. Tippo, T., Thanachayanont, C., Muthitamongkol, P., Junin, C., Hietschold, M., & Thanachayanont, A. (2013). The effects of solvents on the properties of ultra-thin poly (methyl methacrylate) films prepared by spin coating. Thin Solid Films, 546, 180-184. http://dx.doi.org/10.1016/j.tsf.2013.05.022.

14. Mohajerani, E., Farajollahi, F., Mahzoon, R., & Baghery, S. (2007). Morphological and thickness analysis for PMMA spin coated films. Journal of Optoelectronics and Advanced Materials, 9(12), 3901-3906. Retrieved in 9 June 2017, from https://joam.inoe.ro/index.php?option=magazine&op=view &idu=1125&catid=21 15. Hass, D. E., Quijada, J. N., Picone, S. J., & Birnie, P. (2003).

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21. Cicala, G., Milella, A., Palumbo, F., Favia, P., & d’Agostino, R. (2003). Morphological and structural study of plasma deposited fluorocarbon films at different thicknesses. Diamond and Related Materials, 12(10-11), 2020-2025. http://dx.doi. org/10.1016/S0925-9635(03)00293-0.

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Pattern formation on aluminum electrodes. Journal of the Electrochemical Society, 138(9), 2763-2770. http://dx.doi. org/10.1149/1.2086051.

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Journal of Polymer Science, 27(4), 1267-1286. http://dx.doi. org/10.1002/pola.1989.080270413.

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Received: Oct. 06, 2016 Revised: Nov. 18, 2016 Accepted: Jan. 05, 2017 Polímeros, 27(3) , 195-200, 2017

http://dx.doi.org/10.1590/0104-1428.10716

Evaluation of the mechanical and thermal properties of PHB/canola oil films Cláudia Daniela Melo Giaquinto1, Grasielly Karine Martins de Souza1, Viviane Fonseca Caetano1 and Glória Maria Vinhas1* Laboratório de Materiais Poliméricos e Caracterização, Departamento de Engenharia Química – DEQ, Universidade Federal de Pernambuco – UFPE, Recife, PE, Brazil

1

*gmvinhas@yahoo.com.br

Abstract Packages are essential for the food processing industry. Among the innovative alternatives there is antimicrobial packaging, which aims to reduce or inhibit microbial growing on the food surface. One potential to produce this type of package is poly(3-hydroxybutyrate)-PHB additivated with canola oil. In this work, films of PHB additivated with canola oil were produced in different compositions. Mid-infrared records, tensile mechanical testing and thermal analyses were performed on the films. The results of the mechanical tests indicated that the addition of canola oil to the polymeric matrix of PHB increases the material flexibility. The thermal analyses results showed that the addition of canola oil changes the thermal properties of PHB, such as the melting and crystallization temperatures, maximum crystallization rate and relative crystallinity. The knowledge of these properties is fundamental for the manufacturing process of polymeric materials, due to the specifications required for these materials in the intended applications.

Keywords: antimicrobial packaging, canola oil, mechanical properties, poly(3-hydroxybutyrate), thermal properties.

1. Introduction Packages are essential to the food processing industry since they are responsible for maintaining the quality of food products[1,2]. Among the types of packaging used for foods, there are passive and active packaging[3]. The latter is used due to its performance under changing environmental conditions to maintain the sensorial properties of the food contents, thus providing quality assurance, increase of shelf lifetime, besides considerations of hygiene and food safety[4,5].

There are various types of active packaging[3], among which antimicrobial stands out, developed in order to reduce the microbial growing on the food surface[6].

One alternative for antimicrobial packaging is the use of biodegradable polymers with vegetable oils that present antimicrobial activity. Biodegradable packages have been studied as an alternative to those made with synthetic plastics, which are known to have a negative impact on the environment[7]. Among the options of biodegradable polymers available for packaging[8], there is the poly(3-hydroxybutyrate)-PHB, which is a potential thermoplastic used to replace polymers obtained from petrochemical industry[9]. PHB is an aliphatic polyester that is biodegradable in water and carbon dioxide under environmental conditions, sustainable, durable, produced from several microorganisms, and with some characteristics similar to polypropylene, a synthetic polymer[10-13].

Vegetable oils are derived from a larger group of chemical compounds known as fats or lipids[14]. One of these is canola oil, which has a huge potential for use in the manufacturing of bioplastics[15]. Canola oil is compounded mainly from oleic, linoleic and alpha-linolenic acids[16].

These components are long chain unsaturated fatty acids that

Polímeros, 27(3) , 201-207, 2017

have antibacterial activity[17]. Thus, PHB films additivated with canola oil are an alternative for use in the production of antimicrobial packages. In the literature, there are studies that have evaluated the antimicrobial activity of eugenol in films of PHB/eugenol[18] and the influence of d-limonene in the mechanical, thermal and barrier properties in the PHB/PLA/d-limonene system[19].

In this work, we evaluated both the mechanical and thermal properties of PHB films additivated with canola oil in different compositions for purposes of antimicrobial packaging . Also in this work, we have evaluated the antibacterial activity of canola oil. The knowledge of these properties is fundamental in the manufacturing process of polymeric materials due to the specifications that these materials must meet in their intended applications.

2. Materials and Methods 2.1 Preparation of films PHB in powder was donated by PHB Industrial S/A. The canola oil used was the Purilev trademark of the Cargill company. This oil is compounded mainly from oleic (50-70%), linoleic (15-30%) and alpha-linolenic acids (5-13%). The solvent used was chloroform from the Vetec brand. The films were produced by the solution casting technique using 1.3 g of PHB and 50 mL of chloroform.

PHB films additivated with canola oil were prepared in different amounts (0, 2, 6 and 10% w/w). The films had an average thickness of 0.09 ± 0.02 mm. In total, 36 samples were prepared by casting.

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

Giaquinto, C. D. M., Souza, G. K. M., Caetano, V. F., & Vinhas, G. M. 2.2 Antimicrobial activity of canola oil

3. Results and Discussions

The activity of the canola oil was evaluated by disk diffusion assay[20] with medium Plate Count Agar (PCA).

Filter paper disks of 2 cm diameter were utilized. Aliquots of 0.5 ml of E. coli (ATCC 8739) in the order of 107 CFU/ml, were quantified by turbidity on the Mcfarland comparison scale. They were inoculated into the PCA by the pour plate method. After solidification of the PCA, these were placed on discs soaked with canola oil, in the center of a Petri dish.

The plates were incubated at 35 °C for 48 h.

3.1 Mid-Infrared spectra of the films

2.3 Mechanical tests Mechanical tests of the films were made in the universal machine for tensile testing, model DL-500 MF from EMIC brand following the ASTM D882-12 standard[21]. The tests took place at room temperature without humidity control.

Assays were performed under the following conditions: load cell of 500 N; jaw speed of 5 mm/min; initial distance between the jaws of 40 mm; and specimen dimension of 2.5 x 7.5 cm. For the mechanical test, nine films for each composition were analyzed.

2.4 Spectral acquisition Mid-infrared (MIR) spectra of the films and canola oil were recorded using a Spectrum 400 FT-IT/FT-NIR spectrometer from Perkin Elmer brand with Horizontal Attenuated Total Reflectance-HATR accessory under the following conditions: spectral region from 4000 to 650 cm-1 with resolution of 4 cm-1 and 16 scans. From the spectra recorded, a Principal Component Analysis-PCA was performed in the software The Unscrambler version 9.7.

2.5 Thermal analyses Thermal analyses of the films were performed using a calorimeter from Mettler Toledo brand, DSC STAR and SYSTEM model. Each sample used had mass of 4.0 ± 1.0 mg. The heating took place from -10 ºC to 200 ºC at a heating/cooling rate of 12 ºC/min. The experiments were performed under a nitrogen cooling system. The data were processed using the software Integral© - Version B[22].

It was assumed that the latent heat of the 100% crystalline PHB was 146 kJ/kg[23].

The evolution of the relative crystallinity x observed at different temperatures was computed from the latent heat of the exothermic crystallization peaks as: t2 1 t x= ∫ J ( t ′ ) − J 0 ( t ′ ) dt ′, E0 = ∫ J ( t ) − J 0 ( t ) dt (1) E0 t1 t1

The variables in Equation 1 are: J(t) is the heat flux recorded by DSC; J0(t) is a suitable virtual baseline; t1 e t2 are the onset and end times of the crystallization event; and t is an intermediate time related to a given value of x (0< x < 100%). E0 represents the total heat for the observed crystallization process.

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Figure 1 illustrates representative spectra obtained from mid-infrared scans of canola oil, PHB film, and PHB film additivated with canola oil in the compositions of 2% w/w (PHB/2% CO), 6% w/w (PHB/6% CO) and 10% w/w (PHB/10% CO).

PHB has the following characteristic bands: 1720-1650 cm−1 corresponding to the carbonyl stretching vibrations of ester groups of the polymer; 1460-1380 cm−1 corresponding to asymmetric stretching and CH2 groups; 1300-1100 cm-1 corresponding to symmetric and asymmetric stretching of the COC group; 1274 and 1226 cm−1 corresponding to crystalline phase of polymer; and 1261 and 1180 cm−1 corresponding to amorphous phase of the polymer[24]. Canola oil has the following characteristic bands: 3012 cm-1 correspond to the C-H stretching vibration; 2920 and 2854 cm-1 corresponding to the symmetric and asymmetric stretching vibration of the aliphatic CH2 group, respectively; 1739 cm-1 corresponding to the ester carbonyl functional group of triglycerides; 1457 cm−1 corresponding to the bending vibrations of the CH2 and CH3 aliphatic groups; 1378 cm−1 corresponding to the bending vibrations of CH2 groups; 1244 and 1165 cm−1 corresponding to the stretching vibration of the C–O ester groups; and 713 cm−1 corresponding to the overlap of the CH2 rocking vibration and the out-of-plane vibration of cis-disubstituted olefins[25]. The incorporation of canola oil (CO) can be identified in the spectra of films of PHB/CO by increasing the maximum peak intensity corresponding to the ester carbonyl functional group which is detected between 1620 cm-1 and 1810 cm-1. In PHB films to which 10% canola oil was added, it is also possible to visualize bands with lower absorption intensities, such as those referring to the wavenumbers 2920 cm-1 and 2854 cm-1.

To confirm that it had the incorporation of canola oil into the matrix analysis, Principal Component Analysis (PCA) was performed. PCA[26] is a chemometric technique used to evaluate the variations in the data. The PCA extracts the key information variations from the original variables and reduces the dimensional space, creating a reduced number of variables called principal components (PC’s)[27-29].

The PCA technique consists of decomposing the data matrix X (m samples and n absorbances) into a product of two matrices, the matrix of the scores (T) and loadings (P), together with the error matrix (E)[30,31]. The graph of the scores represents the coordinates of the samples in the PC space and the loadings represent the relevance of the original variables in each PC. From the PCA it is possible to identify samples that have the same spectral characteristics.

This can be verified by grouping the samples in a graph of the scores. The scores plot is illustrated in Figure 2.

Figure 2 shows that the PCA separated the samples into distinct groups. Groups identified by the colors dark blue, green, red and light blue refer to films of PHB, PHB/2% CO, PHB/6% CO and PHB/10% CO, respectively.

The PCA results showed that there was an incorporation of canola oil in the polymer matrix since the samples were grouped according to the PHB/CO composition. Note that as the concentration of canola oil increases, a distancing of the Polímeros, 27(3) , 201-207, 2017

Evaluation of the mechanical and thermal properties of PHB/canola oil films

Figure 1. Mid-Infrared spectra of canola oil, PHB film and PHB/CO film in different compositions.

Figure 2. Score plots of PC1 x PC2 of PHB film and PHB/CO films.

samples in relation to the PHB/0% CO occurs, evidencing the incorporation of the oil into the polymeric matrix.

The percentages of variance explained by PC1 and PC2 were 88% and 10%, respectively.

3.2 Mechanical properties Table 1 shows the values obtained for the mechanical properties tensile strength (TS), percentage elongation at break (%E) and elastic modulus (EM). The mean values of ​​ the mechanical properties obtained through the mechanical tests were compared statistically by Duncan’s test with a significance level of 5% (p <0.05).

By the Duncan’s test, the tensile strength the canola oil incorporation can be seen to cause a decrease in the value of this property. It can also be verified that the mean values obtained for the tensile strength with the addition of 6 and 10% showed no statistical differences considering a level of significance of 5%. The tensile strength expresses Polímeros, 27(3) , 201-207, 2017

the maximum strength of the material when under tension[32].

The decrease in the value of this property, in PHB films, when the percentage of additive canola oil is increased, shows that the film has become mechanically less resistant when compared with the pure PHB film. The decrease in the tensile strength value was also observed in the work of Souza et al.[33] with the incorporation of cinnamon essential oil in to polymer starch.

For the percentage elongation at break it can be seen that there was a decrease in the average value of this property only from the incorporation of 10% of canola oil. Elongation at break is a measure of the film’s stretch ability prior to breakage[34]. The other mean values were statistically equivalent when compared to the mean values of pure PHB films. The percentage elongation measures the ability of a film to stretch to its breakup. The incorporation of canola oil from 10% w/w led to a decrease in PHB’s elongation ability. This decrease in the value of %EM 203

Giaquinto, C. D. M., Souza, G. K. M., Caetano, V. F., & Vinhas, G. M. was also observed in the work of Hauser et al.[35] with the incorporation of myrtle essential oil to methylcellulose.

There was a decrease in the value of the elastic modulus.

An incorporation of 6 and 10% of canola oil showed no statistical differences for the 5% significance level. The elastic modulus is related to the stiffness of the material[36]. The lower the value of the modulus of elasticity, the less rigid the film will be and vice versa. Thus, the decrease caused by incorporating canola oil results in less rigid films. As PHB is a rigid polymer, this change in stiffness is very positive when the films are intended for food purposes, whose packaging requires greater flexibility. The decrease in the value of EM was also observed in the work of Noronha et al.[37] with the incorporation of nanocapsules of α-tocopherol in methylcellulose film.

3.3 Thermal properties The output data from the Integral© program regarding the crystallization in the molten state and the cold crystallization are shown in Table 2 together with the melting temperature.

Tm shows the melting temperature of the material under study, Tc,i is the temperature of crystallization, ΔHc,i is the latent heat of crystallization and Xc,i is the degree of crystallinity, considering i=1 for the crystallization from the molten state and i=2 for the cold crystallization. Regarding the melting temperature, a decrease is observed due to the increase in the percentage of canola oil in the polymeric matrix. The melting temperature of PHB is very close to its thermal degradation temperature (~180°C)[38]. The decrease in the PHB melting temperature increases its possibilities for applications, because Its applicability is limited due to their small window of processability[39]. It is also observed that There was an increase in the crystallization temperature for the additivated films when compared to PHB films, both from the molten state as from the cold. On cooling, this implies an early crystallization and on heating a delayed crystallization.

The values of the crystallized fraction from molten state of the PHB/6% CO and PHB/10% films are much smaller compared to PHB and PHB/2% CO films. This shows that the crystallization of PHB/CO and 6% PHB/10% films occurred, the most part, when cold.

3.3.1 Crystallization from the molten state Figures 3a and 3b illustrate the curves of the rates of crystallization and the relative crystallinity versus temperature, respectively, for the crystallization from the molten state, with respect to the films of PHB and PHB additivated with canola oil in different concentrations.

Figure 3a shows that the maximum rate of crystallization occurred between 65°C and 75°C for the films additivated with canola oil, while for the PHB film the temperature was approximately 54°C. It is also possible to observe that the maximum crystallization rate values were lower for the films containing 2 and 6% of canola oil compared to PHB film.

These results show that the percentage of canola oil has a direct influence on the polymer crystallization process.

Figure 3b shows that the film additivated with 2% canola oil begins to crystallize at higher temperatures and finishes at higher temperatures than the PHB film, whereas the additivated films with 6 and 10% start to crystallize at lower temperatures, and finish at higher temperatures than PHB films.

3.3.2 Cold crystallization For the analysis of cold crystallization behavior, Figures 4a and 4b illustrate the curves of the crystallization rate as a function of the temperature and the relative crystallinity as a function of the temperature, respectively, for films of PHB and PHB additivated with canola oil in different compositions.

Figure 4a shows that the maximum crystallization peaks of films were 2.2, 1.31, 1.31 and 1.27 for additions of 0, 2, 6 and 10% w/w CO. In this figure it is verified that there was a reduction in the value of the maximum crystallization rate for the additivated PHB film. Figure 4a also shows the temperatures that occur at the crystallization rate peaks. The temperatures are 41.9, 48.5, 48.5 and 50°C. This shows that each increase in temperature causes a faster crystallization rate, thus a higher peak, for the additivated films. Figure 4b shows that the increase in temperature causes a delay in the cold crystallization of the added PHB when compared to the non-additive film. This fact can be seen at a fixed temperature, where it is possible to see that the PHB has a greater relative crystallinity than the films with additives. Figures 4a and 4b show that the curves for

Table 1. Average values obtained for the mechanical properties tensile strength, percentage elongation at break and elastic modulus.

Sample PHB PHB/2% CO PHB/6% CO PHB/10% CO

TS (MPa) 20.87+0.85a 16.50+1.58b 10.45+0.15c 10.39+0.32c

%E 4.12+0.41a 4.11+0.40a 4.01+0.12a 3.37+0.13b

EM (MPa) 801.30+24.46a 592.03+12.62b 375.20+22.06c 345.13+24.37c

Means with the same letter in the same column do not differ with p <0.05 for the Duncan’s test.

Table 2. Data melting, crystallization from the molten state and cold crystallization of the samples.

Sample

Tm (ºC)

Tc,1 (ºC)

ΔHc,1 (kJ/kg)

Xc,1 (%)

Tc,2 (ºC)

ΔHc,2 (kJ/kg)

Xc,2 (%)

PHB PHB/2% CO PHB/6% CO PHB/10% CO

175.3 174.8 168.6 167.6

54.06 72.17 70.20 66.77

20.33 26.56 6.24 3.61

13.92 18.19 4.27 2.47

41.97 48.72 49.88 48.63

19.15 9.83 26.52 29.33

13.12 6.74 18.16 20.09

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Evaluation of the mechanical and thermal properties of PHB/canola oil films

Figure 3. (a) Graph of the crystallization rate versus temperature of PHB film and PHB films additivated with canola oil in different compositions; (b) Graph of the relative crystallinity versus temperature of PHB film and PHB films additivated with canola oil in different compositions.

Figure 4. (a) Graph of the crystallization rate versus temperature of PHB film and PHB/CO films with different compositions; (b) Graph of the relative crystallinity versus temperature of PHB film and PHB/CO films with different compositions.

the PHB added with 2 and 6% w/w CO had a similar result, with overlaps occurring. This can be attributed to the additive values being very close and relatively low when compared to the mass of the polymer matrix.

3.4 Antimicrobial activity of canola oil Figures 5a and 5b show the result of the antimicrobial activity of the sample without canola oil (control sample) and with canola oil, respectively, in the presence of E. coli bacteria. Figure 5a shows that there was no antimicrobial activity in the control sample. In Figure 5b the inhibition halo was observed, showing that there was antimicrobial activity in the presence of E. coli bacteria. The inhibition halo is observed around the sample, evidencing the absence of microorganisms. This result proves that canola oil is an antimicrobial agent and has the potential to be used in antimicrobial packages.

4. Conclusions Infrared spectra in various compositions showed that canola oil was incorporated into a polymeric matrix of PHB. The results of an analysis of the mechanical properties showed Polímeros, 27(3) , 201-207, 2017

Figure 5. Antimicrobial activity in the presence in E. coli bacteria of the (a) sample without canola oil (control sample); and (b) sample with canola oil.

that the addition of canola oil increases the flexibility of the films. From the results of the thermal properties, it can be seen that the addition of canola oil caused changes in the melting and crystallization temperatures, in the maximum crystallization rate and in the relative crystallinity. Moreover, 205

Giaquinto, C. D. M., Souza, G. K. M., Caetano, V. F., & Vinhas, G. M. the added oil accelerated crystallization from the molten state and delayed cold crystallization. The tested canola oil is an antimicrobial agent that can be used as a natural additive.

These results indicate tendencies of PHB behavior when canola oil is added, suggesting a potential improvement for use in antimicrobial packaging.

5. Acknowledgements The authors thank the Fundação de Amparo Ciência e Tecnologia do Estado de Pernambuco (FACEPE) andthe Coordenação de Aperfeiçoamento de Pessoal de Nível Superior (Capes) for theprovidedscholarships. The English text of this paper was revised by Sidney Pratt, Canadian, MAT (The Johns Hopkins University), RSAdip - TESL (Cambridge University).

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Received: Sept. 13, 2016 Revised: Dec. 21, 2016 Accepted: Jan. 17, 2017

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http://dx.doi.org/10.1590/0104-1428.03416

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

Epoxidized natural rubber and hydrotalcite compounds: rheological and thermal characterization Vanessa Macedo da Silva1,2, Regina Célia Reis Nunes1* and Ana Maria Furtado de Sousa3 Instituto de Macromoléculas Professora Eloisa Mano – IMA, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brazil 2 Centro Universitário Estadual da Zona Oeste – UEZO, Rio de Janeiro, RJ, Brazil 3 Departamento de Processos Químicos, Instituto de Química, Universidade do Estado do Rio de Janeiro – UERJ, Rio de Janeiro, RJ, Brazil 1

*rcnunes@ima.ufrj.br

Abstract Epoxidized natural rubber (ENR) and synthetic non-modified hydrotalcite (HT) compounds were prepared and evaluated. Natural rubber (NR) was epoxidized with 20.6% of epoxy groups from a chemical modification of the latex.

A sulfur-based curing system formulation with accelerators was used. The amounts of HT in the ENR-HT compositions was varied between 0, 2, 3 and 5 phr. All compositions were evaluated as to cure parameters, rheological properties, thermal resistance and crosslink density. The results showed that the mineral filler does not have a significant influence on the cure parameters. Different methods of crosslink density determination were used (swelling at equilibrium and elastic modulus). The results turn out to be equivalent and rise as the amount of filler is increased. The best results were found for the 5 phr hydrotalcite compound (ENR-HT5).

Keywords: crosslink density, epoxidized natural rubber, layered double hydroxide, rheological behavior, thermal behavior.

1. Introduction Layered double hydroxides (LDH), known as anionic clays, are minerals having positively charged layers interleaved with ions and water molecules, so as to maintain neutrality of charge. Anionic clays are not abundant in nature and hydrotalcite (Mg6Al2(OH)16CO3.4H2O) is an example of anionic clay found in nature. Properties like lamellar structure, ease of synthesis and readily controllable particle size are characteristics that have attracted researchers to study LDH and their applications on polymers, especially elastomers[1].

The use of LDH in polymer matrices has been directed to different purposes. In nanocomposites, organically-modified LDH was studied to improve thermal resistance and inflammability properties[2]. The capacity to act like a crosslink agent for different kinds of elastomers has been studied aiming at replacing zinc oxide in rubber formulations[3,4].

Chemical modification of natural rubber, especially epoxidation reaction, is a versatile route to obtain a polymer with better properties comparative to unmodified natural rubber like oil and organic solvent resistance, air permeability and performance for wet roads (for tire applications)[5].

The introduction of epoxy groups into NR chains increases the polarity of ENR and improves the compatibility of this polymer with polar fillers or blends with polar polymers[6].

In this work hydrotalcite was studied in compositions of epoxidized natural rubber synthetized in our laboratory.

The effect of the hydrotalcite content on cure characteristics, rheological properties and on thermal resistance was investigated. Also, for the sake of comparison, crosslink density was studied by two different methods.

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2. Materials and Methods 2.1 Materials High ammonia natural rubber latex (NRL) and 60% of dry rubber content was supplied by Teadit. Nonionic surfactant Ultranex 95 (ethoxilatednonylphenol), supplied by Oxiteno, was used to stabilize NRL during the epoxidation reaction. Magnesium aluminium layered double hydroxide (hydrotalcite) was obtained from Sigma Aldrich and used in its unmodified form. Double ventilated sulfur, Irganox 1010 (Pentaerythritoltetrakis(3-(3,5-di-tert-butyl-4-hydro xyphenyl)propionate)), stearic acid, zinc oxide and TBBS (N-tert-butyl-2-benxothiazolesulfenamide) were used for the curing process. The other chemical products used in this work were formic acid, hydrogen peroxide 30%, ethanol 95%, anhydrous sodium carbonate, toluene and chloroform, all obtained from Sigma.

2.2 Methods 2.2.1 Characterization of hydrotalcite Infrared spectra were collected in a Varian 3100 FT-IR equipment using a total reflection attenuated accessory, between 600 and 4000 cm-1, mode transmittance, 4cm-1 of resolution and 100 scans.

Thermal stability was studied using a thermogravimetric analyser (TGA) Q500 from TA Instruments, between 30°C and 700°C at 10°C/minute under inert atmosphere (nitrogen).

Polímeros, 27(3) , 208-212, 2017

Epoxidized natural rubber and hydrotalcite compounds: rheological and thermal characterization 2.2.2 Chemical modification of NRL – Epoxidation reaction Epoxidized natural rubber (ENR) was obtained containing 20% of epoxy groups as described in[7] with some modifications. NRL, previously stabilized with surfactant, was epoxidized using organic peracid obtained in situ by the reaction between formic acid and hydrogen peroxide.

The reaction was processed during 8 hours at 70°C. In the end, the pH of the epoxidized latex was neutralized by addition of anhydrous sodium carbonate solution, coagulated, dried and characterized.

2.2.3 Characterization of epoxidized natural rubber (ENR) The chemical structure of ENR was investigated in a Varian 3100 FT-IR equipment, between 500 and 4000 cm-1, resolution of 4 cm-1 and 32 scans. The polymer film for analysis was obtained after chloroform evaporation of solubilized ENR. Epoxy groups content was determined by hydrogen nuclear magnetic resonance (1HNMR) in a Varian Mercury VX300 instrument with 300,5Hz of frequency. For the tests, the ENR polymer was solubilized in deuterated chloroform.

2.2.4 ENR/Hydrotalcite compounds preparation ENR/Hydrotalcite compounds were prepared in a LRMR-S two-roll mixing mill by TECH Engineering Company LTD at 30°C according to ASTM D3184. The formulations studied were (in phr): ENR 100.0; stearic acid 2.0; zinc oxide 5.0; sulfur 2.5; TBBS 1.5; Irganox 1010 1.0. The hydrotalcite content varied between 0, 2, 3 and 5 phr. Compounds were identified as ENR-HT0, ENR-HT2, ENR-HT3, ENR-HT5 corresponding respectively 0, 2, 3 and 5 phr of hidrotalcite.

2.2.5 Rheometric properties The rheometric parameters of the compounds were studied using an Alpha Technologies RPA2000 rubber process analyzer. Tests were processed according to ASTM D5289 at 150 °C, oscillating arc 1° and 1.7 Hz of frequency during 60 minutes. As for torque, the measurement tolerance is 0.5% of the working range and for the temperature, accuracy is ± 0.3ºC of the test temperature.

2.2.6 Rheological properties – Sweep temperature The compounds cure was studied by sweep temperature using a RPA 2000 Analyser between 50 and 200°C, 5°C/minute of heating rate, 0.5° of strain and 1 Hz of frequency.

2.2.7 Thermal properties The thermal stability of compounds was studied using a thermogravimetric analyzer (TGA) Q500 from TA Instruments between 30°C and 700°C at 10°C/minute under inert atmosphere (nitrogen).

2.2.8 Crosslink density Crosslink density was determined using two methods, comparatively. One of then was the method based on swelling resulting from contact with an organic solvent as reported in Flory studies[8] and the other one was based on the rheological property of elastic modulus (G’), obtained by the RPA 2000 Analyser[9].

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2.2.8.1 Swelling method Specimens of 20mm x 20mm x 2mm were immersed in toluene kept in closed containers in the dark during 7 days.

The swollen rubber volume (Vr) was measured based on Equation 1 developed by Flory-Rehner[8].  W W   W W   W − W3    Vr = 1 − h  /  1 − h  +  2    (1)  ρ2 ρh   ρ2 ρh   ρ1   

Where Vr is the swollen rubber volume; W1 is the sample weight before swelling; W2 is the swollen sample weight; W3 is the dried sample weight; Wh is the hydrotalcite weight in the sample; ρ1 is the solvent density;ρ2 is the sample density; ρh is the hydrotalcite density.

It is possible to measure the crosslink density through the rubber volume in a swelling network, as described in Equation 2.  1 V = − ln (1 − Vr ) + Vr + .Vr2  / V0  Vr 3 − r   2 

 (2)  

Where ν is the number of chains inside the reticulum in mol/cm3; χ is the interaction parameter between toluene and epoxidized natural rubber (χ=0.39); V0 is the molar volume of the solvent (V0 = 106.2).

2.2.8.2 Rheological properties method To measure crosslink density by rheological properties, the elastic modulus (G’) was determined for uncured and cured compounds at 100°C, 0.25° of strain, 5 Hz and 0.5 Hz, respectively, using a RPA 2000 Analyser[9]. Through this method it is possible to measure the physical crosslink density (Xinitial) relative to entanglements and the total crosslink density (Xtotal), relative to entanglements plus chemical crosslinks. The values of Xinitial and Xtotal can be obtained using Equations 3 and 4, respectively[9].

X total =

G'cured ( 0.5 Hz )

X initial =

2RT

(3)

G'uncured ( 5 Hz ) 2RT

(4)

Where: G’cured is the elastic modulus measured at 0.5 Hz of frequency for cured compounds; G’uncured is the elastic modulus measured at 5 Hz of frequency for uncured compounds; R is the gases constant (8.31 L.KPa/K.mol); and T is the temperature in Kelvin degrees.

The elastic modulus (G’) values were corrected using the Guth-Gold Equation to determine the equivalent rubber modulus without any filler[9]. It was calculated through Equation 5.

(

)

= G ' filled G 'unfilled 1 + 2.5∅ + 14.1∅ 2 (5)

Where: ∅ is the effective volume fraction of filler.

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Silva, V. M., Nunes, R. C. R., & Sousa, A. M. F.

3. Results and Discussions 3.1 Characterization of hydrotalcite The chemical structure of hydrotalcite was investigated using FT-IR. The infrared spectra (Figure 1) showed a large absorption band at 3416 cm-1 referred to hydroxyl groups and hydration water of the hydrotalcite structure. At 1365 cm-1 the asymmetric stretch of carbonate present between the hydrotalcite lamellae could be seen. The absorption bands located at 775 cm-1 and 637 cm-1can be attributed to magnesium oxide and aluminum oxide[10,11].

The thermal stability of hydrotalcite was studied by TGA. Figure 2 showed three degradations steps. The first one, between 50°C and 230°C, was referred to the loss of absorbed water and interlayer water. In step 2, between 230°C and 370°C, the degradation of interlayered carbonate and the loss of hydroxyl groups associated to aluminum and magnesium was observed. Between 370°C and 550°C, in step 3, the total degradation of hydroxyl groups of metals happened simultaneously forming metallic oxides (Al2O3 and MgO) and steam[11].

The efficiency of the chemical modification of natural rubber latex was investigated by 1HNMR. The spectrum of the polymer (Figure 4) showed chemical shifts at 2.7ppm and 5.1ppm referred to the methine hydrogen of epoxy rings and olefin hydrogen of cis-1,4-polyisoprene, respectively.

Chemical shifts referred to methyl and methylene groups adjacent to the epoxy ring were located at 1.6 ppm and 2.1 ppm, respectively[13,14]. The area ratio of signals at 2.7 ppm and 5.1 ppm, as described in Equation 6, yielded the epoxidation content of 20.6%[14].

(

)

= Epoxidation %  Area2.7 ppm / Area2.7 ppm + Area5.1 ppm  .100  

(6)

3.3 Rheometric properties The rheometric parameters of ENR-HT compounds are shown in Table 1.

Generally, the rheometric parameters did not show significant difference. This means that filler addition did not influence negatively on kinetics cure. Also, it indicates that the compounds have the same crosslink density.

3.2 Characterization of epoxidized natural rubber

3.4 Rheological properties – Sweep temperature

The epoxidation of natural rubber latex was confirmed by infrared spectroscopy (Figure 3). Absorption bands at 1249 cm-1 and 870 cm-1 refer to symmetric and asymmetric deformations by epoxy rings, respectively[12].

Elastic torque curves of the sweep temperature study (Figure 5) showed that the presence of hydrotalcite did not influence the rheometric parameters, considering the amounts used in this work, corroborating the rheometry tests.

The torque values decrease with increasing temperature.

Figure 1. Infrared spectra of hydrotalcite.

Figure 3. Infrared spectra of ENR.

Figure 2. Weigth loss derivative curves of Hydrotalcite.

Figure 4. H1-NMR spectra of ENR.

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Epoxidized natural rubber and hydrotalcite compounds: rheological and thermal characterization Table 1. Rheometric characteristics of ENR-HT compounds.

Rheometric characteristics ENR-HT0 ENR-HT2 ENR-HT3 ENR-HT5

ts1 (min)

t90 (min)

ML (dNm)

MH (dNm)

MH–ML (dNm)

2.3 2.3 2.1 2.0

5.3 4.8 4.9 5.4

0.2 0.2 0.2 0.1

21.8 21.7 21.7 21.7

21.6 21.5 21.5 21.6

Table 2. Thermal degradation temperatures.

Thermal degradation temperatures ENR-HT0 ENR-HT2 ENR-HT3 ENR-HT5

Tonset(°C)

Tmax (°C)

355 350 350 348

379 374 376 375

Figure 5. Curves of elastic torque versus temperature of ENR-HT compounds.

The initial torque is 1.3 dNm at 50°C and it decreases to 0.1 at 140°C. This is due to the fact that the polymer chain acquires molecular mobility and, consequently, lower flow resistance. As from 140°C a significant increase in torque values with increasing temperature was observed. It resulted from crosslink formation, making compounds more rigid.

The maximum torque was obtained at 167°C and from that temperature on, torque decreased until 7dNm at 200°C, when the test was stopped. A slightly increase in maximum torque was observed as a function of the hydrotalcite amount in the compound.

Figure 6. Weigth loss derivative curves of ENR-HT compounds.

The sweep test showed that different compositions presented the same cure kinetics. By relating sweep temperature curves with Table 1 it was found that the differences observed in cure time were not significant and that the compositions keep the same behavior until 5% of hydrotalcite content.

3.5 Thermal properties Table 2 shows compounds data for onset (Tonset) and maximum temperature degradation (Tmax). It was observed that compounds containing hydrotalcite had lower thermal resistance. Compounds had lower initial temperature degradation (Tonset) as compared with a compound without filler (ENR-HT0). This fact could be related with the degradation of interlayer carbonates between 230 and 370°C. Derivative weight loss curves (Figure 6) showed shoulders at 420°C for hydrotalcite-containing compounds.

This coincides with hydrotalcite dehydroxilation, when the hydroxyl bound to aluminum and magnesium is degraded.

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Figure 7. Crosslink density obtained by sweeling and elastic modulus method.

3.6 Crosslink density Comparative results of two different methods to obtain crosslink density are described in Figure 7. It is shown that the hydrotalcite amount did not influence crosslink density, corroborating results described in Table 1. Although the results obtained by swelling and elastic modulus show different values it should be considered that such results show the same variation profile between compounds and they are similar. The analysis of crosslink density by elastic modulus properties is a fast and efficient alternative to 211

Silva, V. M., Nunes, R. C. R., & Sousa, A. M. F. measure that property. Also, it is ecofriendly because it does not make use of organic solvents harmful to human health and the environment.

4. Conclusions Epoxidized natural rubber was synthetized with 20.6% of epoxy groups. The presence of hydrotalcite in rubber compounds did not affect the cure kinetics comparatively with compounds without filler. Sweep temperature studies showed that all compounds cured at the same temperature and they showed a slightly increase in torque with hydrotalcite content. Thermal resistance of filler-containing compounds was lower than that of ENR-HT0 due to interlayer carbonates and hydroxyl metals of the hydrotalcite structure which undergo degradation. Crosslink density analyses have comparable results by swelling or modulus elastic method.

5. Acknowledgements The authors are indebted to CNPq, CAPES and FAPERJ for financial support and to Teadit Indústria e Comércio Ltda by supply of natural rubber latex, additives and further materials employed during the course of this research.

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12. Chuayjuljit, S., Yaowasang, C., Ranomg-NA, N., & Potiyaraj, P. (2006). Oil resistance and physical properties of in situ epoxidized natural rubber from high ammonia concentrated latex. Journal of Applied Polymer Science, 100(5), 3948-3955. http://dx.doi.org/10.1002/app.22998.

13. Saito, T., Klinklai, W., & Kawahara, S. (2007). Characterization of epoxidized natural rubber by 2D NMR spectroscopy. Polymer, 48(3), 750-757. http://dx.doi.org/10.1016/j.polymer.2006.12.001.

14. Klinklai, W., Kawahara, S., Mizumo, T., Yoshizawa, M., Sakdapipanich, J. T., Isono, Y., & Ohno, H. (2003).

Depolymerization and ionic conductivity of enzymatically deproteinized natural rubber having epoxy group. European Polymer Journal, 39(8), 1707-1712. http://dx.doi.org/10.1016/ S0014-3057(03)00060-0.

Received: May 28, 2016 Revised: Jan. 13, 2017 Accepted: Feb. 21, 2017

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http://dx.doi.org/10.1590/0104-1428.04116

Influence of addition of silanized nanosilica and glycerol on hydrophobicity of starch using a factorial design Fernando Luis Panin Lopes1, Vicente Lira Kupfer2, Júlio César Dainezi de Oliveira1, Eduardo Radovanovic2, Andrelson Wellington Rinaldi2, Murilo Pereira Moisés3 and Silvia Luciana Favaro1* Departament of Mechanical Engineering, Maringá State University – UEM, Maringá, PR, Brazil 2 Departament of Chemistry, Maringá State University – UEM, Maringá, PR, Brazil 3 Paraná Federal Technological University – UTFPR, Apucarana, PR, Brazil

1

*slfavaro@hotmail.com

Abstract The thermoplastic starch (TPS) is regarded as a promising material for manufacturing packaging and products with biodegradable properties. This study aimed at obtaining hydrophobic starch using silanized silica nanoparticles (nSS) with hexamethyldisilazane. A factorial design 22 with central point was developed to evaluate the influence of glycerol (plasticizer) and nSS addition on the properties of water absorption, solubility and TPS contact angle. The materials morphology was also evaluated by means of scanning electron microscopy. The amount of glycerol and nSS influenced on starch hydrophobic character, for the increase of the glycerol dosage contributed to the increase of absorption and solubility of TPS in water. On the other hand, nSS has greater influence on the characteristics related to the TPS surface, favoring an increase of up to 27% in the contact angle values. Therefore, the sample with the greatest hydrophobic character was obtained by using lowest amounts of glycerol (30%) and highest amounts of nSS (5%).

Keywords: biodegradable polymer, thermoplastic starch, silanized nanosilica.

1. Introduction The search for new polymeric materials, mainly the ones that derivate from renewable sources, has been the reason for the realization of diverse studies[1-4]. A material with promising potential for that end is the starch, for its raw material is very accessible in many parts of the world.

Besides that, in general, it is obtained with low costs, being used at the formation of blends or composites with polymeric materials[1-6].

Starch can be obtained from different vegetable sources, such as rice, maize, potatoes and cassava. These are the most important sources of that carbohydrate[7]. Among the mentioned sources, the cassava starch can be highlighted for the production of thermoplastic starch (TPS), for presenting mechanical, chemical and physical properties that are interesting from the processing and application point of view[8]. Because of these characteristics, cassava starch has been studied by many Brazilian researchers, aiming at producing biodegradable plastics[9-11].

In the process of making biodegradable plastics with starch, it is necessary to add a plasticizer, once starch presents a degradation temperature lower than its fusion point. The most commonly used plasticizers are water and glycerol[11,12]. Besides that, compatibilizing agents, such as block or grafted copolymers can be used. These agents cause the material tenacity to change in relation to the thermoplastic, because they make the size of the disperse phase decrease[13].

The thermoplastic starch has good characteristics for the oxygen barrier. Nevertheless, many limitations are noticed in relation to its hydrophilic characteristics and permeability

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to water vapor, which are responsible for the deterioration of mechanical properties and for dimensional stability[8-10,14-18].

This way, aiming at the improvement of the thermoplastic starch (TPS) properties, we have used, in this work, silanized silica nanoparticles. It consists of high-purity silica, treated with hexamethyldisilazane. This treatment aimed to substitute hydroxyl groups of the pyrogenic silica with trimethylsilyl groups, which makes silica extremely hydrophobic. The silanized silica nanoparticles have been used in stickers, sealants, paints, toners and for making products for skin and beauty care[19].The aim of this work was obtaining a biodegradable and hydrophobic starch-based polymer. The addition of glycerol and of silica hydrophobic nanoparticles was evaluated for the solubility, absorption and contact angle of the polymer produced, using factorial design.

2. Materials and Methods The materials used in this study were: cassava starch, donated by INPAL S.A. Indústria Química (a Chemical Factory), located in the region of São Tomé – Paraná, Brazil; Glycerol (Nuclear, 98%) and CAB-O-SIL silanized silica nanoparticles (nSS) TS-530 (99.8 % SiO2).

The biodegradable polymers with hydrophobic characteristics were obtained from the grinding of in natura cassava starch, with different ratios of glycerol and silanized silica nanoparticles (nSS), which was later processed by extrusion, in a twin-screw extruder MiniLab II HAAKE Rheomex CTW 5. The extrusion speed was kept between 50 and 100 rpm, at a temperature of 140 °C.

213

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

Lopes, F. L. P., Kupfer, V. L., Oliveira, J. C. D., Radovanovic, E., Rinaldi, A. W., Moisés, M. P., & Favaro, S. L. In order to evaluate the influence and the interaction of nSS and glycerol for obtaining starch with hydrophobic characteristics, was done a study of a complete factorial design 22 complete with central point. The statistical data were analyzed by using a computer program, Design‑Expert.

Table 1 presents the inferior (-1) and superior (+1) levels and the central points (0) of the variables nSS(A) and glycerol (B). All the experiments were performed in duplicate and at random, which generated 12 experiments in total, as shown in Table 2. In that Table the name of each sample and the conditions used at each experiment are also mentioned.

The samples identification follows this order: TPS, glycerol percentage/silica percentage, number of the replication.

For instance, sample TPS_30G/1Si02 was prepared with 30% of glycerol, 1% of silica, and is a duplicate of sample TPS_30G/1Si 01.

2.1 Characterization The silica nanoparticles were characterized by Scanning Electron Microscopic (SEM) and Transmission Electron Microscopy (TEM). The SEM was carried out by a Quanta 250 from FEI Company (Hillsboro, OR, USA) with and accelerating voltage of 5kV. The TEM images were recorded using JEOL JEM-140 equipment operated at 120KV. The biodegradable polymers with hydrophobic characteristics were characterized by trials of water absorption, solubility and water contact angle. The scanning electron microscopy (SEM) was employed to evaluate the dispersion of the nSS particles in the polymeric material and its morphology.

The water absorption trial was carried out for the comparison of TPS with the different percentages of glycerol and nSSin relatedto the polymer mass. Firstly, the biodegradable polymers with hydrophobic characteristics were dried, until the mass became constant, by use of a desiccator with blue silica gel (4-8mm, Synth), at room temperature. After that, they were kept in a glass jar with humidity controlled at 75% (supersaturated sodium chloride aqueous solution) at room temperature, according to the standard ASTM-E-104-85. The masses of the biodegradable polymers were analyzed chronologically, until they remained constant. This way, the quantity of water absorption was calculated by Equation 1 and by the weighing of the results of three samples of each formulation.  M − MS  AbH 2O ( % )  U =  ×100  MS 

(1)

in which: AbH 2O is the water absorption of the polymer, in percentage; MU is the wet mass of the polymer in successive time intervals; M S is the dry mass of the polymer.

The analysis of solubility in water was performed in triplicate, following the method proposed by Gontard et al.[20].

The samples were dried in an oven with air circulation

and renovation for a 24-hour period, at 105 °C.After this process, the initial dry weight was obtained. The samples were immersed in 50 mL of distilled water, kept under slow and periodic stirring for a 48-hour period, at 25 °C. After this period, the remaining film fragments were removed from water and dried in an oven (105 °C, 24 hours) for determining the final dry mass. The solubility was expressed according to Equation 2, where M Si is the sample initial dry mass and M Sf is the sample final dry mass: = Solubility (%)

M Si − M Sf M Sf

×100

(2)

For the evaluation of the contact angle, films made from pellets of the plastified starch were prepared. For obtaining the films, a heated hydraulic press was used; three samples were heated for three minutes at 140 °C and, after that, they were submitted to the pressure of 1.103 kg/cm2.They were kept under that pressure condition until the polymeric film returned to room temperature. The contact angle measured with the water sessile drop method was obtained with Tantec equipment. Each value was taken as an average of five measurements in different parts of two samples produced in the same experimental conditions.

The SEM images were obtained from a microscope Shimadzu model SS 550. To this, the polymers were prepared by means of a cryogenic fracture with liquid nitrogen (-196 °C), placed on a specimen holder with adhesive tape, covered with a thin gold layer by sputtering, leaving the area where the fracture occurred exposed to be analyzed.

3. Results and Discussions 3.1 Characterization of silica Figure 1 presents the images obtained from scanning electron microscopy and from transmission electron microscopy of silanized silica nanoparticles (nSS), as received. By means of the scanning electron microscopy image, it can be seen that the nSS present agglomerations of nanoparticles with sizes varying from 5 to20 µm. These results were confirmed by transmission electron microscopy.

3.2 Starch modified with silica The responses obtained from the trials of water absorption, solubility and contact angle were used at the experimental design 22with central point (Table 2). Those responses are presented in Table 3, which also contains the data for the sample without nSS (TPS_30G) addition.

The main effects of the variables silica quantity (A) and glycerol quantity (B), as well as the combination among the factors at the values of solubility, absorption and contact angle, according to the design described in Table 3, were evaluated on the software Design-Expert. With the

Table 1. Factors and levels of the factors used at the factorial design 22 complete for hydrophobic starch.

Factor A B

Name Silica Glycerol

Unity* % %

Type Numerical Numerical

Level (-1) 1.0 30.0

Level (0) 3.0 35.0

Level (+1) 5.0 40.0

*Silica and glycerol percentage in relation to the total mass.

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Influence of addition of silanized nanosilica and glycerol on hydrophobicity of starch using a factorial design application of the experimental designing, the models, represented by Equations3, 4 and 5, were obtained for water absorption, solubility and contact angle, respectively.

In order to evaluate the reliability of the models proposed, Table 2. Complete factorial design22, study in duplicate with central point.

Samples TPS_30G/1Si 01 TPS_30G/1Si 02 TPS_30G/5Si 01 TPS_30G/5Si 02 TPS_40G/1Si 01 TPS_40G/1Si 02 TPS_40G/5Si 01 TPS_40G/5Si 02 TPS_35G/3Si 01 TPS_35G/3Si 02 TPS_35G/3Si 03 TPS_35G/3Si 04

A Silica [%] (-1) 1.0 (-1) 1.0 (+1) 5.0 (+1) 5.0 (-1) 1.0 (-1) 1.0 (+1) 5.0 (+1) 5.0 (0) 3.0 (0) 3.0 (0) 3.0 (0) 3.0

Factors B Glycerol [%] (-1) 30.0 (-1) 30.0 (-1) 30.0 (-1) 30.0 (+1) 40.0 (+1) 40.0 (+1) 40.0 (+1) 40.0 (0) 35.0 (0) 35.0 (0) 35.0 (0) 35.0

Table 3. Responses obtained at the experiments of the complete factorial design 22.

RESPONSES SAMPLES

Water absorption [%]

Solubility[%]

Contact angle

TPS_30G TPS_30G/1Si 01 TPS_30G/1Si 02 TPS_30G/5Si 01 TPS_30G/5Si 02 TPS_40G/1Si 01 TPS_40G/1Si 02 TPS_40G/5Si 01 TPS_40G/5Si 02 TPS_35G/3Si 01 TPS_35G/3Si 02 TPS_35G/3Si 03 TPS_35G/3Si 04

24.1 ± 0.1 16.9 ± 0.1 16.5 ± 0.4 17.0 ± 0.1 17.0 ± 0.4 19.8 ± 0.1 20.0 ± 0.1 19.5 ± 0.1 19.1 ± 0.1 18.5 ± 0.1 17.4 ± 0.2 18.5 ± 0.1 18.3 ± 0.1

44.1 ± 0.2 37.7 ± 1.5 33.6 ± 0.8 31.2 ± 0.7 29.7 ± 0.7 45.0 ± 0.3 40.7 ± 0.8 42.8 ± 1.5 36.2 ± 0.5 35.3 ± 0.7 38.0 ± 0.4 38.1 ± 2.2 30.2 ± 1.9

[degrees] 61.2 ± 0.1 66.5 ± 1.7 70.0 ± 1.0 91.8 ± 1.6 83.6 ± 2.4 85.4 ± 1.0 80.0 ± 2.0 89.1 ± 2.0 87.3 ± 1.0 72.5 ± 2.3 83.7 ± 1.6 81.2 ± 2.2 77.5 ± 2.5

the analysis of variance (ANOVA)was used. From ANOVA, it could be observed that all the models (Tables 4, 5 and 6) fit well to the experimental data, without presenting any significant lack of fit.

Water absorption = +6.35 − 0.75 x Silica +

0.34 x Glycerol + 0.02 x Silica x Glycerol

(3)

Solubility = + 16.77 − 2.71 x Silica + x

Glycerol + 0.04 x Silica x Glycerol

(4)

Contact Angle = +9.57 + 15.32 x Silica +

1.79 x Glycerol − 0.35 x Silica x Glycerol

(5)

By means of the analysis of variance of the factorial 22 for the data of water absorption trial presented in Table 4, it can be observed that variable B – quantity of glycerol – showed to be statistically significant (p-value<0.05).

Therefore, as the quantity of glycerol used for the starch plastification increases, the quantity of water absorbed also raises. On the other hand, variable A (quantity of nSS) and the interaction between the factors AB are not statistically significant for the model, because, by test F, they present a value p much higher than 0.05. In Figure 2, the response surface obtained for the model is presented. A significant increase in water absorption with the addition of glycerol percentage can be confirmed.

Farahnaky et al.[21] performed a study with edible polymeric films of native starch by varying the amount of glycerol. They observed that in high moisture environment up to 84%, those polymeric films have mass gain increase until 54%.

In a study carried out by Plotegher and Ribeiro[22] with TPS reinforced withZSM-5 Zeolite and colloidal silica, it could be observed that the introduction of both fillers reduced the permeability to water vapor in up to 20% when compared to the permeability of pure TPS. There are other studies that mention the use of TPS and paraffin; they show that the addition of paraffin in all the formulations reduces the water absorption[23].

Figure 1. Scanning electron microscopy (left) and transmission electron microscopy (right) of silanized silica nanoparticles (nSS) TS‑530 (SiO2) from CAB-O-SIL.

Polímeros, 27(3) , 213-219, 2017

215

Lopes, F. L. P., Kupfer, V. L., Oliveira, J. C. D., Radovanovic, E., Rinaldi, A. W., Moisés, M. P., & Favaro, S. L. Table 4. ANOVA obtained for the design of Table 3 with an adjustment for the data obtained in the water absorption trial.

Water absorption Model A (Silica) B (Glycerol) AB Pure Error Total

Sum of squares 15.57 0.045 15.12 0.40 1.01 16.59

Freedom degree 3 1 1 1 7 11

Average of squares 5.19 0.045 15.12 0.40 0.14

Value F 36.07 0.31 105.09 2.81

Value P 0.0001 0.5935 < 0.0001 0.1374

R2 = 0.9212 (percentage of prediction of model = 92.12 %).

Table 5. ANOVA obtained from the central composite design of Table 3 with an adjustment for the solubility trial.

Solubility Model A (Silica) B (Glycerol) AB Pure Error Total

Sum of squares 170.47 36.42 132.28 1.78 81.89 260.03

Freedom degree 3 1 1 1 7 11

Average of squares 56.82 36.42 132.28 1.78 11.70

Value F 4.86 3.11 11.31 0.15 0.66

Value P 0.0391 0.1210 0.0120 0.7083

R2 = 0.6556 (percentage of prediction of model = 65.56 %).

Table 6. ANOVA obtained for the design of Table 3, with an adjustment for the results of contact angle.

Contact Angle Model A (Silica) B (Glycerol) AB Pure Error Total

Sum of squares 520.30 311.25 111.75 97.30 127.07 671.18

Freedom degree 3 1 1 1 7 11

Average of squares 173.43 311.25 111.75 97.30 18.15

Value F 9.55 17.15 6.16 5.36

Value P 0.0072 0.0043 0.0421 0.0538

R2 = 0.7752 (percentage of prediction of model = 77.52 %).

Figure 2. Response surface obtained at factorial 22 presented in Table 4.

Figure 3. Response surface obtained at factorial 22 presented in Table 5.

Table 5 presents the analysis of factorial variance 22 for the solubility trial. It can be observed that variable A and the interaction between factors AB do not present a significant level for the model (p-value>0.05). In its turn, variable B is statistically significant. Figure 3 presents the obtained response surface. A significant increase of the starch solubility with the addition of glycerol percentage cane noticed.

In a study carried out by Matta et al.[24] for analyzing the solubility of biofilms obtained from peas starch with xanthan

gum and glycerol, one could observe that the addition of glycerol increased the film solubility. It occurred because glycerol interacts with the film matrix, increasing the free space between the chains; it facilitates the entrance of water, thus increasing the solubility[24-27].

The analysis of variance of factorial 22 generated from the data of the contact angle trial (Table 6) presents significant statistics, for nSS (variable A) as well as for glycerol (variable B). However, it can be observed that the quantity of nSS had high influence to the model, because,

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Influence of addition of silanized nanosilica and glycerol on hydrophobicity of starch using a factorial design by test F, a p-value much lower than 0.05 can be noticed.

On the response surface illustrated in Figure 4, it can be observed that highest values of contact angles are obtained with around 5% of nSS. This fact can be related to the presence of nonpolar group in nSS on the TPS surface.

It causes an increase in the contact angle[28,29] and a change of its hydrophobic characteristics.

The highest contact angle obtained was 91.8° for the sample TPS_30G/5Si, a value 33.3% higher when compared to the TPS without addition of nSS with 30% of glycerol (61.2°). This trend is noticed in polymers based on starch as well. The raising of glycerol content increases the hydrophilicity of these films, which might be related with the OH functional groups of the glycerol[30-33].

Figure 4. Response surface obtained at factorial 22 presented in Table 6.

Figure 5 presents the images of scanning electron microscopy of the polymers with hydrophobic characteristics on the fracture surface of the TPS samples, according to

Figure 5. Scanning electron microscopies for the TPS_30G and other formulations generated by the factorial study, according to Table 2.

Polímeros, 27(3) , 213-219, 2017

217

Lopes, F. L. P., Kupfer, V. L., Oliveira, J. C. D., Radovanovic, E., Rinaldi, A. W., Moisés, M. P., & Favaro, S. L. the nomenclature presented in Table 2. It can be observed that both polymeric materials present similar morphology, and from the cryogenic fractures, the presence of nSS nanoparticles is seen. Furthermore, it could also be observed that the nSS nanoparticles showed a very close similarity in their homogeneities comparing all percentages of polymeric materials used.

The results obtained in the factorial design can be better understood with the analysis of the morphology of the samples.

The formulations of TPS and nSS, in general, present a homogeneous appearance and a good dissolution of the starch granules after the plastification process. In Figure 5, we highlight the fact that the samples TPS_30G/1Si 01 and TPS_40G/5Si 01 present a better dispersion of nSS in the polymeric matrix. Those results can be justified: the sample TPS_30G/01Si 01 has obtained the best result in the water absorption trial (16.9 %), while the sample TPS_40G/05Si 01 has obtained a high value of contact angle. The other samples presented some areas with nSS agglomerates in the TPS matrix, which can be associated to the extrusion process.

4. Conclusion From the complete factorial design 22with central point, we could evaluate that both factors, quantity of glycerol and silanized nanosilica, influence on the process to obtain starch with hydrophobic characteristics; glycerol has a higher influence on the characteristics related to the polymer bulk, presenting higher water absorption and higher solubility with the increase of the plasticizer glycerol. On the other hand, the quantity of silanized nanosilica has more influence on the characteristics related to the surface. Therefore, there is a significant increase in the values of contact angle with addition of 5% of silanized nanosilica.

The material developed in this study has a high potential to be applied in the preparation of blends with synthetic and hydrophobic polymers, such as polyethylene and polypropylene.

5. Acknowledgements We gratefully acknowledge the Brazilian agencies (CAPES/ CNPQ) for the financial support. We would like to acknowledge to COMCAP/UEM for the electron microscopy analysis.

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24. Matta, M. D., Jr., Sarmento, S. B. S., Sarantópoulos, C. I. G. L., & Zocchi, S. S. (2011). Propriedades de barreira e solubilidade de filmes de amido de ervilha associado com goma xantana e glicerol. Polímeros: Ciência e Tecnologia, 21(1), 67-72. http:// dx.doi.org/10.1590/S0104-14282011005000011.

25. Leyva, M. B., Chávez, P. T., Wong, B. R., Jatomea, M. P., & Bojórquez, F. B. (2008). Physical and mechanical properties of durum wheat (triticum durum) starch films prepared with a and b type granules. Stärke, 60(10), 559-567. http://dx.doi. org/10.1002/star.200800227.

26. Laohakunjit, N., & Noomhorm, A. (2004). Effect of plasticizer on mechanical and barrier properties of rice starch film. Stärke, 56(8), 348-356. http://dx.doi.org/10.1002/star.200300249.

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27. Zhang, Y., & Han, J. H. (2006). Mechanical and thermal characteristics of pea starch films plasticized with monosaccharides and polyols. Journal of Food Science, 71(2), E109-E118. http:// dx.doi.org/10.1111/j.1365-2621.2006.tb08891.x.

28. Garcia, M. A., Pinotti, A., & Zaritzky, N. E. (2006). Physicochemical water vapor barrier and mechanical properties of corn starch and chitosan composite films. Stärke, 58(9), 453-463. http:// dx.doi.org/10.1002/star.200500484.

29. Thiré, R. M. S. M., Simão, R. A., & Andrade, C. T. (2003).

High resolution imaging of the microstructure of maize starch films. Carbohydrate Polymers, 54(2), 149-158. http://dx.doi. org/10.1016/S0144-8617(03)00167-X. 30. Müller, C. M. O., Yamashita, F., & Laurindo, J. B. (2008).

Evaluation of the effects of glycerol and sorbitol concentration and water activity on the water barrier properties of cassava starch films through a solubility approach. Carbohydrate Polymers, 72(1), 82-87. http://dx.doi.org/10.1016/j.carbpol.2007.07.026.

31. Olivato, J. B., Grossmann, M. V. E., Yamashita, F., Eiras, D., & Pessan, L. A. (2012). Citric acid and maleic anhydride as compatibilizers in starch/poly(butylene adipate-co-terephthalate) blends by one-step reactive extrusion. Carbohydrate Polymers, 87(4), 2614-2618. http://dx.doi.org/10.1016/j.carbpol.2011.11.035.

32. Olivato, J. B., Grossmann, M. V. E., Bilck, A. P., & Yamashita, F. (2012). Effect of organic acids as additives on the performance of thermoplastic starch/polyester blown films. Carbohydrate Polymers, 90(1), 159-164. PMid:24751025. http://dx.doi. org/10.1016/j.carbpol.2012.05.009.

33. Nobrega, M. M., Olivato, J. B., Müller, C. M. O., & Yamashita, F. (2012). Biodegradable starch-based films containing saturated fatty acids: thermal, infrared and Raman spectroscopic characterization. Polímeros: Ciência e Tecnologia, 22(5), 475480. http://dx.doi.org/10.1590/S0104-14282012005000068.

Received: Apr. 04, 2016 Revised: Dec. 02, 2016 Accepted: Dec. 20, 2016

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http://dx.doi.org/10.1590/0104-1428.04016

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Cowper-Symonds parameters estimation for ABS material using design of experiments with finite element simulation Alexandre Luis Marangoni1* and Ernesto Massaroppi Junior2** Whirlpool S.A., Rio Claro, SP, Brazil Laboratório de Mecânica Aplicada e Computacional, Departamento de Engenharia Mecânica, Escola de Engenharia de São Carlos – EESC, Universidade de São Paulo – USP, São Carlos, SP, Brazil 1

2

*alexandre.marangoni@usp.br, **massarop@sc.usp.br

Abstract Polymers exhibit significant strain rate dependence in their mechanical strength. The impact simulations accuracy is associated with the use of mechanical properties obtained at high strain rates. These properties are often not available to engineers introducing a risk on the product development step. This paper presents a method for adjusting the parameters of the Cowper-Symonds, used for a constitutive material model, through computational experiments carried out considering the simulation of the Izod impact test.The proposed adjustment method allows reducing the Izod impact strength error from 44% to 2.4%.

Keywords: cowper-symonds, finite element analysis, izod, space filling design, strain rate.

1. Introduction The structural strength of certain products and its components can not be easily evaluated by analytical calculations. For these cases the use of computational simulations performed with the Finite Element Method (FEM) is a frequently used tool for project aid, which generates significant gains in development time and prototypes cost reduction. The FEM is used to obtain a numerical solution of partial diferential equations in an approximate and discretized way. The vality domain of the diferential equation is subdivided in several subdomains, named elements, that are described by characteristics points, named nodes, normally positioned on the vertex of a polygon or a polyhedron.

To each node is associated a polynomial shape function, whose linear combination is adopted as the solution in the element subdomain (approximate). The shape functions are fixed for a gived geometry and the solution functions become only dependent of the nodal value of the interesting variable (discretized)[1]. Furthermore, compatibility conditions are imposed to the solution functions in the interfaces between the elements. Among the possible diferential equations, those which relate the stress and strain in solid bodies submitted to forces, are commercially interesting and implemented in FEM softwares as, for exemple, LS-DYNA[2]. Before starting the FEM solver, it is necessary to define the domain, which is normally presented as a geometric description in a Computer Aided Design (CAD) software, and to generate the FEM mesh in an appropriated software as Altair HyperMesh 13.0[3]. The simulation results are graphically represented in a post processing software as Altair HyperView 13.0[4].

The accuracy of the simulations results is strongly related to the finite element model quality, which includes the choice of the formulations of the elements that will be used to represent the structure, the interpretation of the boundary conditions, the constitutive model of the material and the solution method being used.

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Impact simulations are performed to quantify the behavior of a mechanical structure under impact loads. The lack of information regarding material properties could lead to erroneous results and, as a consequence, wrong conclusions regarding the impact strength of the simulated part or product.

Polymers are widely used in the industry due to its good balance between cost and mechanical characteristics.

Its mechanical strength varies as a function of the applied strain rate, which is associated with the polymer´s viscoelastic behavior[5]. The well-known Cowper-Symonds Equation 1 is frequently used to describe the material behavior at different strain rates[6].

1   ε  P   sD = sE. 1 +    (1) C   

Where sD: dynamic stress; sE: quasi-static stress; C e P: Cowper-Symonds equation´s parameters and; ε : strain rate.

It is possible to find the Cowper-Symonds parameters with correlations above 0.95. However, these parameters assume different values for different reference stresses, because each stress has its own sensitivity to strain rate, leading to different parameters in Cowper-Symonds equation.

This feature can have an important effect on the numerical simulation of dynamic processes[7,8].

The LS-DYNA has several materials constitutive models and some of them use the Cowper-Symonds equation to include the strain rate effect at the yield stress.

Năstăsescu and Iliescu[9] obtained values for the Cowper-Symonds parameters of a polymer through finite element simulations of the Izod impact test.

The mechanical properties of the material at high strain rates could be obtained experimentally by the use of

Polímeros, 27(3) , 220-224, 2017

Cowper-Symonds parameters estimation for ABS material using design of experiments with finite element simulation a servo-hydraulic tensile testing machine or a Hopkinson bar testing[10].

Computational experiments could be generated using space filling method, where the desired respose is evaluated under several configurations, obtained from an organized combination of the design factors levels. As the result a metamodel is generated to predict a system behavior as function of the studied factors[11,12].

The objective of this work is to use design of experiments techniques to find and to adjust the Cowper-Symonds parameters of a LS-DYNA material model, which can be used to improve the accuracy of the product level simulations.

2. Materials and Methods The notched Izod impact test specimen, the impact hammer and the fixation jaw finite element model, showed at Figure 1, were created using the Altair HyperMesh 13.0[3] software in accordance to the ISO 180[13] standard.

The LS-DYNA *MAT_089 (Plasticity Polymer) material model was set to the specimen because the true stress-strain curve is used as input data, allowing the representation of the non-linear elastic behavior of most polymers. The strain rate effect could be applied at the yield or ultimate stress by the use of the C and P parameters of the Cowper-Symonds

Figure 1. Notched Izod Impact test FE model: (a) fixation jaw (green), specimen (blue) and impact hammer (gray); (b) Detail of the specimen mesh at the notch.

Table 1. ABS mechanical properties.

Mechanical Property Young´s Modulus Yield Stress (23°C) Yield Strain (23°C) Izod Impact Strengh

Value 2317.65 44.05 3.34 21.691

Unit MPa MPa % kJ/m2

Source Stress-strain curve Stress-strain curve Stress-strain curve Datasheet

Figure 2. ABS stress-strain curve. Source: Campus[14].

Polímeros, 27(3) , 220-224, 2017

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Marangoni, A. L., & Massaroppi, E., Jr. equation or as a load curve at the failure strain by the use of a natural log of the strain rate by the true strain to failure.

The absorbed energy by the specimen was calculated through Equation 4.

The mechanical properties and the stress-strain curve of an ABS (Acrylonitrile butadiene styrene) grade used on this study were obtained from the CAMPUS[14] (Computer Aided Material Preselection by Uniform Standards) database, and are showed at Table 1 and Figure 2, respectively.

1 ∆EK = EK − m × v 2f (3) 2

For the material model *MAT_89 an internal check is performed by LS-DYNA and the yield occurs when the current stress-strain curve slope becomes lower than the value defined at the elastic model field[15].

The rigid material (*Mat_020) was used at the jaw and impact hammer. The material density was calculated in order to result into a 15 kg impact hammer mass (m). The ABS Izod absorbed energy is 2.0 J which correspond to an impact energy (EK) of 2.7 J. The impact hammer initial velocity (v) of 600 mm/s was calculated as shown in Equation 2. = v 1000 × 2 EK / m (2)

A space filling design experiment, considering as factors the parameters C and P of the Cowper-Symonds equation, was carried out to adjust these material properties due to its capability for modeling non-linear phenomena.

The factors levels were estimated through simulations considering arbitrary values to C and P. The impact test simulation was also performed without any strain rate effect.

The computational experiment was generated using the data analysis software SAS JMP 10.0[11], considering the uniform design and 10 simulations, as showed on Table 2.

Although the experimental generation uses the range of values for the factors C and P, here they will be presented at normalized form between -1 and 1 for industrial confidence purposes.

The Altair HyperMesh 13.0 software was used to generate 10 finite element model files considering the combination of the C and P factors. The simulations were performed at LS-DYNA considering an impact time of 0.015 s. The final hammer velocity value obtained at the post-processing, performed by Altair HyperView 13.0 software, was used at Equation 3 to calculate the kinetic energy variation.

Table 2. Normalized parameters for each simulation of the space filling experiment.

Simulation 1 2 3 4 5 6 7 8 9 10

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C -0.703 0.310 -0.898 0.897 0.507 -0.288 0.121 -0.115 0.704 -0.512

P 0.836 0.370 0.637 0.701 0.775 0.432 0.966 0.901 0.563 0.502

Where ΔEK: kinetic energy variation; EK: initial kinetic energy of the hammer; m: hammer mass and; vF: final velocity of the hammer.

Eabs = ∆EK / A fr (4)

Where EABS: specific absorved energy; ΔEK: kinetic energy variation; and AFR: fractured area of the specimen.

The software SAS JMP 10.0 was also used to analyze the experiment results. A metamodel was adjusted by the Gaussian method. The desirability feature was used to reach an absorbed energy of 21.691 kJ/m2 through the combination of the values of the factors C and P. The obtained combination of the factors was then used in another simulation to check the metamodel adjustment quality/accuracy.

3. Results and Discussions The final hammer velocity obtained at the simulation without the strain rate effects was 554.95 mm/s. The absorbed energy by the specimen was 39.026 kJ/m2.

A fractured specimen and the respective plot of the hammer velocity as function of the impact time are shown on Figures 3a and 3b, respectively. The hammer velocity was reduced from 600 mm/s to 517.758 mm/s. Despite an incomplete fracture of the specimen was observed at impact time of 0.015 s, the asymptotic behavior of the curve after 0.014 s indicates that the hammer velocity will not be significantly changed by the use of a higher simulation termination time. Table 3 shows the absorbed energy calculated through Equation 4 for the all simulations of the experiment.

Figure 4 shows the absorbed impact energy obtained by the simulation as a function of the predicted results by the metamodel. The quality of the metamodel can be verified by the proximity of the points to the dashed red line. For this case, as predicted, only one point showed a significant distance from the dashed line. Although the computational experiment was carried out with a reduced number of simulations, the resultant metamodel has shown a satisfactory quality of fit.

The prediction energy model as a function of the C and P parameters behavior are shown on Figure 5.

The normalized values of -0.378 and 0.5107 for the C and P parameters, respectively, were obtained through the use of the SAS JMP 10.0 Desirability function used to match the target energy of 21.691 kJ/m2. The impact simulation result for this case is shown on Figure 6. A final velocity of 515.016 mm/s of the impact hammer resulted on an absorbed energy of 22.209 kJ/m2, which is 2.4% higher than the target impact value as shown at the polymer datasheet.

Polímeros, 27(3) , 220-224, 2017

Cowper-Symonds parameters estimation for ABS material using design of experiments with finite element simulation

Figure 3. Simulation number 1: (a) Fractured virtual specimen; (b) Impact hammer velocity (mm/s) as function of the time (s).

Table 3. Final impact hammers velocities and absorbed energy by the specimen for all simulations.

Simulation number 1 2 3 4 5 6 7 8 9 10

Final Hammer Velocity (mm/s) 517.758 517.171 516.591 518.917 518.936 516.562 519.322 519.004 519.831 516.779

Final Hammer Kinetic Energy (J) 2.011 2.006 2.001 2.020 2.020 2.001 2.023 2.020 2.027 2.003

Absorbed Energy (kJ/m2) 21.545 21.688 21.828 21.264 21.259 21.835 21.165 21.243 21.041 21.783

Figure 4. Absorbed energy values (Eabs) as function of the predicted values by the Gaussian Modeling.

Figure 5. Prediction of specimen energy absorption as function of the normalized C and P parameters.

PolĂ­meros, 27(3)â&#x20AC;&#x160;, 220-224, 2017

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Marangoni, A. L., & Massaroppi, E., Jr.

Figure 6. Results of the validation simulation: (a) Fractured Specimen; (b) Impact hammer velocity (mm/s) as function of the time (s).

4. Conclusions While the execution of an Izod impact test simulation without considering any material strain rate effect led to a 44% error relative to the Izod result shown at polymer datasheet, the absorbed energy error was significantly reduced to 2.4% with the use of the adjusted Cowper-Symonds strain rate parameters.

Based on this result it can be concluded that the adjustment method of Cowper-Symonds equation parameters for the LS-DYNA *MAT_089 (Plasticity Polymer) material model through the use of a computational space filling design experiment was effective. This methodology could improve the simulation results in cases where the complete dynamic material characterization is not available.

5. Acknowledgements The authors are grateful to Whirpool S.A. and to Laboratory of Applied and Computational Mechanical of the Departament of Mechanical Engineering of the School of Engineering of São Carlos of the University of São Paulo for providing the necessary computational resources.

6. References 1. Zienkiewicz, O. C., & Taylor, R. L. (2000). The finite element method (Vol. 1, 5th ed., 689 p.). Oxford: Butterworth Heinemann.

2. Livermore Software Technology Corporation. (2014). LS-DYNA keyword user’s manual, revision 5471 (Vol. 1). Livermore: LSTC. 3. Altair. (2014). HyperMesh 13.0 manual. Retrieved in 1 April 2016, from http://www.altairhyperworks.com/product/ HyperMesh 4. Altair. (2014). HyperView 13.0 manual. Retrieved in 1 April 2016, from http://www.altairhyperworks.com/product/ HyperView 5. Xiao, X. (2008). Dynamic tensile testing of plastic materials.

Polymer Testing, 27(2), 164-178. http://dx.doi.org/10.1016/j. polymertesting.2007.09.010.

6. Peixinho, N., & Doellinger, C. (2010). Characterization of dynamic material properties of light alloys for crashworthiness applications. Materials Research, 13(4), 471-474. http://dx.doi. org/10.1590/S1516-14392010000400008.

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7. Peixinho, N., & Pinho, A. (2007). Study of viscoplasticity models for the impact behavior of high-strengh steels. Journal of Computational and Nonlinear Dynamics, 2(2), 114-123. http://dx.doi.org/10.1115/1.2447129.

8. Alves, M. (2000). Material constitutive law for large strains and strain rates. Journal of Engineering Mechanics, 126(2), 215-218. http://dx.doi.org/10.1061/(ASCE)0733-9399(2000)126:2(215).

9. Năstăsescu, V., & Iliescu, N. (2010). Upon accompanying of the experimental testing of materials by numerical analysis with FEM. Acta Technica Napocensis - Applied Mathematics and Mechanics, 2(53), 173-178. Retrieved in 1 April 2016, from artens2010.utcluj.ro/acta%20tehnica%20nr53%20 vol2%202010/10.doc 10. Zrida, M., Laurent, H., Grolleau, V., Rio, G., Khlif, M., Guines, D., Masmoudi, N., & Bradai, C. (2010). High-speed tensile tests on a polypropylene material. Polymer Testing, 29(6), 685-692. http://dx.doi.org/10.1016/j.polymertesting.2010.05.007.

11. SAS Institute. (2010). User Guide - SAS JMP 10.0. Cary: SAS Institute.

12. Baco, S. B. (2016). Uso de experimentos computacionais no desenvolvimento de produtos: um estudo de caso na indústria de linha branca (Dissertação de mestrado). Universidade Federal de São Carlos, São Carlos.

13. International Standard Organization. (2000). ISO 180:2000 - Plastics -- Determination of Izod impact strength. Geneva: ISO. 14. Computer Aided Material Preselection by Uniform Standard.

(2016). Plastics Material Database. Frankfurt: Campus.

Retrieved in 1 April 2016, from http://www.campusplastics. com 15. Lobo, H., & Croop, B. (2009). A robust methodology to calibrate crash material models for polymers. In NAFEMS World Congress (14 p). Ithaca: DatapointLabs. Retrieved in 1 April 2016, from http://www.datapointlabs.com/testpaks/2009/ NAFEMS09.pdf Received: Apr. 05, 2016 Revised: Dec. 01, 2016 Accepted: Mar. 01, 2017 Polímeros, 27(3) , 220-224, 2017

http://dx.doi.org/10.1590/0104-1428.13216

Chemical resistance of core-shell particles (PS/PMMA) polymerized by seeded suspension Luiz Fernando Belchior Ribeiro1, Odinei Hess Gonçalves2, Cintia Marangoni3, Günter Motz4 and Ricardo Antonio Francisco Machado1* Materials Engineering Department, Universidade Federal de Santa Catarina – UFSC, Florianópolis, SC, Brazil 2 Post-graduation Program in Food Technology, Universidade Técnológica Federal do Paraná – UTFPR, Campo Mourão, PR, Brazil 3 Universidade Federal de Santa Catarina – UFSC, Blumenau, SC, Brazil 4 Lehrstuhl Keramische Werkstoffe, Universität Bayreuth, Bayreuth, Germany 1

*ricardo.machado@ufsc.br

Abstract Core-shell particles were produced on seeded suspension polymerization by using polystyrene (PS) as polymer core, or seed, and methyl methacrylate (MMA) as the shell forming monomer. Two synthesis routes were evaluated by varying the PS seed conversion before MMA addition. The main purpose of this work was to investigate the influence of synthesis routes on the morphology and chemical resistance of the resulting particles. 1H NMR spectroscopy showed that the use of PS seeds with lower conversion led to the formation of higher amount of poly(styrene-co-MMA). The copolymer acted as a compatibilizer, decreasing the interfacial energy between both homopolymers. As a consequence, a larger amount of reduced PMMA cluster were formed, as was revealed by TEM measurements. Samples in this system showed enhanced resistance to cyclohexane attack compared with pure PS, with a PS extraction of only 37% after 54 hours test.

Keywords: seeded suspension polymerization, poly(methyl methacrylate), polystyrene, suspension polymerization.

1. Introduction Styrene suspension polymerization has been extensively studied over the past decades[1-5]. It consists in polymerizing relatively water-insoluble monomer droplets in a water media formed by vigorous stirring in the presence of a steric stabilizer. The synthesis leads to the formation of spherical polymer particles dispersed in an aqueous system.

The main use of polystyrene (PS) produced by suspension is the production of expandable polystyrene (EPS). EPS foams possess excellent properties, which combine low density with high mechanical resistance. Therefore, EPS foams are used in a wide range of applications, especially in the construction and packaging industries. The range of applications of these foams would be broader by improving some of the EPS properties, in particular their chemical resistance.

Low chemical resistance of PS is a limiting factor for many applications, such as the transport of heavy equipment containing lubricant greases. In this case, the EPS foam needs to be covered with a chemically resistant polymer film, which introduces a subsequent manufacturing step, increases the cost of production and makes the use of it near prohibitive.

The use of poly(methyl methacrylate) (PMMA) foam as a substitute to EPS is reported elsewhere[6]. PMMA chemical resistance is much higher than PS[7] and despite the promising results, the application is limited by the higher cost of monomers.

Polímeros, 27(3) , 225-229, 2017

An interesting way to overcome this problem is the synthesis of core-shell particles using lower amounts of PMMA to cover PS particles. Gonçalves et al.[8] investigated the morphology of core-shell particles composed of PS core and PMMA shell. The synthesis consisted in feeding methyl methacrylate (MMA) monomers with the initiator at low temperature (50 °C) to swell previously obtained PS seeds.

In a subsequent step, the temperature reaction is increased up to 70 °C in order to polymerize the MMA forming shell.

The obtained particles presented a complex morphology with a PS core and a shell composed of PMMA domains distributed in the PS matrix. Further works[9] evaluated the expansion of the aforementioned synthesized core-shell particles. The results revealed that the expansion needs higher temperatures than the conventional EPS processing.

Moreover, the presence of PMMA at the surface seems to hinder particle expansion, and a non-uniform foamed particle was obtained. This result is in agreement with a more recent study published by Heydarpoor et al.[10].

The authors investigated similar systems of PS/PMMA core-shell expandable particles, by using pentane as the expansion agent. They concluded that pentane was more concentrated in the PS core than in the PMMA domains, which was responsible for a non-uniform expansion behavior.

In this work a similar system used by Gonçalves et al.[8] was employed but the PS seed was polymerized in situ.

The use of in situ polymerized PS seeds allows to control the polymer conversion before MMA addition, which could be used to change particle morphology by improving the

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

Ribeiro, L. F. B., Gonçalves, O. H., Marangoni, C., Motz, G., & Machado, R. A. F. MMA diffusion inside the seeds. Two different synthesis route were evaluated and the resulting particles were investigated in respect to their morphology and chemical resistance.

2. Materials and Methods 2.1 Materials Technical grade styrene (Innova S.A.) and methyl methacrylate (Rohm&Hass S.A.) were used as monomers.

Benzoyl peroxide (BPO, Sigma-Aldrich) was used as an initiator. Distilled water was used as continuous phase and poly(vinyl pyrrolidone) (PVP K90, Sigma-Aldrich) as stabilizer. Ascorbic acid (Sigma-Aldrich) was used to avoid inhibition caused by oxygen, and sodium chloride PA (Vetec) was used to decrease methyl methacrylate solubility in water. Chloroform-d (Sigma-Aldrich) was used to solubilize samples for 1H NMR investigations and cyclohexane (Vetec) was used for the chemical resistance tests. All chemicals were used as received without further purifications.

2.2 Core-shell synthesis The synthesis of the core-shell particles was carried out in a 5 L jacketed reactor fitted with stirring rate and temperature control. The synthesis consisted in a two-step reaction and the formulation employed is shown in detail at Table 1. The first step consisted in synthesizing the PS seed using conventional suspension polymerization.

Water, styrene and BPO were added in the reactor at room temperature and the system was heated up to 90°C and stirred at this temperature until the determined reaction time. PVP was added only 1h after the reaction media reached 90°C. In the second step, ascorbic acid and sodium chloride were added in the reactor and the system was cooled (50°C).

The additional amount of BPO was dissolved in the MMA monomer and feed in the reactor at a constant rate (5 g/min).

After the monomer/initiator feed completion, the system was heated up to 70°C and allowed to react for 4 h. The period of time in which the system was kept at low temperature before reaching 70°C was arbitrarily defined as swelling time. During this period, no significant amount of PMMA is expected to be formed.

Two synthesis routes were evaluated by varying the PS seed conversion before MMA addition. The PS seed conversion was determined by reaction time, 7 hours for synthesis route 1 and 5 hours for synthesis route 2. The choice Table 1. Formulation used for core-shell particles synthesis.

Reactant 1ª Step – PS seed synthesis Water Styrene PVP BPO 2ª Step- Core-shell synthesis MMA Ascorbic acid Sodium cloride BPO

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Weight (g) 3624.0 969.0 130.0 5.3 302.0 5.4 181.2 1.7

for reaction times was based on a mathematical model of styrene homopolymerization to obtain PS conversions near to 100 and 90%[11]. The conversion was analyzed by gas chromatography and the values founded were 98% for synthesis route 1 and 92% for synthesis route 2. The temperature profiles used in these reactions are represented in Figures 1a and 1b.

2.3 Characterizations Styrene conversion was investigated by gas chromatography (GC) analysis. A sample of the polystyrene seed was taken from the reaction medium before the addition of MMA and solubilized in chloroform. An aliquot of 20 µl was taken from the solution, added in a sealed vial and analyzed with a Shimadzu GC-2010AF (Shimadzu Corporation, Japan) equipped with a headspace auto sampler (Shimadzu AOC‑5000), a flame ionization detector at 220°C and a Restek 30m RTX-5 column. A calibration curve was previously performed to determine the amount of unreacted styrene.

The composition and structural characterization of core‑shell particles was investigated by 1H NMR spectroscopy.

The samples were solubilized in chloroform-d and the NMR‑spectra of the polymer solutions were collected on a Bruker AC200 (200 MHz) spectrometer. The molar percentage of PMMA was calculated using Equation 1.

A/3   %PMMA =   *100 A / 3 + B/5 

(1)

where A is the integrated area related to CH3-O peaks in PMMA and B is the integrated area related to different aromatic protons in PS. Transmission electron microscopy (JEM-1011 TEM at 100 kV) was used to investigate the morphology of the particles. The particles were sliced using an ultramicrotome and deposited on a copper grid. The samples were stained by ruthenium tetroxide (RuO4) vapor for 1 h to reveal the phases.

Chemical resistance was tested by particle dissolution in cyclohexane, a selective solvent of PS. For sampling, the agitation was ceased and a liquid phase sample was collected. The collected samples were dried in an oven and the amount of PS extracted was determined gravimetrically.

For all characterization technique employed particles with diameter size in the range of 1.0-1.18 mm were selected.

3. Results and Discussions 3.1 NMR measurements The 1H NMR spectrum (Figure 2) shows the specific signals for PS and PMMA homopolymers for both systems.

The resonances at 6.2-7.5 ppm are due to the different aromatic protons in PS while the resonances at 3.2-3.9 ppm are characteristic to the O-CH3 group of PMMA. In addition to typical peaks from PS and PMMA homopolymers, a new peak appears at 2.9 ppm. This signal is commonly assigned to O-CH3 sites bonded to PS sequences[10,12], which proves that a copolymerization reaction is taking place. By using the integrated area of the peak centered at 2.9, the percent of copolymer formed for synthesis route 1 and 2 were determined to be 1 and 6 mol% respectively. The higher Polímeros, 27(3) , 225-229, 2017

Chemical resistance of core-shell particles (PS/PMMA) polymerized by seeded suspension

Figure 1. Temperature profile used in the experiment of (a) synthesis route 1 and (b) synthesis route 2.

Figure 2. 1H NMR spectrum of the synthesized core-shell particles.

amount of poly(styrene-co-MMA) formed in synthesis route 2 particles is due to the lower PS seed conversion in this system, which contains a higher amount of residual styrene to react with MMA. The PMMA molar percentages in core-shell particles were calculated as described in the experimental procedure and the results are shown in Table 2. Incorporation efficiency was defined as the ratio of the MMA fed to the reactor and the MMA effectively incorporated into the particles.

There was no significant difference in PMMA incorporation.

Both systems presented high incorporation efficiency. For the systems under study, swelling time seems to rule the MMA incorporation in the PS seeds. However, as will be further discussed, morphology was drastically influenced by the conversion of the PS seed.

3.2 TEM analysis TEM micrographs of the core-shell particles are presented in Figure 3a and 3b. On these images, PMMA appears as light gray and PS as dark gray. The cross section Polímeros, 27(3) , 225-229, 2017

Table 2. Concentration of PMMA and incorporation efficiency of core-shell particles.

Synthesis Route

PMMA (mol %)

1 2

18 20

Incorporation efficiency (%) 74 82

representation shows the approximated region where the images were obtained.

The TEM measurements reveal that core-shell morphology consisted of PMMA clusters dispersed in the PS matrix and the size and concentration of these clusters decrease along the radius. Previous studies showed that this kind of morphology was expected[8,10]. This occurred because the polymerization rate is much faster than the diffusion of PMMA/MMA in the swollen PS particles.

When comparing the TEM images of the synthesis route 1 and 2 particles, it was noteworthy the influence of PS seed conversion in the distribution and size of the PMMA clusters. In the system using a lower PS seed conversion it 227

Ribeiro, L. F. B., Gonçalves, O. H., Marangoni, C., Motz, G., & Machado, R. A. F.

Figure 4. Polystyrene extraction by dissolution in cyclohexane.

Figure 3. TEM images and the approximated representation of the cross section of core-shell particles from (a) synthesis route 1 and (b) synthesis route 2.

The core-shell morphology for both reaction systems improved cyclohexane resistance in comparison to pure PS. After 8 h tests, the pure PS particles were completely solubilized. In contrast, for core-shell particles derived from synthesis route 1 and 2, only 55 and 19% of PS was extracted respectively. As cyclohexane is a non-solvent to PMMA, its presence at the particle surface hinder solvent diffusion, thus retarding PS core extraction. Even after long period test, the PS were not completely extracted in core-shell particles, which reflect the improvement in chemical resistance as the benefit of the presence of PMMA clusters at the particle surface. The remarkable result for particles from synthesis route 2 is clearly justified by their morphology, which contains more PMMA clusters densely grouped at the surface.

4. Conclusions was possible to identify the presence of PMMA domains closer to the center of the particle (Figure 4b). By contrast, for the particles from the system employing higher PS seed conversion, the PMMA clusters are more concentrated at the surface. These differences can be explained in part by the higher diffusion rate of MMA in particles from synthesis route 2. In this system, the higher amount of residual styrene present in PS particles act as a solvent. As a result, intermolecular forces between the PS chains are decreased, which improve MMA diffusion and allows the nucleation of PMMA clusters closer to the center of the particle. In addition, the formation of a poly(styrene-co-MMA), confirmed by 1 H NMR measurements, reduced the PS/PMMA interfacial energy. The presence of a higher amount of copolymers in particles from synthesis route 2 act as a compatibilizer, and it was responsible for the formation of a higher amount of reduced size clusters.

In this work, core-shell particles were produced through seeded suspension polymerization by using polystyrene (PS) as a seed, and methyl methacrylate (MMA) as the shell forming monomer. The core-shell particles presented a complex morphology with a PS core and a shell composed

by PMMA clusters distributed in the PS matrix. The size and concentration of these clusters along the particle radius were highly influenced by the styrene conversion. The use of PS seed with a lower conversion promoted MMA diffusion inside the PS seeds, and PMMA clusters were observed near the center of the particle. Moreover, the formation of poly(styrene-co-MMA) decreased the PS/PMMA interfacial energy, which led to formation of smaller PMMA clusters.

Finally, the synthesized core-shell particles presented enhanced chemical resistance compared to PS. The presence of PMMA domains grouped at the particles surface prevents the completely PS extraction, even after long period tests.

3.3 Chemical resistance tests

5. Acknowledgements

The chemical resistance results are shown in Figure 4.

The percentage of polystyrene extracted by the solvent (cyclohexane) was determined gravimetrically in relation with the total of PS present in the particles. All the measurements were made in triplicate.

The authors would lilke to thank CAPES-PROBRAL and CAPES-PVE (Coordination for the Improvement of Higher Education Personnel) and CNPq (National Council for Scientificand Technological Development) for their financial support. We are also thankful LCME

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Chemical resistance of core-shell particles (PS/PMMA) polymerized by seeded suspension – UFSC (Electron Microscope Laboratory) for the TEM measurements.

6. References 1. Vivaldo-Lima, E., Wood, P. E., Hamielec, A. E., & Penlidis, A. (1997). An updated review on suspension polymerization.

Industrial & Engineering Chemistry Research, 36(4), 939-965. http://dx.doi.org/10.1021/ie960361g.

2. Dowding, P. J., & Vincent, B. (2000). Suspension polymerisation to form polymer beads. Colloids and Surfaces. A, Physicochemical and Engineering Aspects, 161(2), 259-269. http://dx.doi. org/10.1016/S0927-7757(99)00375-1.

3. Scheirs, J., & Priddy, D. (2003). Modern styrenic polymers: polystyrenes and styrenic copolymers. Chichester: John Wiley & Sons, Ltd. 4. Brooks, B. W. (2005). Free-radical polymerization: suspension.

In T. H. Meyer & J. Keurentjes (Eds.), Handbook of Polymer Reaction Engineering (pp. 213-247). Weinheim: Wiley-VCH Verlag GmbH. 5. Kotoulas, C., & Kiparissides, C. (2008). Suspension polymerization.

In J. Asua (Ed.), Polymer Reaction Engineering (pp. 209-232).

Oxford: Blackwell Publishing Ltd. 6. Gonçalves, O. H., Staudt, T., Araújo, P. H. H., & Machado, R. A. F. (2009). Foaming of poly(methyl methacrylate) particles.

Materials Science and Engineering C, 29(2), 479-484. http:// dx.doi.org/10.1016/j.msec.2008.08.034.

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7. Brandrup, J., Immergut, E. H., & Grulke, E. A. (1999). Polymer handbook. New York: John Wiley & Sons, Inc. 8. Gonçalves, O. H., Asua, J. M., Araújo, P. H. H., & Machado, R. A. F. (2008). Synthesis of PS/PMMA core−shell structured particles by seeded suspension polymerization. Macromolecules, 41(19), 6960-6964. http://dx.doi.org/10.1021/ma800693m.

9. Gonçalves, O. H., Leimann, F. V., Araújo, P. H. H., & Machado, R. A. F. (2013). Expansion of core-shell PS/PMMA particles.

Journal of Applied Polymer Science, 130(6), 4521-4527. http:// dx.doi.org/10.1002/app.39731.

10. Heydarpoor, S., Abbasi, F., Jalili, K., & Najafpour, M. (2015).

Synthesis of core-shell PS/PMMA expandable particles via seeded suspension polymerization. Journal of Polymer Research, 22(8), 151. http://dx.doi.org/10.1007/s10965-015-0789-0.

11. Machado, R. A. F., & Bolzan, A. (1998). Control of batch suspension polymerization reactor. Chemical Engineering Journal, 70(1), 1-8. http://dx.doi.org/10.1016/S1385-8947(98)00006-0.

12. Songkhla, P. N., & Wootthikanokkhan, J. (2002). Effect of the copolymer composition on the K and a constants of the Mark-Houwink equation: Styrene-methyl methacrylate random copolymers. Journal of Polymer Science. Part B, Polymer Physics, 40(6), 562-571. http://dx.doi.org/10.1002/polb.10119.

Received: Oct. 25, 2016 Revised: Feb. 13, 2017 Accepted: Mar. 02, 2017

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http://dx.doi.org/10.1590/0104-1428.02316

Evaluation of degree of conversion, microtensile bond O strength and mechanical properties of three etch-and-rinse O dental adhesives O Samantha Ariadne Alves de Freitas , Marco Daniel Septimo Lanza , Karina Kato Carneiro , Alessandro Dourado Loguercio and José Bauer * O Disciplina de Materiais Dentários, Curso de Odontologia, Universidade Federal do Maranhão – UFMA, O São Luís, MA, Brazil Departamento de Dentística Restauradora, Faculdade de Odontologia, Universidade Federal de Minas O Gerais – UFMG, Belo Horizonte, MG, Brazil Curso de Odontologia, Departamento de Dentística Restauradora, Universidade Ceuma – Uniceuma, O São Luís, MA, Brazil Departamento de Dentística Restauradora, Faculdade de Odontologia, Universidade Estadual de Ponta O Grossa – UEPG, Ponta Grossa, PR, Brazil O O Abstract This study evaluated microtensile bond strength (µTBS), degree of conversion, modulus of elasticity and ultramicrohardness etch-and-rinse adhesives systems. The materials evaluated were: Ambar (FGM), Optibond (Kerr) and Magic O ofBondthree(Vigodent).

The degree of conversion was analyzed by FTIR/ATR. To evaluate bond strength (μTBS) in dentin, 15 teeth (n = 5) were restored and sliced to obtain the specimens (0.8mm ). The dynamic ultra microhardness tester was O used to evaluate the hardness and modulus of elasticity. The Magic Bond adhesive system showed lower µTBS than and Optibond (p <0.001). For degree of conversion, comparisons between groups of adhesive systems evaluated O Ambar showed statistically significant difference (p<0.001), with higher values for Ambar and Optibond when compared a Bond. For modulus of elasticity and ultramicrohardness, Ambar and Magic Bond showed lower values than O Magic Optibond. The best results in all properties evaluated were obtained by the Optibond adhesive system.

O Keywords: dental materials, adhesive, mechanical properties, microtensile bond strength, degree conversion.

O 1. Introduction 1

2

4

3

1

1

2

3

4

*bauer@ufma.br

2

Based on the management of the smear layer substrate contemporary adhesive systems are categorized as etch‑and-rinse and self-etch systems. When using etchand-rinse adhesive systems, the surface and subsurface mineral components of dentin are removed totally by acid etching[1-3]. When the conditioning step is followed by a priming step and subsequent application of the adhesive resin, they are recognized and available as three-step adhesive systems. On the other hand when the primer and adhesive resin are combined into one application and presented as two-step procedures that reduce the number of clinical steps[1].

In this single bottle there is a mixture of hydrophilic, hydrophobic monomers and solvent[3-5]. The presence of solvents favors the penetration of monomers into the collagen network and contributes to removing water[3]. Vapor pressure is an important feature to ensure good solvent evaporation after applying the adhesive on dental tissue[6,7], as this is related to the evaporation rate of the solvent mixture[7].

The residual solvent dispersed within the matrix monomers may difficult the polymerization and reduce degree of conversion of resin materials and thus compromise the mechanical property of the material[8]. Some studies have

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demonstrated that the physical properties of the adhesive resin may have a profound influence on the resin-dentin bond strengths. Low mechanical property values can compromise the bond strength of the adhesive system[8-12]. Another way, presence of acid monomers may influence on the bondig to dentin and the mechanical properties of adhesive systems[9,10]. One of the keys of success with adhesives is the chemical bonding capability of their functional monomers to hydroxyapatite (HAp).

Therefore, it is essential for clinicians to know the mechanical behavior of some adhesive systems on the market to enable them to perform successful restorations.

Some studies have suggested that laboratory-screening tests continue to be indispensable in providing data to predict clinical effectiveness of dental products[13,14].

Therefore, the aim of this study was to compare several properties of three two-step etch-and-rinse adhesives available on the market; among these, the microtensile bond strengths (µTBS), degree of conversion, modulus of elasticity and ultramicrohardness. The null hypothesis tested was that would be no difference in the mechanical-physical properties of the adhesive systems evaluated.

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Evaluation of degree of conversion, microtensile bond strength and mechanical properties of three etch-and-rinse dental adhesives

2. Materials and Methods 2.1 Microtensile bond strength After extraction, fifteen teeth were cleaned and stored in distilled water until used for the study. The adhesive systems used were Ambar (FGM, Joinville, SC, Brazil), Optibond (Kerr, Orange, CA, USA) and Magic Bond (Vigodent, Rio de Janeiro, RJ, Brazil), details of the composition of materials are described in Table 1. The teeth were divided into groups according the adhesive system used (n=5). Occlusal enamel of each tooth was removed, using a diamond disc (Isomet 1000, Buehler, Lake Bluff, Illinois, USA) under constant water-cooling. The occlusal surface was abraded with SiC paper #320 under water irrigation on a polishing machine (Aropol-E, Arotec, Cotia, SP, Brazil) to expose middle-depth dentin. A standardized smear layer was then created with SiC paper #600, under continuous irrigation for 60 s. All adhesives were applied in a controlled environment using the bonding protocols summarized in Table 1.

Resin composite build-ups (Opallis A2, FGM, Joinville, SC, Brazil) were placed on the bonded surfaces (three increments for 1.5 mm each) that were individually light‑activated for 40s each (Optilux 501, Kerr, Orange, CA, USA) at 600 mW/cm2. After storage in distilled water at 37 °C for 24 h, the specimens were sectioned in both “x” and “y” directions, perpendicular to the adhesive/tooth interface using a water-cooled diamond saw (Isomet 1000, Buehler, Lake Bluff, Illinois, USA) to obtain rectangular beams. The number of prematurely debonded beams (PDS, specimens that failed prematurely during preparation and handling) per tooth was recorded. After 24 hrs the specimens were tested.

Before testing, the cross-sectional area of each stick was measured with the digital caliper to the nearest 0.01 mm and recorded to calculate the μTBS (Absolute Digimatic, Mitutoyo; Tokyo, Japan). The average cross-sectional area of specimens ranged from 0.79 to 0.82mm2. Each stick was fixed with cyanocrylate glue (Pegamil Bond Gel, Buenos Aires, Argentina) and tested under tension at a crosshead speed of 1.0 mm/minute using an Instron testing machine (Instron 3342, Canton, MA, USA) equipped with a load cell of 500 N. The fractured surface of each test specimen was evaluated under a stereoscopic microscope (Kozo Optical and Electronical Instrumental, Nanjing, Jiangsu, China) at 40x magnification and classified as cohesive (failure exclusively within substrate or resin composite) and adhesive/mixed

(failure at resin/substrate interface or mixed with partially cohesive failure of the neighboring substrates).

All values obtained from each tooth were averaged for statistical purposes. For each bonding substrate, the bond strength values were subjected to One-way repeated measures ANOVA and the Tukey test for comparisons (α=0.05).

2.2 Degree of conversion analysis The DC was analyzed by FTIR spectrometer (IRPrestige-21, Shimadzu Corporation, Kyoto, Japan) equipped with an attenuated total reflectance crystal (ATR‑MIRacleSingle Reflection Horizontal, Pike Technologies, Inc. Madison, WI, USA). The absorption spectra of each uncured adhesive were obtained by placing two drops of each adhesive solution directly onto the surface of the ATR diamond crystal. The absorption spectra of each polymerized adhesive specimen were obtained by dispensing two drops of the tested adhesive on an individual acetate strip and this was subsequently light-cured. After polymerization, the flat cured surface of the adhesive was firmly placed against the ATR crystal to collect the spectra. FTIR readouts were made at 22 ± 1°C with 50% relative humidity.

For the adhesive systems containing aromatic vinyl bonds of bisphenol and aliphatic bonds of the methacrylate functional group, the DC measurements were made with the relative intensity of the aromatic component band with the main peak around 1608 cm-1, relative to the band with aliphatic carbon-to-carbon double-bond absorbance main peak around 1638 cm-1, which changes with the polymerization of the composite (Figure 1)[15]. Thus, for the adhesive systems Ambar, Optibond and Magic Bond the absorption spectra were obtained from the region between 1650 cm-1 and 1595 cm-1 with 30 scans at a resolution of 4 cm-1.

The DC (%) was calculated using the following equation: DC (%) = 100 x [1-(R-cured/R- uncured)][16], where R represented the ratio between the absorbance peak around 1638 cm-1 and 1608 cm-1. The data were analyzed by one-way analysis of variance (ANOVA) and post hoc Tukey tests (α=0.05).

2.3 Ultramicrohardness and modulus of elasticity Six teeth were cut and 12 halves of tooth fragments were obtained. These fragments were distributed into groups among adhesives system used (n=4). These halves were restored in the same way as described for the microtensile bond strength test. Crowns fragments were embedded in

Table 1. Composition, manufacturer, lot number and mode of application of adhesive systems tested in this study.

Adhesive Systems Ambar [ABA, FGM] 10311 Optibond S [OPB, KERR] 3462530 Magic Bond DE [MGB, VIGODENT] 012/11

Composition Application Mode UDMA, HEMA, acidic methacrylate monomers, 10-MDP, (1): acid-etch (15 s) methacrylate hydrophilic monomers, silanized silicon (2): rinse (15 s) dioxide, camphorquinone, 4-EDAMB, ethanol.

(3): air-dry at 20cm distance (30 s) BIS-GMA, HEMA, GDM, GPDM, ethanol, silica, barium (4): dentin rewetted with water glass, camphorquinone. Approximately 15% of filler (5): 1st one coat of adhesive under agitation (10 s) weight to 0.4 microns.

(6): air-dry (10s/20cm) HEMA, dimethacrylates, neopentyl fluoride, acrylate (7): 2nd coat of adhesive under agitation (10 s) fluoride, adhesive monomer (MEP), highly dispersed (8): air-dry (10s/20cm) silicon dioxide, photoinitiators and stabilizers in an (9): light-cure (10 s – 600 mW/cm2) alcoholic solution

Polímeros, 27(3) , 230-236, 2017

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Freitas, S. A. A., Lanza, M. D. S., Carneiro, K. K., Loguercio, A. D., & Bauer, J.

Figure 1. Fourier transform infrared spectra of adhesive systems. (A) Ambar; (B) Optibond; (C) Magic Bond, 1638 cm–1 = absorbance intensity for aliphatic carbon–carbon double bonds, 1608 cm–1 = internal reference for aromatic carbon–carbon double bonds.

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Evaluation of degree of conversion, microtensile bond strength and mechanical properties of three etch-and-rinse dental adhesives transparent acrylic resin (Orthodontic Resin, Dentsply Caulk, Milford, DE, USA) and flattened in a polishing machine (Unipol 1210, MTI Corporation, CA, USA) under water‑cooling with different grit silicon carbide abrasive papers and followed by a final polishing performed using diamond pastes from 1 μm and lower. The modulus of elasticity and hardness across the interface were evaluated after 24 h of storage in distilled water at 37°C. The measurements were obtained using a dynamic ultra microhardness tester (DUH-211S, Shimadzu Corporation, Kyoto, Japan). The indenter used was a pyramidal-triangular shaped Berkovich indenter (115°) and the maximum force applied was 19.61 mN. Each indentation was accomplished by a load-unload step at a speed of 4.44mN/sec. The hold time at the maximum load was 7sec and the hold time with minimum load was 5sec. The elastic modulus calculation for each read out was measured between of 100 and 70% of the test curve. The software associated with the DUH-211S analyzes the data, calculated and expressed the hardness and the elastic modulus values from the load-unload curve recorded. Four indentations were performed in the adhesive layer. The distance between each indentation was kept constant by adjusting the distance intervals to 10 (±1) μm.

The modulus of elasticity and ultramicrohardness data were analyzed by one-way analysis of variance (ANOVA) and post hoc Tukey tests (α=0.05).

3. Results 3.1 Microtensile bond strength Statistical analysis was performed with SPSS Statistics for Windows version 20 (SPSS; Chicago, IL, USA).

The Shapiro-Wilk test indicated normal distribution of the data (p = 0.610) that were analyzed by using one-way ANOVA and Tukey tests. The Tukey post hoc test showed the highest µTBS was obtained for Ambar and Optibond with no statistical difference (p>0.005). The lowest µTBS values were found for Magic Bond with statistical difference (p=0.0034) (Table 2). The failure mode distributions for all adhesive system tested are depicted in Figure 2.

The predominant failure mode was adhesive/mixed between the adhesive and dentin.

3.2 Degree of conversion The DC was significantly influenced for some of the adhesive systems analyzed (p<0.005). The Tukey post hoc test showed the highest DC was obtained for Ambar and Optibond with no statistical difference (p>0.005).

The lowest DC values were found for Magic Bond with statistical difference (p=0.0098) (Table 2).

3.3 Ultramicrohardness and Modulus of elasticity The ultramicrohardness and modulus of elasticity were significantly influenced for some of the adhesive systems analyzed (p<0.005). The Tukey post hoc test showed the

Figure 2. Incidence of failure modes (%) analyzed by stereomicroscopy of the Magic Bond, Ambar and Optibond systems.

Table 2. Mean values of microtensile bond strength (µTBS), degree of conversion (%), ultramicrohardness and modulus of elasticity (GPa) of adhesive systems used*.

Materials Ambar Optibond Magic Bond

µTBS (MPa) 60.4 ± 8.7 a 59.7 ± 16.1 a 27.8 ± 14.2 b

Degree Conversion (%) 53.9 ± 4.5 a 58.6 ± 0.9 a 44.5 ± 9.4 b

Ultramicrohardness 21.0 ± 1.3 b 28.7 ± 1.3 a 23.4 ± 1.6 b

Modulus (GPa) 0.69 ± 0.09 b 0.87 ± 0.008 a 0.75 ± 0.12 b

*Different letters in the vertical mean statistical difference.

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Freitas, S. A. A., Lanza, M. D. S., Carneiro, K. K., Loguercio, A. D., & Bauer, J. highest ultramicrohardness and modulus of elasticity were obtained for Optibond with statistical difference from Ambar and Magic Bond (p=0.0003 and p=0.002 respectively) (Table 2).

4. Discussion The null hypothesis tested in this study was rejected because the adhesives tested show different mechanicalphysical properties. Microtensile bond strength is the property that suggests that the quality and quantity of monomer infiltrated into the demineralized substrate and high bond strength values can be an excellent indicator of the clinical behavior of the material[17].

The high bond strength values obtained with the Optibond adhesive system may be related to the presence of the monomer GPDM in its composition. The greater capability of the GPDM monomer could be speculated to be due two factors: (1) the phosphate group was responsible for acid etching on dentin and (2) development of a chemical bond with the mineral component[18-21]. On the other hand, it has two methacrylate functional groups for copolymerization with other methacrylate monomers to provide increased crosslinking density and enhanced mechanical strength for the polymerized adhesive[14]. Perhaps for this reason this adhesive system showed a high degree of conversion, ultramicrohardness and modulus of elasticity.

In this same way, the high microtensile bond strength values for Ambar can be explained by the presence of multifunctional monomer 10-MDP. This monomer contains a polymerizable methyl-methacrylate group and a phosphate group responsible for ionic interaction with calcium[19,22], 10-methacryloyloxydecyl dihydrogenphosphate (MDP) has demonstrated a very effective and durable bond to dentine, due to the low solubility of the calcium salt that forms on the hydroxyapatite surface.

Moreover, the Magic Bond adhesive system showed the lowest bond strength values. Fontes et al.[23] found poor results with this material, and most of the specimens obtained from this material in the present study were lost during the cutting procedure. These same authors explained that this low microtensile bond strength obtained by the Magic Bond system is due to lack of accurate information about the evaporation of the solvent and water. In the present study, the authors standardized the application mode for all adhesive systems and the type of solvent (ethanol).

While the adhesive systems Ambar and Optibond have monomers with the ability to bind to the hydroxyapatite component of dentin substrate, the adhesive MGB presents a generics monomers and dimethacrylate monomer (MEP).

Perhaps, the monomers present in Magic Bond adhesive did not have the ability to form a chemical bond with the dental substrate.

Several studies have endeavored to correlate the mechanical property values with the bond strength results[17,24-26]. Hasegawa et al.[24] reported that tensile bond strength values to dentin significantly correlated with the values of mechanical properties of the resin composites such as tensile strength, flexural strength and modulus of the adhesive.

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In the present study the Magic Bond adhesive system presented elastic modulus and ultramicrohardness values similar to those obtained with the adhesive system Ambar.

The importance of the mechanical properties of the material with regard to bond strength is due to the hybridization/ infiltration of the adhesive system into the collagen network and dentinal tubules, thereby creating greater micromechanical interlocking with this tissue after light polymerization.

An adhesive system with a high mechanical strength could also resist the stresses generated during a μTBS test and thus present a high bonding to dentin[13]. Perhaps the high modulus of elasticity and hardness of the adhesive system was not sufficient to ensure high bonding strength to dentin[17]. According to the results obtained in this study, it is clear that the acid monomers with a high affinity for the mineral must be present in the adhesive system to ensure chemical interaction with the substrate. However, an adhesive system with low mechanical properties may compromise the resin/dentin bond interface[25,26].

The adhesive systems that presented a high degree of conversion also showed good performance in the μTBS test.

Conversion of monomer into polymer plays an important role in successful dentin bonding[27]. A low degree conversion of dental adhesives is associated with low bond strength values, mechanical properties and high permeability[28-30].

Hass et al.[31] was recently able to demonstrate that there is a strong correlation between the degree of conversion of adhesive systems and μTBS values. Another way, a good performance of an adhesive system to bond with the dentin depends on countless factors. However, it seems clear that a high degree of conversion[32] and the presence of monomers with an affinity for hydroxyapatite[33,34] it is desirable for a good bonding system.

5. Conclusions The Optibond adhesive system was capable that show a high degree of conversion, ultramicrohardness, modulos of elasticity and resin–dentin microtensile bond strength.

6. Acknowledgements This study was supported by the Foundation for the Support of Scientific and Technological Research of Maranhão (FAPEMA - BEPP-03730/13).

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Operative Dentistry, 39(6), 652-662. PMid:24720263. http:// dx.doi.org/10.2341/13-236.

15. Borges, B. C. D., Souza-Junior, E. J., Brandt, W. C., Loguercio, A. D., Montes, M. A. J. R., Puppin-Rontani, R. M., & Sinhoreti, M. A. C. (2012). Degree of conversion of simplified contemporary adhesive systems as influenced by extended air-activated or passive solvent volatilization modes.

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16. Pianelli, C., Devaux, J., Bebelman, S., & Leloup, G. (1999).

The Micro-Raman spectroscopy, a useful tool to determine the degree of conversion of light-ativated composites resins.

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PMid:10490681. http://dx.doi.org/10.1002/(SICI)10974636(1999)48:5<675::AID-JBM11>3.0.CO;2-P. 17. Takahashi, A., Sato, Y., Uno, S., Pereira, P. N., & Sano, H. (2002). Effects of mechanical properties of adhesive resins on bond strength to dentin. Dental Materials, 18(3), 263-268. PMid:11823019. http://dx.doi.org/10.1016/S01095641(01)00046-X. 18. Goracci, C., Margvelashvili, M., Giovannetti, A., Vichi, A., & Ferrari, M. (2013). Shear bond strength of orthodontic brackets bonded with a new self-adhering flowable resin composite.

Clinical Oral Investigations, 17(2), 609-617. PMid:22538472. http://dx.doi.org/10.1007/s00784-012-0729-x.

19. Yoshida, Y., Nagakane, K., Fukuda, R., Nakayama, Y., Okazaki, M., Shintani, H., Inoue, S., Tagawa, Y., Suzuki, K., De Munck, J., & Van Meerbeek, B. (2004). Comparative study on adhesive performance of functional monomers. Journal of Dental Research, 83(6), 454-458. PMid:15153451. http://dx.doi.org /10.1177/154405910408300604.

20. Sezinando, A., Perdigão, J., & Regalheiro, R. (2012). Dentin bond strengths of four adhesion strategies after thermal fatigue and 6-month water storage. Journal of Esthetic and Restorative Dentistry, 24(5), 345-355. PMid:23025319. http://dx.doi. org/10.1111/j.1708-8240.2012.00531.x.

21. Stona, P., Borges, G. A., Montes, M. A., Burnett, L. H., Jr., Weber, J. B., & Spohr, A. M. (2013). Effect of polyacrylic acid on the interface and bond strength of self-adhesive resin cements to dentin. The Journal of Adhesive Dentistry, 15(3), 221-227. PMid:23560256. http://dx.doi.org/10.3290/j.jad. a29531.

22. Luque-Martinez, I. V., Perdigão, J., Muñoz, M. A., Sezinando, A., Reis, A., & Loguercio, A. D. (2014). Effects of solvent evaporation time on immediate adhesive properties of universal adhesives to dentin. Dental Materials, 30(10), 1126-1135.

PMid:25139815. http://dx.doi.org/10.1016/j.dental.2014.07.002.

23. Fontes, S. T., Cubas, G. B., Flores, J. B., Montemezzo, M. L., Pinto, M. B., & Piva, E. (2010). Resin-dentin bond strength of 10 contemporary etch-and-rinse adhesive systems after one year of water storage. General Dentistry, 58(6), 257-261.

PMid:21062710.

24. Hasegawa, T., Itoh, K., Koike, T., Yukitani, W., Hisamitsu, H., Wakumoto, S., & Fujishima, A. (1999). Effect of mechanical properties of resin composites on the efficacy of the dentin bonding system. Operative Dentistry, 24(6), 323-330. http:// dx.doi.org/10.2341/1559-2863-24-6-1. PMid:10823080.

25. Carrilho, M. R., Carvalho, R. M., Tay, F. R., & Pashley, D. H. (2004). Effects of storage media on mechanical properties of adhesive systems. American Journal of Dentistry, 17(2), 104-108. PMid:15151336.

26. Carrilho, M. R. O., Tay, F. R., Pashley, D. H., Tjäderhane, L., & Carvalho, R. M. (2005). Mechanical stability of resindentin bond components. Dental Materials, 21(3), 232-241.

PMid:15705430. http://dx.doi.org/10.1016/j.dental.2004.06.001.

27. Eick, J. D., Gwinnett, A. J., Pashley, D. H., & Robinson, S. J. (1997). Current concepts on adhesion to dentin. Critical Revision of Oral Biology and Medcine, 8(3), 306-335. PMid:9260046. http://dx.doi.org/10.1177/10454411970080030501.

28. Kanehira, M., Finger, W. J., Hoffmann, M., Endo, T., & Komatsu, M. (2006). Relationship between degree of polymerization and enamel bonding strength with self-etching adhesives. The Journal of Adhesive Dentistry, 8(4), 211-216. PMid:16958284.

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Freitas, S. A. A., Lanza, M. D. S., Carneiro, K. K., Loguercio, A. D., & Bauer, J. 29. Wang, Y., Spencer, P., Yao, X., & Ye, Q. (2006). Effect of coinitiator and water on the photoreactivity and photopolymerization of HEMA/camphorquinone-based reactant mixtures. Journal of Biomedical Materials Research. Part A, 78(4), 721-728. PMid:16739171. http://dx.doi.org/10.1002/ jbm.a.30733.

30. Cadenaro, M., Antoniolli, F., Codan, B., Agee, K., Tay, F. R., Dorigo, E. S., Pashley, D. H., & Breschi, L. (2010). Influence of different initiators on the degree of conversion of experimental adhesive blends in relation to their hydrophilicity and solvent content. Dental Materials, 26(4), 288-294. PMid:20018363. http://dx.doi.org/10.1016/j.dental.2009.11.078.

31. Hass, V., Dobrovolski, M., Zander-Grande, C., Martins, G. C., Gordillo, L. A. A., Accorinte, M. L. R., Gomes, O. M. M., Loguercio, A. D., & Reis, A. (2013). Correlation between degree of conversion, resin-dentin bond strength and nanoleakage of simplified etch-and-rinse adhesives. Dental Materials, 29(9), 921-928. PMid:23830512. http://dx.doi.org/10.1016/j. dental.2013.05.001.

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32. Ferracane, J. L. (2006). Hygroscopic and hydrolytic effects in dental polymer networks. Dental Materials, 22(3), 211-222.

PMid:16087225. http://dx.doi.org/10.1016/j.dental.2005.05.005.

33. Van Landuyt, K. L., Yoshida, Y., Hirata, I., Snauwaert, J., De Munck, J., Okazaki, M., Suzuki, K., Lambrechts, P., & Van Meerbeek, B. (2008). Influence of the chemical structure of functional monomers on their adhesive performance. Journal of Dental Research, 87(8), 757-761. PMid:18650548. http:// dx.doi.org/10.1177/154405910808700804.

34. Yoshihara, K., Yoshida, Y., Nagaoka, N., Fukegawa, D., Hayakawa, S., Mine, A., Nakamura, M., Minagi, S., Osaka, A., Suzuki, K., & Van Meerbeek, B. (2010). Nano-controlled molecular interaction at adhesive interfaces for hard tissue reconstruction.

Acta Biomaterialia, 6(9), 3573-3582. PMid:20346420. http:// dx.doi.org/10.1016/j.actbio.2010.03.024.

Received: Mar. 02, 2016 Revised: June 27, 2016 Accepted: Aug. 28, 2016

PolĂ­meros, 27(3)â&#x20AC;&#x160;, 230-236, 2017

http://dx.doi.org/10.1590/0104-1428.14216

Thermal and mechanical properties of bio-based plasticizers mixtures on poly (vinyl chloride) Boussaha Bouchoul1*, Mohamed Tahar Benaniba1 and Valérie Massardier2 Laboratoire des Matériaux Polymérique MultiPhasiques – LMPMP, Département de Génie des Procédés, Faculté de Technologie, Université Sétif 1, Algérie 2 Laboratoire Ingénierie des Matériaux Polymères UMR 5223, Centre National de la Recherche Scientifique – CNRS, Institut National des Sciences Appliquées de Lyon – INSA, Lyon, Villeurbanne, France

1

*B_Bouchoul@yahoo.com

Abstract The use of mixtures of nontoxic and biodegradable plasticizers coming from natural resources is a good way to replace conventional phthalates plasticizers. In this study, two secondary plasticizers of epoxidized sunflower oil (ESO) and epoxidized sunflower oil methyl ester (ESOME) were synthesized and have been used with two commercially available biobased plasticizers; isosorbide diesters (ISB) and acetyl tributyl citrate (ATBC) in order to produce flexible PVC. Different mixtures of these plasticizers have been introduced in PVC formulations. Thermal, mechanical and morphological properties have been studied by using discoloration, thermogravimetric analysis (TGA), differential scanning calorimetry (DSC), dynamic mechanical thermal analysis (DMTA), tensile - strain and scanning electron microscopy (SEM). Studies have shown that PVC plasticization and stabilization were improved by addition of plasticizers blends containing ISB, ATBC, ESO and ESOME. An increase in the content of ESO or ESOME improved thermal and mechanical properties, whereas ESOME/ATBC formulations exhibited the best properties.

Keywords: PVC, epoxidized sunflower oil, epoxidized sunflower oil methyl ester, isosorbide diesters, acetyl tributyl citrate.

1. Introduction Poly (vinyl chloride) (PVC) is very present in daily living applications due to its diverse properties and low cost[1,2].

Its properties depend on the amount of different kinds of additives such as plasticizers[3]. The plasticizer is a very important additive of PVC; it can improve the flexibility of PVC without changing its chemical properties[3-6]. Phthalates are the most commonly used plasticizers in PVC with applications in food packaging, medical devices, children’s toys, building materials, and other common products[3].

Unfortunately, phthalates contaminate indoor environments, human food and are environmental contaminants. It has been reported by Bhakti et al.[3] that several phthalates and especially diethyl hexyl phthalate (DEHP) also known as di-octyl phthalate (DOP) are suspected of having carcinogenic and toxic effects.

Recently, several alternatives exist to substitute DEHP in PVC applications[3]. Nowadays, there is an increasing interest in the use of nature based plasticizers for PVC[7,8].

Some studies have been done on the use of epoxidized sunflower oil (ESO) as secondary plasticizer to partially replace di-2-ehylhexyl phthalate (DEHP) in PVC formulations[9,10].

The compatibility of plasticizers with PVC also needs to be considered[11]. Solubility parameters are often used to predict PVC/plasticizer interactions. Decrease in the glass transition temperature (Tg) of PVC can also be studied to assess plasticization efficiency[12].

Epoxidized sunflower oil (ESO) and epoxidized sunflower oil methyl esters (ESOME) were synthesized and then characterized by oxirane index titration and FTIR spectroscopy.

The aim of the work reported here is to investigate the use

Polímeros, 27(3) , 237-246, 2017

of these new products as secondary plasticizers, mixed with isosorbide diesters (ISB) and acetyl tributyl citrate (ATBC) to plasticize PVC with 60 parts of plasticizers mixtures.

The plasticization of PVC with ISB, as well as with its mixtures containing ESO and DEHP, has been studied in our previous work[13]. Although ATBC and ISB have been studied as individual plasticizers for this polymer[13,14], ATBC, ISB and their mixtures with ESO or ESOME in PVC is the axis of the present research to evaluate thermal, mechanical and morphology characteristics of plasticized PVC. Discoloration degree of sheets, thermogravimetry analysis (TGA), dynamic mechanical thermal analysis (DMTA), differential scanning calorimetry (DSC), mechanical properties and scanning electron microscopy (SEM) were used to investigate the properties of our PVC plasticized with ATBC, ISB and their blends with ESO or ESOME.

2. Materials and Methods 2.1 Materials PVC suspension grade resin (SE 950, K= 65.7-67.1), was kindly supplied by Shintech (Houston, USA). Plasticizers used were as follows: acetyl tributyl citrate (ATBC) (Sigma Aldrich, USA), isosorbide diesters (ISB) (ID47, Roquette Frères, France), sunflower oil (SO) with an iodine value index, Iiod, of 130g I2/100g (Cevital Bejaia, Algeria), epoxidized sunflower oil (ESO) with 6.1% of oxirane oxygen (conversion = 98.2%, yield = 80.5%) and epoxidized sunflower oil methyl ester (ESOME), were prepared in our

237

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

Bouchoul, B., Benaniba, M. T., & Massardier, V. laboratory. Ca/Zn stearates and stearic acid (SA) were used as thermal stabilizer and lubricant, respectively.

was averaged over 32 scans. Spectra were obtained by direct measurement.

2.2 Methods

2.2.4 Preparation of PVC sheets

2.2.1 Epoxidation of sunflower oil Sunflower oil (SO) and formic acid were combined in a 250 mL three neck flask equipped with stirrer, reflux condenser and thermocouple. The three neck flask was immersed in a heating mantle. Reaction occurred during mixing of the feed components and the reaction temperature was controlled at 55°C. To start the epoxidation, hydrogen peroxide solution (30%) was then added dropwise into the mixture for the first 30 minutes of reaction. The mole ratio of carbon double bonds to hydrogen peroxide (C=C: H2O2) was 1: 1.5. After feeding H2O2 was completed, the reaction continued by mixing at a stirring speed of 700 rpm and controlling the temperature at 55°C for a further two hours. After that, the product of reaction was next cooled and decanted for the separation of the organic soluble compounds (epoxidized oil) from the water soluble phase.

The epoxidized sunflower oil (ESO) was then washed with warm water to remove residual contaminants. Diethyl ether was used to enhance the water separation[15]. The washed organic layer was also dried with centrifugation to remove water traces in the epoxidized oil. The epoxy functionality was determined using AOCS Cd 9-57 method[16].

2.2.2 Transesterification of epoxidized sunflower oil Reactions were performed in a 250 mL flat bottom glass vessel. Epoxidized sunflower oil was introduced to the glass vessel along with a magnetic stir bar. Sodium methoxide catalyst was dissolved in methanol and then added to the epoxidized sunflower oil in the reactor[17]. This mixture was heated to 50°C on a hot plate and magnetically stirred at 700 rpm. The reaction was stopped after 2 hours and the reaction mixture was transferred to a separatory funnel to recover the products. Then, the oxirane index is measured to confirm the no opening of oxirane cycles.

2.2.3 Fourier transform infrared spectroscopy (FTIR) The samples were identified using Fourier transform infrared spectroscopy with a 680 Nicolet thermo spectrometer (Thermo Scientific, USA), employing an Attenuated Total Reflection (ATR) accessory, equipped with a diamond crystal cell, angle 45°. The spectra were acquired in the range 4000-500 cm-1 at a resolution of 4 cm-1 and the signal

The plasticized PVC sheets were prepared using PVC, ISB, ATBC, ESO and ESOME. Various combinations (on weight basis) were prepared and used as plasticizing mixtures. Four series of blends were prepared: PVC/(ISB/ESO), PVC/(ISB/ESOME), PVC/(ATBC/ESO), PVC/(ATBC/ESOME) and one formulation of PVC/DEHP as reference. In each case, 60 parts of this mixture was mixed with 100 parts of PVC. The different combinations used for the preparation of PVC sheets are displayed in Table 1.

Samples prepared with different plasticizer combinations were further mixed for 10 minutes by using a two roll mill at a temperature of 160°C to obtain a homogeneous blend.

2.2.5 Discoloration sheets by static heat Circular test pieces (19 mm diameter) were placed in a circulating air oven maintained at 177± 2°C. Test pieces were removed at intervals of 5, 10, 15, 20, 30, 45, 60, 80, 100, 130, 160 and 200 minutes. The color changes with the heating time were observed and measured with color Synmero scale as described by Ocskay et al.[18].

2.2.6 Thermogravimetry analysis (TGA) Thermal degradation studies were conducted using a thermogravimetry analyzer TGA (Q500 TA Instruments, USA). All the samples were evaluated from ambient to 500°C under nitrogen flow (40 mL/min) at 10°C/min heating rate.

2.2.7 Differential scanning calorimetry (DSC) Glass transition temperature (Tg) was determined by differential scanning calorimetry using a DSC Q10 (TA Instruments, USA) calibrated with indium. Samples of about 8 mg were conditioned in aluminum pans, equilibrated at 25°C and held isothermally for 1 min. Then, they were heated to 220°C at 20°C/min and held isothermally for 3 min at 220°C to eliminate previous thermal history. The samples were then cooled to -60°C and held isothermally for 3 min and finally heated to 220°C at a constant rate of 10°C/min.

2.2.8 Dynamic mechanical thermal analysis (DMTA) All tests were conducted with RSA Rheometrics.

The analyses were carried out on specimens measuring 30x7x0.5 mm3 in tensile mode with a temperature ranging from -60°C to 80°C at a 3°C/min heating rate with frequency

Table 1. Compositions of our plasticized formulations.

Materials (phr) PVC SA Ca/Zn ISB ATBC ESO ESOME

238

1 100 1 2 60 -

2 100 1 2 50 10 -

3 100 1 2 40 20 -

4 100 1 2 30 30 -

5 100 1 2 60 -

6 100 1 2 50 10 -

Formulation number 7 8 9 100 100 100 1 1 1 2 2 2 50 40 30 20 30 10

10 100 1 2 40 20

11 100 1 2 30 30

12 100 1 2 50 10

13 100 1 2 40 20

14 100 1 2 30 30

Polímeros, 27(3) , 237-246, 2017

Thermal and mechanical properties of bio-based plasticizers mixtures on poly (vinyl chloride) of 1, 3, 5 or 10 Hz. Tα was taken at the peak in the tan δ curve accompanied by a step reduction in the storage modulus.

2.2.9 Mechanical properties Test specimens were cut according to NF T 51-034 H1.

The Young’s modulus, tensile strength and elongation at break were determined at room temperature with a universal dynamometer (Instron, ref. 33R4469, USA) at constant rate of 50 mm/min by using a 0.5 kN sensor.

2.2.10 Scanning electron microscopy (SEM) Morphological studies were conducted using a scanning electron microscope JEOL (model JFM-6360 LV). Specimens were sputter-coated with gold to a thickness of 10 nm before surface characterization to prevent charging. The SEM was equipped with a lanthanum hexaboride (LaB6) crystal as electron emitter source. An accelerating voltage of 13 kV was used to collect the SEM images.

3. Results and Discussion 3.1 Epoxidation of sunflower oil (ESO) There are two main reactions involved in the epoxidation reaction as mentioned above. During the first stage, peroxy acid is formed from the reaction of formic acid and hydrogen peroxide while, in a second stage, epoxidized sunflower oil is produced from the reaction between peracid and double bonds in the sunflower oil[19].

The following reaction scheme considers hydrogen peroxide as oxygen donor and acetic acid as oxygen carrier[20]: Step 1: Formation of Peroxy formic acid CH3COOH + H2O2

Schuchardt et al.[21]. have reported that the overall process is a sequence of three consecutive and reversible reactions, in which di and monoglycerides are formed as intermediates.

The stoichiometric reaction requires 1 mol of ESO and 3 mol of methanol. However, an excess of alcohol is used to increase the yields of epoxidized alkyl esters and to allow its phase separation from the glycerol formed[21].

3.3 Fourier transform infrared spectroscopy (FTIR) Products were characterized by FTIR (Figure 1) in order to follow the disappearance of double bonds and formation of epoxy groups. For SO, the characteristic peaks at 3008.44 cm-1 and 723.95 cm-1 are attributed to the stretching vibration of the double bonds =C-H2 and -CH=CH-, respectively. The FTIR spectrum displays peaks of ESO at 863.58 cm-1 and of ESOME at 845.30 cm-1, characteristic of the C-O-C oxirane stretch and disappearance of the double bond at 3008.44 cm-1. It indicates the almost complete conversion of C=C unsaturations to epoxy groups and confirms the success of the epoxidation reaction of ESO[22,23] on fatty acid chains.

The other new peaks at 3373.35 cm-1 for ESO and at 3374.51 cm-1 for ESOME are attributed to the hydroxyl functional group, derived from the epoxy functional group via partial epoxy ring opening reaction. The epoxy ring opening reaction could occur either by acid catalysis in the presence of water

CH3COOOH + H2O (1)

Step 2: Epoxidation Reaction O -CH2=CH2- + CH3COOOH

-HC

CH- + CH3COOH

(2)

3.2 Transesterification of epoxidized sunflower oil In the transesterification, ESO reacts with methanol in the presence of NaOH, and produces a mixture of epoxidized fatty acids alkyl esters and glycerol (Scheme 1).

Figure 1. FTIR of (A) SO; (B) ESO and (C) ESOME.

Scheme 1. Reaction for the synthesis of ESOME from ESO. Polímeros, 27(3) , 237-246, 2017

239

Bouchoul, B., Benaniba, M. T., & Massardier, V.

Figure 2. Discoloration degree of plasticized PVC films with: (A) ISB/ESO; (B) ATBC/ESO; (C) ISB/ESOME; (D) ATBC/ESOME.

associated with aqueous solution of H2O2 used[15]. The peak in the spectrum of ESOME at 1739.92 cm-1 indicates the shifting of C=O absorption band and the characteristic ester band at 1027.80 cm-1, these results indicate that the functional group has been converted into –COOC– for the three ester functions[23,24].

3.4 Discoloration sheets by static heat A study of coloration change is an important method for investigating the stabilizing performance of plasticizers in PVC when the variation of the PVC color from transparency towards yellow to brown then black, indicate the evolution of the degradation state of PVC sheets.

The discoloration degree of PVC samples before and after degradation was determined by using the Synmero scale.

The samples are colored because of the dehydrochlorination and formation of conjugated double bonds in PVC sheets.

During dehydrochlorination, HCl molecules are extracted, which initiates the formation of double bonds and polyenes in PVC chains and noticeable color change from yellow to black, when the conjugated polyene sequences contain more than four to five double bonds[25]. Figure 2 shows the discoloration degree of sheets containing different mixtures of plasticizers. All sheets turn brownish with discoloration degree (C≈5) after 40 min.

The sheets containing ISB and ATBC alone were blackening after 120 minutes. But the dispersion of ESO, the degree of discoloration decreases by increasing the rate 240

of ESO, either mixed with ISB (Figure 2A) or with ATBC (Figure 2B) in the plasticizer system. In the case of ESOME, (Figure 2C) and (Figure 2D), the sheets have a yellow color for times in the range [0-200 minutes].

This yellowish appearance could be related to polyene formation due to dehydrochlorination of PVC with the reaction of HCl and ESOME[26]. Thus, the combination of ESO and ESOME with ISB or ATBC increases the thermal stability of formulations compared with samples containing ISB or ATBC alone. The changes in discoloration degrees for the PVC were minimal in the films plasticized in presence of ESO or ESOME. Epoxy groups of ESO and ESOME neutralize the HCl evolved by dehydrochlorination of PVC. Thus, auto acceleration of dehydrochlorination by evolved HCl is prevented because it is captured by epoxy groups, owing epoxy groups in ESO and ESOME to react with HCl, generated from PVC[27], as shown in reaction 3.

(3)

3.5 Thermogravimetry analysis Figure 3 shows the TGA and DTG curves of plasticized PVC by some plasticizer blends. The curves have similar shapes for all formulations and each one presents two distinct stages.

Polímeros, 27(3) , 237-246, 2017

Thermal and mechanical properties of bio-based plasticizers mixtures on poly (vinyl chloride) Table 2 summarizes the thermal performance data of plasticized PVC by the different systems in the first and second stage respectively.

The characteristical thermal parameters selected were the 1% weight loss (T 1%), which is the initial weight loss temperature, maximum degradation temperature (T1 max and T2 max), which is the highest thermal degradation rate temperature obtained from the peak of weight loss, maximum speed of degradation (S1 max and S2 max), the 5% weight loss (T 5%), the 50% weight loss (T 50%)

and the residue. The elimination of a large amount of HCl happened at about 200°C. It could be attributed to the first thermal degradation stage. The second stage above 400°C is attributed to crosslinking of chains containing C=C bonds as the process of thermal degradation of polyenes involves cyclization and splitting of chains[8]. The results obtained in this range show that the PVC mixed with ATBC begins to lose weight at a lower temperature than that with ISB. However, formulations with ESO (formulation number 4, 8) or with ESOME (formulation number 11, 14) present the minimal weight loss in the first step.

Figure 3. TGA and DTG curves for PVC plasticized with ISB, ATBC and their blends with 30 phr of ESO and ESOME. Table 2. Thermal gravimetric analysis (TGA) results of the formulations.

Formulation number 1 2 3 4 5 6 7 8 9 10 11 12 13 14

First stage T1% (°C)

T1 max (°C)

S1 max (%/°C)

184.0 193.1 196.9 204.0 160.9 165.4 165.8 168.9 196.0 197.8 198.0 164.3 167.0 172.3

292.2 293.4 297.6 302.8 281.8 284.7 303.3 305.7 292.4 295.5 302.0 285.4 252.1 254.4

1.58 1.12 0.84 0.85 1.39 0.87 0.75 0.79 1.18 0.95 0.85 1.00 0.90 0.84

Polímeros, 27(3) , 237-246, 2017

Second stage

T 5% (°C)

T 50% (°C)

Weight loss (%)

T2 max (°C)

S2 max (%/°C)

224.2 230.7 235.6 241.9 193.0 198.8 201.4 208.0 231.6 237.0 242.1 196.8 203.9 217.2

293.4 297.5 304.1 311.6 281.8 290.5 303.4 309.4 296.5 301.8 307.5 288.5 300.8 306.1

73.4 72.1 69.0 67.0 74.1 72.2 69.0 67.2 72.3 69.5 67.0 73.0 69.5 67.7

458.7 459.3 459.7 457.3 465.4 461.1 460.7 459.2 459.9 460.1 460.4 459.8 461.0 457.5

0.41 0.35 0.38 0.39 0.29 0.34 0.37 0.40 0.33 0.37 0.39 0.35 0.38 0.42

Residue at 480 °C (%) 13.8 11.0 11.3 11.6 12.3 10.9 11.2 11.1 10.8 11.1 11.3 9.9 11.2 11.2

241

Bouchoul, B., Benaniba, M. T., & Massardier, V. In the first step, the presence of 30 phr ESO or ESOME with 30 phr ISB or with 30phr ATBC Instead of 60 phr ISB or 60 phr ATBC retards the temperature at 5% weight loss from respectively 224.2°C and 193.0°C to 241.9°C and 208.0°C and decreases the weight loss from 73.4% and 74.1% to 67.0% and 67.2% respectively. These results are in agreement with those of Joseph et al.[28]. From Table 2, it could be observed that the value of T1 max increases with the addition of ESO or ESOME in the plasticizer system.

On the other hand, S1 max decreases.

The second stage of the degradation begins at temperatures higher than 400°C. Thermal degradation of the polyene sequences occurring during this stage yields volatile aromatic and aliphatic compounds by the intermolecular cyclization of the conjugated sequences[13].

However, the second degradation stage occurs without any marked variation in the temperature range. Hence, ESO and ESOME improve the thermal stability of PVC blends better than ISB or ATBC.

3.6 Differential scanning calorimetry (DSC) The glass transition temperature (Tg) of all the sheets with various plasticizer combinations was determined by using differential scanning calorimetry (DSC).

This temperature is an important parameter in polymer characterization to evaluate the plasticizing effect of plasticizers mixtures added to polymer systems. DSC curves of the realized formulations are displayed in Figure 4.

The analyzed plasticized PVC sheets show Tg values ranging between 62.1°C and 65.5°C. Unplasticized PVC exhibits a glass transition temperature (Tg) at 81.0 °C. We have

observed that with incorporation of plasticizers systems, Tg values decrease steadily down to 65°C. This indicates the excellent plasticizing effect of plasticizers mixtures for 60 phr by weight[29].

The data displayed in Figure 4 indicates that Tg values evolve slightly with ISB or ATBC and when ESO or ESOME quantity is increased from 10 to 30 phr in plasticizer system, which is agreement with a good miscibility of ISB, ATBC, ESO and ESOME with PVC. A low value is observed for the ATBC alone (65.5°C) but it even decreases by addition of ESO or ESOME at 30 phr to 63.6°C and 64.6°C respectively. Another low value is obtained for ISB alone (62.1°C) and it increases slightly by adding ESO and ESOME from 10 phr to 20 phr.

These data observations suggest that the length of the alkoxy group plays a critical role in controlling the available free volume, as it has been reported by Stuart et al.[29].

3.7 Dynamic mechanical thermal analysis (DMTA) DMTA was used to investigate the compatibility of the plasticizers with PVC. Generally, for an immiscible blend, the tan δ curve shows a large peak. For a highly compatible plasticizer system, the curve shows a narrow peak[14,30]. tan δ, loss modulus (E”) and FWHM (full width at half maximum)[31] of realized formulations are displayed in Table 3.

The DMTA data indicate that the Tα peak shifts to lower temperature when increasing concentrations of ISB or ATBC in the plasticized blends (Figure 5). Based on the data in Table 3, for example, the Tα peak for the ISB/ESO system decreases from 23.0°C to 11.4°C when the ISB concentration increases from 30 to 60 phr respectively. The variation of Tα is a very important factor in the evaluation of the

Figure 4. DSC curves of plasticized PVC films with: (A) ISB/ESO; (B) ATBC/ESO; (C) ISB/ESOME; (D) ATBC/ESOME. 242

Polímeros, 27(3) , 237-246, 2017

Thermal and mechanical properties of bio-based plasticizers mixtures on poly (vinyl chloride) Table 3. DMTA analysis of T drawn from E” peak (°C), FWHM (°C), Tα at different frequencies and the activation energy of PVC films with mixtures of plasticizers.

Tα (°C)

Formulation Number

T from E” (°C) 1 Hz

FWHM (°C) 1 Hz

1 Hz

3 Hz

5 Hz

10 Hz

1 2 3 4 5 6 7 8 9 10 11 12 13 14

-30.6 -28.9 -27.1 -23.3 -33.8 -31.6 -29.7 -25.8 -27.1 -24.1 -22.0 -32.9 -30.8 -17.2

51.6 52.7 58.5 60.6 45.1 48.9 53.0 54.8 52.7 53.0 56.8 47.3 47.4 49.4

11.4 18.3 21.7 23.0 6.8 10.2 17.2 21.6 14.1 19.9 23.4 11.1 15.0 19.5

15.6 31.6 12.7 27.5 28.7 23.1

18.4 33.5 13.7 30.9 30.3 24.7

19.0 35.7 15.9 35.0 35.3 28.3

Activation Energy Ea kcal. mol-1 45.02 30.88 39.10 30.50 35.70 46.57

FWHM: Full width at half maximum.

Figure 5. tan δ peaks of PVC films plasticized with various plasticizers mixture (A) ISB/ESO; (B) ATBC/ESO; (C) ISB/ESOME; (D) ATBC/ESOME.

plasticizer effect and it can be seen that the Tα values of formulations with 60 phr ISB or 60 phr ATBC are 11.4°C and 6.8°C, respectively.

This means that all the plasticizers mixtures show good plasticization of PVC. The Tα values with 60 phr ATBC or 60 phr DOP (not shown) based formulations are very close with little difference on each other (1.8°C).

However, it is known that DOP is dangerous for human health[30], so ATBC is a potentially less toxic plasticizer for PVC based compositions. Table 3 shows that the FWHM values of PVC with ATBC and ISB alone are smaller than the FWHM values of plasticized systems with ESO and Polímeros, 27(3) , 237-246, 2017

ESOME secondary plasticizers. When ISB and ATBC content increases in plasticizer systems, the FWHM values decrease; this indicates the good compatibility of ISB and ATBC with PVC. From the frequency dependence of the Tα, the values of activation energy for the plasticized samples are calculated (Table 3). It appears that there is a dependence of the activation energy on the measured Tα. It can be related to the good compatibility between PVC with ISB/ESO, ATBC/ESO and ISB/ESOME. The points corresponding to the plasticized PVC samples are shifted toward lower values of activation energy[9].

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Bouchoul, B., Benaniba, M. T., & Massardier, V. 3.8 Mechanical properties Figure 6 shows the effect of the plasticizers compositions on tensile strength and elongation at break of PVC formulations.

The tensile strength decreases by rising the amount of ISB or ATBC in the plasticizers systems, whereas the elongation at break increases with a rise in content of ESO and ESOME from 0 to 20 phr in plasticizer systems. The increase in tensile strength and decrease in elongation at break with 30 phr of ESO or ESOME in plasticizers systems is attributed to the increased polarity of the plasticizer, which growths the cohesive energy density (intensity of intermolecular attractions), so that with a higher polarity, materials tend to hold themselves together more tightly[32]. As a result, the chain mobility and thus the flexibility of the compounds are reduced, although the increase in tensile strength and decrease in elongation at break with ESO plasticizer may be attributed to its high viscosity. Reduction in tensile strength with an increase in ATBC or ISB content (60 phr) possibly result from a high plasticization efficiency compared to the one of other plasticizers[33], which allows facile sliding of polymer chains past each other[32]. Hence, compounds of ATBC/ESOME (40/20) and ISB/ESOME (50/10) give the highest elongation at break 858% and 871% respectively.

3.9 Scanning electron microscopy (SEM)

Figure 6. Mechanical properties of plasticized PVC films: (A) Tensile strength (Mpa); (B) Elongation at break (%).

Fractured surfaces of the selected formulations were investigated by SEM. The microscopic structures of PVC films plasticized with different plasticizers mixtures are displayed in Figure 7. A homogeneous and regular aspect

Figure 7. SEM micrographs of fracture surfaces of PVC films with different plasticizer compositions: (A) ISB 60 phr; (B) ATBC60 phr; (C) ISB/ESO 30/30; (D) ATBC/ESO 30/30; (E) ISB/ESOME 30/30; (F) ATBC/ESOME 30/30.

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Thermal and mechanical properties of bio-based plasticizers mixtures on poly (vinyl chloride) has been obtained, without presence of PVC aggregates, when ISB or ATBC are used alone or with ESO and ESOME (30/30), which indicates a good miscibility between PVC and plasticizer systems, responsible for the higher levels of mechanical performances of films[34]. The obtained surfaces are rather homogeneous and smooth with little roughness.

This influence on the morphology of plasticized mixtures (with ISB, ATBC, ESO and ESOME) can be attributed to a decrease in the matrix viscosity and denotes a good dispersion of plasticizers in the PVC matrix[35].

4. Conclusions Discoloration change as well as thermal degradation decreases when increasing the amount of ESO or ESOME from 10 to 30 phr in plasticized systems. The effect of the four plasticizers systems on the glass transition temperature of PVC is similar, lowering Tg from 81°C for unplasticized PVC to values in the range [62°C, 65°C] for PVC specimens plasticized with different plasticizers systems. Tα and FWHM decrease when increasing the amount of ISB or ATBC; which indicates they have a good compatibility with PVC. Tensile strength at break of PVC formulations with different plasticizers mixtures remains about constant by decreasing the ESO or ESOME level from 30 to 10 phr, the elongation at break is rather stable with ESO or ESOME levels from 20 to 10 phr.

The analysis of the morphology indicates that all the plasticizers have good compatibility with PVC. ATBC or ISB with 10 to 30 phr of ESO or ESOME are good candidates to substitute phthalates such as DEHP in flexible PVC formulations.

The use of secondary biobased plasticizers, such as ESO and ESOME with other primary biobased plasticizers as ISB and ATBC is a good way for the production of flexible PVC with low health toxicity and low environmental impacts.

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4. David, R. M., Moore, R. M., Finney, D. C., & Guest, D. (2000). Chronic toxicity of di(2-ethylhexyl) phthalate in mice.

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12. Martin, T. M., & Young, D. M. (2003). Correlation of the glass transition temperature of plasticized PVC using a lattice fluid model. Polymers, 44(16), 4747-4754. http://dx.doi.org/10.1016/ S0032-3861(03)00446-4.

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Effect of biobased plasticizers on thermal, mechanical, and Permanence Properties of Poly(vinyl chloride). Journal of Vinyl and Additive Technology, 20(4), 260-267. http://dx.doi. org/10.1002/vnl.21356.

14. Gil, N., Saska, M. & Negulescu, I., & (2006). Evaluation of the effects of biobased plasticizers on the thermal and mechanical properties of poly(vinyl chloride). Journal of Applied Polymer Science, 102(2), 1366-1373. http://dx.doi. org/10.1002/app.24132.

15. Meyer, P., Techaphattana, N., Manundawee, S., Sangkeaw, S., Junlakan, W., & Tongurai, C. (2008). Epoxidation of soybean oil and jatropha oil. Thammasat International Journal of Sciences and Technology, 13(Special edition), 1-5.

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Transesterification of vegetable oils: a review. Journal of the Brazilian Chemical Society, 9(3), 199-210. http://dx.doi. org/10.1590/S0103-50531998000300002.

22. Nihul, P. G., Mhaske, S. T., & Shertukde, V. V. (2014). Epoxidized rice bran oil (ERBO) as a plasticizer for poly (vinyl chloride) (PVC). Iranian Polymer Journal, 23(8), 599-608. http://dx.doi. org/10.1007/s13726-014-0254-7.

23. Sammaiah, A., Padmaja, K. V., & Prasad, R. B. N. (2014).

Synthesis of epoxy jatropha oil and its evaluation for lubricant properties. Journal of Oleo Science, 63(6), 637-643.

PMid:24829128. http://dx.doi.org/10.5650/jos.ess13172.

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Synthesis of plasticizers derived from palm oil and their application in polyvinyl chloride. Procedia Chemistry, 4(1), 313-321. http://dx.doi.org/10.1016/j.proche.2012.06.044.

25. Sun, B., Chaudhary, B. I., Shen, C. Y., Mao, D., Yuan, D. M., Dai, G. C., Li, B., & Cogen, J. M. (2013). Thermal stability of epoxidized soybean oil and its absorption and migration in poly(vinylchloride). Polymer Engineering and Science, 53(8), 1645-1656. http://dx.doi.org/10.1002/pen.23417.

26. Bueno-Ferrer, C., Garrigós, M. C., & Jiménez, A. (2010).

Characterization and thermal stability of poly (vinyl chloride) plasticized with epoxidized soybean oil for food packaging.

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Kinetics of degradation of PVC-containing novel neem oil as stabilizer. Journal of Applied Polymer Science, 112(4), 2151-2165. http://dx.doi.org/10.1002/app.29686.

28. Joseph, R., Madhusoodhanan, K. N., Alex, R., Varghese, S., George, K. E., & Kuriakose, B. (2004). Studies on epoxidised rubber seed oil as secondary plasticiser/stabiliser for polyvinyl chloride. Plastics, Rubber and Composites, 33(5), 217-222. http://dx.doi.org/10.1179/146580104225020974.

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29. Stuart, A., Le Captain, D. J., Lee, C. Y., & Mohanty, D. K. (2013). Poly(vinyl chloride) plasticized with mixtures of succinate di-esters – synthesis and characterization. European Polymer Journal, 49(9), 2785-2791. http://dx.doi.org/10.1016/j. eurpolymj. 2013.06.023.

30. Cadogan, D. F. (1991). Plasticizers: a consideration of their impact on health and the environment. Journal of Vinyl and Additive Technology, 13(2), 104-108. http://dx.doi.org/10.1002/ vnl.730130209.

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33. Chaudhary, B. I., Nguyen, B. D., Smith, P., Sunday, N., Luong, M., & Zamanskiy, A. (2015). Bis(2-ethylhexyl) succinate in mixtures with epoxidized soybean oil as bio-based plasticizers for poly(vinylchloride). Polymer Engineering and Science, 55(3), 634-640. http://dx.doi.org/10.1002/pen.23934.

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Received: Nov. 10, 2016 Revised: Jan. 05, 2017 Accepted: Mar. 02, 2017

Polímeros, 27(3) , 237-246, 2017

http://dx.doi.org/10.1590/0104-1428.03916

Carbon nanotube buckypaper reinforced polymer composites: a review Bruno Ribeiro1*, Edson Cocchieri Botelho1, Michelle Leali Costa1 and Cirlene Fourquet Bandeira1 Departamento de Materiais e Tecnologia, Universidade Estadual Paulista “Júlio de Mesquita Filho” – UNESP, Guaratinguetá, SP, Brazil

1

*dorado.bruno@gmail.com

Abstract This review provides valuable information about the general characteristics, processing conditions and physical properties of carbon nanotube buckypaper (BP) and its polymer composites. Vacuum filtration is the most common technique used for manufacturing BP, since the carbon nanotubes are dispersed in aqueous solution with the aid of surfactant. Previous works have reported that mechanical properties of BP prepared by vacuum filtration technique are relatively weak. On the other hand, the incorporation of polymer materials in those nanostructures revealed a significant improvement in their mechanical behavior, since the impregnation between matrix and BP is optimized.

Electrical conductivity of BP/polymer composites can reach values as high as 2000 S/m, which are several orders of magnitude greater than traditional CNT/polymer composites. Also, BP can improve remarkably the thermal stability of polymer matrices, opening new perspectives to use this material in fire retardant applications.

Keywords: BP composites, carbon nanotubes buckypaper, physical properties.

1. Introduction to CNTs Since the Discovery of carbon nanotubes (CNTs) in 1991 by Iijima, CNTs have attracted a great deal of interest due to their superior mechanical, electrical and thermal properties, which makes them an ideal candidate of nanofiller in preparation of polymer nanostructured composites[1-4]. The possibility of obtaining advanced composites with multifunctional properties has attracted the efforts of researches in both industry and academia. Industry assumes their potential applications such as nanoelectronics devices and ultra-light structural materials. Since the first report of synthesis of polymer nanostructured composites by Ajayan in 1994[5], the number of research articles related to CNTs reinforced polymer composites has increased exponentially, with more than 2000 publications in 2010[6]. On the other hand, one of the limitations for industrial application of CNTs is their high price in relation to polymer value. This barrier can be overcome when CNTs provide significant improvement

in properties of high performance polymers for high-end applications[7-9].

A carbon nanotube can be defined as cylinders composed of rolled-up graphite planes with diameters in nanometer scale. Although similar in chemical composition to graphite, CNTs are highly isotropic, and it is this topology that distinguishes nanotubes from other carbon structures and gives them their unique properties. Also, they are one dimensional carbon material which have an aspect ratio greater than 100[3,10-12]. There are basically two main kinds of CNTs: single walled carbon nanotubes (SWCNT) and multi walled carbon nanotubes (MWCNT) as illustrated in Figure 1. The first one consists of a single graphene layer rolled up into a seamless cylinder, and its diameter is around 0.5-1.5 nm[14]. On the other hand, MWCNTs is defined by two or more concentric cylindrical shells of graphene sheets coaxially arranged around a central hollow core with van der Waals forces between adjacent layers[15].

Figure 1. Schematic representation of single walled carbon nanotube (SWCNT) and multi walled carbon nanotube (MWCNT)[13].

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

Ribeiro, B., Botelho, E. C., Costa, M. L., & Bandeira, C. F.

2. Properties of CNTs

3. Dispersion of CNTs

The physical properties of CNTs are compared to other carbon materials such as graphite, diamond and fullerene in Table 1. According to the literature[3,16,17] quantum mechanics calculations can predict remarkable mechanical properties for SWCNTs. Theoretical and experimental results have demonstrated unusual mechanical properties of SWCNTs with Young’s modulus as high as 1TPa and tensile strength of 150-180 GPa. Byrne and Gun’ko[18] reported in their work that measured Young’s moduli for individual MWCNTs of between 0.27-0.95 TPa and strengths in the 11-63 GPa range.

These make CNTs the strongest and stiffest materials on earth.

Adding carbon fillers to polymers in order to improve mechanical properties and decrease weight is not a novel idea. Carbon black has been utilized to reinforce rubber and plastics[19,20]. Also, carbon fibers composites are very popular materials that have been used in airplanes, cars, bicycles, etc[21]. However, the great potential of CNTs makes them crucial materials to obtain new nanostructured products with remarkable mechanical features. For example, sports equipment, such as tennis racquets containing CNTs, have been produced and marketed. With CNTs becoming easier to produce and cheaper to buy, the CNT industry could potentially overtake that of the carbon fiber industry and become one of the major additives for polymer-composite fabrication[18,22,23].

Similar to mechanical properties, electrical conductivity of CNTs are quite varied, probably due to varying levels of defects as well as an unknown distribution of chiralities.

According to previous works[15], MWCNTs show both a metallic and semiconducting behaviors, with conductivities raging from 2 X 107 to 8 X 105 S/m. In addition, the electrical conductivity of SWCNTs can be calculated as about 5 X 107 S/m. An interesting electrical application of CNTs is their ability to work as field emitters. Field emission is a property by which a material can be induced to eject electrons simply by putting a voltage difference between it and an object. Carbon nanotubes are excellent field emitters because of their highly anisotropic nature and their small diameter.

Thermal conductivity is another property of CNTs that has been attracting great attention by several researches.

Theory predicts that MWCNTs presents thermal conductivity as high as 3000 W/mK at room temperature, which is higher than that found for cooper (385 W/mK)[10,12,15,18].

On the other hand, experimental studies[24,25] found thermal conductivity at room temperature to be significantly lower, 300 W/m K, for a single MWCNT. Simulations reveal that thermal conductivity should depend on nanotube length, increasing as nanotubes become longer. Also, obtaining measurements of this property is very difficult, since the simulations suggest nanotubes can interact with a substrate causing a reduction in the thermal conductivity[10,15].

The successful utilization of CNTs in composite applications depends on their homogenous dispersion throughout the polymer matrix. After several years of research, the full potential of CNTs as reinforcements has been limited due to the issues associated with dispersion of entangled CNT during processing and poor interfacial interaction between nanofillers and polymer matrix[26-29]. Carbon nanotubes have a tendency to form agglomerates during synthesis because of van der Waals attraction between nanotubes, leading in most cases to the formation of large agglomerates in polymer matrices, as can be seen in Figure 2. It has been proved that these bundles and agglomerates result in diminished mechanical, thermal and electrical properties of composites as compared with theoretical predictions related to individual CNTs[31-33].

Also, the processability of CNT-based composites, especially with thermoplastic matrix is not an easy task, since the high aspect ratio (>1000) is responsible for a substantial increase in viscosity of polymer, thus affecting its dispersion process. Such a behavior has been considered as one of the great challenges in obtaining CNT reinforced polymer composites, because its use is generally limited to levels lower than 5%vol in the polymer matrix[12,30].

Ultrasonication is a technique that consists in applying ultrasound energy to agitate particles in a solution. It is the most frequently used method for nanoparticle dispersion.

The equipment called sonicator (Figure 3) produces shock waves that promotes “peeling off” of individual nanoparticles located at the outer part of the nanoparticle bundles, or agglomerates, and thus results in the separation of individualized nanoparticles from the bundles[10]. This technique has been employed to disperse CNTs in liquids with low viscosity, such as water, ethanol and acetone.

However, the sonication treatment plays a crucial role during the dispersion process. If it is aggressive and/or too long, CNTs can be easily damaged, especially when a probe sonicator is employed. The localized damage to nanotubes deteriorates both electrical and mechanical properties of the CNT reinforced polymer composites[35-37].

Table 1. Physical properties of different carbon nanotubes[10].

Property Density (g/cm3) Electrical conductivity (S/cm) Thermal conductivity (W/mK) Thermal stability in air (˚C)

248

SWCNT 0.8 102-106 6000 >600

MWCNT 1.8 103-105 2000 >600

Figure 2. SEM images from fractured surfaces of phenolic resin/ CNT composites[30].

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Carbon nanotube buckypaper reinforced polymer composites: a review

4. Functionalization of CNTs In the last decade the chemical modification of CNTs has been the focus on intense research in the scientific community.

As previously mentioned in this work, CNTs exist in clusters due to van der Waals interactions that make difficult their dispersion in polymer matrix. Therefore, a major challenge in the development of nanostructured polymer composites is to obtain a satisfactory dispersion of the filler in the polymer matrix in order to maximize the properties of the final product. Also, the functionalization process appears to prevent agglomeration of the CNT, improving the interfacial adhesion between polymer and reinforcement.

Basically, the functionalization process can be divided in two groups: covalent and non-covalent functionalization, which is described below. Table 2 provides advantages and disadvantages of functionalization techniques.

oxidation, which results in the formation of carboxylic acid groups (-COOH) on the surface of the nanotubes. During the process, CNTs are refluxed with a mixture of inorganic acids (H2SO4/HNO3), sometimes with the application of high power sonication. This functionalization provides stables dispersions of CNTs in a range of polar solvents, including water[38-41]. The functionalization reaction is exemplified in Figure 4.

However, there are some issues during the covalent functionalization that have been reported. The employment of concentrated inorganic acids combined with high power sonication is responsible for creating a large number of defects on the CNTs sidewalls, and in some extreme cases, CNTs are fragmented into smaller pieces. These damaging effects can result in severe degradation of mechanical, electrical and thermal properties of CNTs[10,12].

4.2 Non-covalent functionalization

4.1 Covalent functionalization Covalent functionalization of CNTs can be achieved by either direct addition reactions of reagents to the sidewalls of nanotubes or modification of appropriate surface-bound functional groups on the nanotubes[29,38]. The most common method employed to functionalize CNTs covalently is nanotube

Non-covalent functionalization is an alternative method for improving the interfacial properties of nanotubes. Also, the process normally involves van der Waals, π-π or CH-π interactions between polymer molecules and CNT surface[42-44].

The two major approaches for non-covalent functionalization is polymer wrapping and surfactant-assisted dispersion.

A typical non-covalent functionalization is known as polymer wrapping. In this case, the suspension of CNTs in the presence of polymers, such as polystyrene[45] or poly(ether-imide)[46], lead to the wrapping of polymer around the CNTs to form supermolecular complexes of CNTs. The polymer wrapping process is achieved through the van der Waals interactions and π–π stacking between CNTs and polymer chains containing aromatic rings.

Surfactant-assisted dispersion consists to transfer CNTs to aqueous phase with the aid of surface-active molecules such as sodium dodecyl-sulfate (SDS) or polyoxyethylene

Figure 3. Sonication tip dispersing carbon nanotubes in aqueous solution[34].

Figure 4. Covalent functionalization reaction of carbon nanotubes[41].

Table 2. Characteristics of different CNT functionalization techniques[10].

Functionalization Covalent Non-covalent

Damage to CNTs Incorporation of functional groups Polymer wrapping Surfactant adsorption

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Yes No No

Interaction with polymer Strong Variable Weak

Re-agglomeration of CNTs in matrix Yes No No

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Ribeiro, B., Botelho, E. C., Costa, M. L., & Bandeira, C. F. octyl phenyl ether (Triton X-100). The physical adsorption of surfactant on the nanotubes surface reduces the surface tension of CNTs, effectively preventing the formation of aggregates. Also, the presence of an aromatic group in the surfactant molecule allows for π–π stacking interactions with the graphitic sidewalls of the nanotubes, which results in their effective coating and dispersion[47,48].

The advantage of using non-covalent functionalization is that it does not alter the structure of the nanotubes and, therefore, both electrical and mechanical properties remain unchanged. However, the efficiency of the load transfer might decrease since the forces between the wrapping molecules and the nanotube surface might be relatively weak[10,18,39,44].

5. CNT buckypaper reinforced polymer composites Polymer composites, consisting of additives and polymer matrices, including thermoplastics, thermosets and elastomers, are considered to be an important group of relatively inexpensive materials for many engineering applications. As effective nanoscale reinforcement, CNTs have attracted great interests in the field of polymer composites.

These nanomaterials displays good mechanical properties, excellent electrical and thermal conductivities, which are considered remarkable attributes for many applications in several fields of industry. However, as previously discussed in this work, their low solubility in common solvents, strong agglomerating tendency and high viscosity of CNT/polymer mixtures caused a poor dispersion and limited their practical applications. In order to solve this issues, CNT sheets, also known as buckypapers (BPs) have been employed to development of polymer nanostructured composites. BPs can be defined as a free-standing porous mats of entangled CNT ropes cohesively bounded by van der Waals interactions[49-52]. Consequently, this material is used in diverse applications such as artificial muscles[53], electrodes[54], field-emission[55], fire shields[56], and for water purification[57]. Also, BPs can be used to prepare polymer composites with uniform tube dispersion, controlled nanostructure and high CNT loading (up to 60 wt%)[58].

The most common technique used for manufacturing BPs is vacuum filtration. The procedure involves basically three steps. Firstly, a small amount of CNTs is ultrasonically dispersed in a solvent with the assistance of a surfactant. The most common solvents used to prepare BPs with good quality are N-Methylpyrrolidone (NMP) and N,N-dimethylformamide (DMF)[59]. However, using an appropriate surfactant in water can be cheaper than using DMF and NMP, which also exhibit the disadvantage of high boiling points. Sodium dodecyl-sulfate (SDS) and polyoxyethylene octyl phenyl ether (Triton X-100) have been employed as water based surfactants for manufacturing BPs by several researches[60,61]. On the second step, a vacuum-assisted filtration of a homogeneously dispersed CNT solution is carried out, using a polytetrafluoroethylene or nylon filter with submicron-sized pores. Finally, CNTs are deposited on the filter surface and form a thin membrane (buckypaper) that can be removed from the filter after drying. Figure 5 shows the buckypaper obtained by vacuum filtration technique.

The major differences between conventional CNT/polymer composites and those incorporating BPs are the carbon content, the bundle distribution and the manufacturing process. Dispersed CNT reinforced composites are usually prepared by melt-mixing[62], mixing solution[63] or in situ polymerization[64]. Also, their carbon content is generally lower than 5 wt.% and the nanotube bundles are dispersed through the matrix without forming a network. On the other hand, BP composites are manufactured by techniques such as hot-compression[65], electro-spinning[66], and intercalation[67].

These materials usually have carbon content higher than 30 wt.% resulting in a network, which acts as a skeleton.

Also, higher mechanical, electrical and thermal properties of the composites could be expected, as a result of better transfer of stress, electrons and phonons of the CNT networks.

5.1 Properties of BP/polymer composites The properties of buckypaper and its polymer composites have been attracted great attention of academic community.

Wide variations are reported in the properties, especially mechanical properties of BP/polymer composites. The mechanical

Figure 5. Representation of MWCNT buckypaper and its microstructure observed by SEM[61].

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Carbon nanotube buckypaper reinforced polymer composites: a review properties of BPs prepared by vacuum filtration technique are relatively weak, leading to a Young’s modulus and a tensile strength of 0.2-2 GPa and 2-33 MPa, respectively[68-72].

However, when polymer matrices are incorporated into BPs these properties display a significant improvement.

Han et al.[73] in a study of BP/polyurethane composites reported that both Young’s modulus and tensile strength increase dramatically by 340% and 960%, respectively, as MWCNT loading reach 46 vol%. These improvements were compared to neat BPs prepared by the authors. Similar results were found by Pham et al.[58] in a study of BP/polycarbonate composites, since the Young modulus and tensile strength increased by about 120 and 200%, respectively. Also, the mechanical properties of BP/polymer composites can be highly influenced by the processing technique employed during the consolidation of the material. Ashrafi et al.[74] compared the correlation between Young’s modulus and impregnation quality of SWCNT BP/epoxy composites.

They found Young’s modulus as high as 11.4 GPa when vacuum technique was employed, whereas for hot-press composites this value was 3.5 GPa. This behavior can be attributed to a higher quality of impregnation as well as a higher content (40-45 wt.%) of CNT than other buckypaper composites reported in the literature. In addition the Young’s modulus of BP/polymer composites is measured by dynamic mechanical analysis. Díez-Pascual et al.[75] found values of E at room temperature as high as 2.2 and 3.7 GPa for BP/PPS and BP/PEEK, respectively, which means an improvement by 38 and 32% compared to neat matrices.

As previously discussed in this work, CNTs possess high values of electrical conductivity. Materials with electrical conductivities higher than 10-8 S/cm are required for electrostatic dissipation, while for electrostatic painting and EMI shielding applications, conductivities greater than 10-6 to 10-1 S/cm, respectively are required[76-78].

The measured electrical conductivities of traditional CNT/polymer composites typically ranged from 10-5 to 10-3 S/cm above the percolation threshold[76-82]. The incorporation of CNTs within a polymer is responsible for creating a CNT network, which allows a transition behavior from a semi-conductive or conductive material. This transition is a phenomenon known as electrical percolation threshold, when conductive pathways are formed at a critical filler concentration in an insulating polymeric matrix. While further increase in CNT content above the percolation threshold can enhance marginally the electrical conductivity of composites, the solution viscosity becomes too high to produce void-free composites when the CNT content is higher than 1.0 wt.%.

The incorporation of buckypapers into polymer matrices offers an attractive route to minimize aforementioned issues.

As studied by several researches the electrical conductivity of BPs prepared by vacuum filtration process is in the range of 50-6000 S/m[58,59,75,83]. In a recent work Han et al.[83] measured electrical conductivity of BP/epoxy composites as high as 2000 S/m. This value is several orders of magnitude higher than the conductivity of conventional CNT/epoxy composites, which due to their preparation method possess a low content of CNTs. This result opens new perspectives in the field of semi conductive materials.

Carbon nanotubes are excellent thermal conductors but their use as fillers in polymeric matrices have not reached Polímeros, 27(3) , 247-255, 2017

the kind of highly thermal conductive composites that one might expect. However, since polymers are usually poor thermal conductors, with thermal conductivity on the order of 0.1 W/mK, the incorporation of carbon nanotubes still offers significant thermal-conductivity improvements in the resulting CNT/polymer composites[84-87]. According to previous works[88], thermal conductivity of epoxy-based composites reinforced with 1.0 vol.% of MWCNTs increased by more than 100% reaching a value around 0.5 W/mK. Díez-Pascual et al.[89] reported similar results for 1.0 wt.% SWCNT/PEEK composites. They found a value around 0.6 W/mK, which means an increase of 150% compared to neat polymer matrix.

Since in buckypapers and buckypaper-based composites, CNTs can form dense networks, a high thermal conductivity is expected. Gonnet and collaborators[90] found a value around 18 and 42 W/mK for the aligned and the random SWCNTs buckypaper. These values are much lower than the theoretical thermal conductivity predicted for SWCNTs and MWCNTs (6000 and 3000 W/mK, respectively).

This difference can be attributed to the high thermal resistance at nanotube/nanotube junctions[91].

Thermal conductivity of BP/polymer composites has presented similar results to conventional CNT-based composites.

Charpategui et al.[92] prepared BP/epoxy composites where the CNTs concentration was in the range of 35-60 wt.%.

The result revealed a thermal conductivity of 0.43 W/mK, which is very close to that reported in traditional CNT/epoxy composites[88]. This behavior can be attributed to the small thermal conductance of the nanotube-polymer interface, the high interfacial thermal resistance between CNTs and, by a reduction of the number of contact points between CNTs[88-92], which limit considerably the heat transfer.

Several studies have reported that only small addition of CNTs into polymers can improve the thermal stability of composites significantly, resulting in large increase of thermal decomposition temperatures by about 5-15 °C[93-96].

According to literature[97] PPS-based composites reinforced with 5.0 wt.% of MWCNTs increased thermal decomposition temperature by about 14˚C compared to neat PPS. Díez-Pascual et al.[98] found similar results for SWCNT reinforced PEEK composites at 1.0 wt.% loading.

Since BP-based composites can be produced with uniform tube dispersion and high CNTs content (up to 60 wt.%), the improvement in thermal decomposition properties should be higher to those presented in conventional CNT/polymer composites. Previously Díez-Pascual et al.[99] have manufactured SWCNT BP reinforced PPS and PEEK laminates with a CNT loading of 25%, using hot-press technique. The results revealed an increase of 62 and 45 °C for BP/PPS and BP/PEEK composites, respectively. This exceptional enhancement can be explained by different factors: Firstly, the good impregnation between both matrices and buckypaper improves the interfacial adhesion between them, thus the SWCNT can effectively act as protective barriers to prevent the transport of volatile decomposed products out of BP-based composites during thermal degradation process, resulting in the enhancement of the thermal stability of both polymer matrices. Also, thermal interfacial resistance between the CNTs and the polymer decrease in the presence of chemical 251

Ribeiro, B., Botelho, E. C., Costa, M. L., & Bandeira, C. F. bonding, resulting to an enhancement of the thermal conductivity, making easy the heat dissipation within the composite. Since CNT concentration is high the barrier effect becomes stronger and the thermal conductivity rises, leading in higher degradation temperatures[93,99]. Analogous stability effects have been reported in the literature for BP/polyimide[91] and BP/epoxy[56] composites.

6. Conclusions This review provides a comprehensive overview of the research in carbon nanotube reinforced polymer composites.

The main challenge is the development of methods to improve the nanofiller dispersion within the matrix in order to enhance mechanical, electrical and thermal properties of the resulting composites. Giving all this information, carbon nanotube buckypapers have been considered as an option to CNT agglomeration issue, resulting in composites with up to 60 wt.% of nanofiller. Ultrasonication with the assistance of a dispersant followed by vacuum filtration are the most popular techniques employed to manufacture carbon nanotube buckypapers, which is proved to be an effective way to obtain homogeneous nanotube sheets.

The impregnation between the polymer matrix and buckypaper plays a keyhole during the processing of BP/polymer composites. Good impregnation leads to an improvement in the interfacial adhesion of the composite, resulting in the upgraded mechanical properties. With regard to electrical and thermal properties, BP-based composites show important and significant results. The high electrical conductivity of the material (around 2000 S/m) is several orders of magnitude higher than the conventional CNT composites, giving them many engineering applications such as electromagnetic interference shielding materials that require conductivities above 10-1 S/cm. Also, the incorporation of buckypapers can improve dramatically the thermal stability of the polymer matrix, resulting in better flame-retardant properties.

7. Acknowledgements This work was supported by Conselho Nacional de Desenvolvimento Científico e Tecnológico (CNPq), project nº 502211/2014-8.

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Electrical and thermal conductivities of multiwalled carbon nanotubes-reinforced high performance polymer nanocomposites.

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Carbon nanotube buckypaper to improve fire retardancy of high-temperature/high-performance polymer composites.

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Received: Mar. 30, 2016 Revised: June 07, 2016 Accepted: June 26, 2016 255

http://dx.doi.org/10.1590/0104-1428.09316

Polymers and its applications in agriculture

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

Priscila Milani1, Débora França1, Aline Gambaro Balieiro1 and Roselena Faez1* 1

Laboratory of Polymer and Biosorbent Materials, Master in Agriculture and Environment Program, Universidade Federal de São Carlos – UFSCar, Araras, SP, Brazil *faez@cca.ufscar.br Graphical Abstract

Abstract Polymers are a class of soft materials with numerous and versatile mechanical and chemical properties that can be tuned specific to their application. Agriculture is an expanding area due to the requirement for indispensable food to meet the demands of a growing global population. Consequently, development of technology to improve the quality of the soil and agriculture manages is still under development. Intelligent agricultural supplies (controlled or slow release agrochemicals and superabsorbents) and biosorbents contribute to an expanding niche using technology from polymers.

This review elucidates the state-of-the-art and will discuss some important aspects of using polymers in intelligent fertilizers, as well as superabsorbent, biosorbent and biodegradation processes in agriculture that are environmentally, technically, socially, and economically sustainable.

Keywords: controlled delivery system, biosorption, superabsorbent, biodegradation.

1. Introduction The main growing, harvesting, and processing practices in agriculture were markedly achieved through the exploitation of natural resources, such as water and soil, combined with the excessive use of agrochemicals[1,2].

The high demand for fresh water, increasing manufacturing costs of agrochemicals as well as heightened awareness of their ecological adversities have rightfully shed light on these unsustainable practices. Developments in materials science and technology, however, have contributed to the

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improvement of soil conditions, nutrients, as well as water and plant management. Polymers are a class of versatile materials that have been used in many agricultural applications due to the ability to engineer application-specific polymers[3-5].

Depending on their uses, lifetime and economic viability, different classes of polymers have been used extensively in agriculture[4,6-8]. With regard to this review, our aim is to demonstrate how polymer science impacts agricultural technology and practices concerning four major applications:

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Polymers and its applications in agriculture (i) controlled delivery systems; (ii) heavy metal removal in soil and water; (iii) superabsorbents; and (iv) trends in polymer biodegradation.

Figure 1 is a panoramic overview of literature based on a bibliography search in Science Direct based on a period of ten years using “polymer agriculture”, “polymer agriculture and controlled-release”, “polymer agriculture and superabsorbent” and “polymer agriculture and biosorbent.” There is a clear increase in the quantity of studies that investigate general polymer applications in agriculture, however, the more specific and novel themes of controlled-release, superabsorbent and biosorbent are evidently publication deficient. Hence, this overview demonstrates the importance and optimism for the study on polymers and their application in agriculture.

2. Controlled Delivery of Agrochemicals by Polymeric Matrices An important point in agriculture is to supply nutritional needs of plants, which ranges from macronutrients such as N, P, K, Ca, Mg and S to micronutrients such as B, Cl, Co, Cu, Fe, Mn, Mo, Ni and Zn[9]. However, these nutrients are often not available in the environment at sufficient levels

for good plant growth. Thus, the use of agrochemicals is essential in agriculture, as fertilizers, herbicides and pesticides. On the other hand, as knowledge of the lifecycle of agrochemicals has increased, concern over how some undesirable environmental impacts, including bioaccumulation processes in the food chain and potential contamination of neighboring ecosystems have also grown more relevant[4].

The development and use of agrochemicals combined with polymeric materials has been an alternative to this problem in order to deliver agrochemicals into the soil to directly fulfill the nutritional needs of plants without causing contamination[4]. The combination of agrochemicals and polymers can be classified as physical or chemical processes.

For example, control over the rate of release is physical while how chemical is stabilized within the polymer matrix is a chemical mechanism. The biological and chemical properties of the agent and its physiochemical interactions help to direct selection of the best system to release the active agent[4]. In addition, the release profile is classified into two classes, slow- and controlled-release systems. Slow-release fertilizer (SRF) relates to the mechanism itself that delays the release, compared to controlled-release which depends on the type of mechanism to liberate the nutrient into the environment[10,11]. Controlled-release fertilizer (CRF) has

Figure 1. Literature review of polymer applications in agriculture, including controlled-release, superabsorbent and biosorbent practices.

Polímeros, 27(3) , 256-266, 2017

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Milani, P., França, D., Balieiro, A. G., & Faez, R. known factors that influence the rate and pattern of release, hence they can be controlled[10,12,13]. The pattern of release in water of slow- and controlled-release materials have been reported[12,14-16].

The concept controlled nutrient release is not a recent innovation. In 1962, Oertli and Lunt[17] published their study about controlled-release of fertilizer minerals by membrane encapsulating[18], and in 1971 Allan et al.[19] reported on controlled-release of pesticides. Their seminal work has had a dramatic influence on the field and development of the technology continues to grow as shown in Figure 1.

Controlled-release materials used to acrylic-based organic polymers, such as acrylamide[19]. Through time, environmentally green materials were proposed as alternatives to encapsulate agrochemicals, such as inorganic soil components including clay minerals[20] or calcium carbonate[21]. Recent studies have introduced the concepts of nanotechnology to improve the release mechanism through the use of nanoparticles. In chemical terms, three types of material are being used to develop this technology: organic polymers, inorganic compounds and polymers, and hybrid materials or composites[22].

As previously described, polymers that are most frequently employed to encapsulate agrochemicals are acrylamide-based gels. Other polymers include: natural rubber; polyethylene (PE); copolymers of VC-acrylic acid esters and copolymers of cyclopentadiene with a glyceryl ester of an unsaturated fatty acid; polysaccharides; and cellulose-based materials[4].

Herbicides, which are the most frequently used agrochemicals, are often employed as pendant groups in polymers. Examples include: pentachlorophenol (PCP), 2,4-dichlorophenoxyacetic

acid (2,4-D) and 4-chloro-2-methylphenoxyacetic acid[23]; fertilizer with NPK (nitrogen, phosphorous and potassium), urea, KNO3, Ca2CO3; and fertilizer with micro and macro nutrients (phosphorus, potassium, manganese, zinc, copper, molybdenum, boron and others) produced according to the agronomical needs (type of crop and soil)[4]; and pesticides as bifenthrin, tebuconazole, chlorpyrifos and atrazine[6].

Primary studies have reported that fabrication of controlled-released materials has been made exclusively with synthetic polymers. More recent studies, on the other hand, are concerned with the application of biodegradable polymers and their composites as materials for controlled‑release of agrochemicals in soil[6,24]. Table 1 shows some polymers that have been studied as matrices for controlled-release in agriculture in the past ten years.

2.1 Advantages and disadvantages of controlled-release There are many advantages of controlled-release materials such as regular and continuous nutrient supply to plants, lower frequency of applications in soil, reduced nutrient loss due to leaching, volatilization and immobilization, mitigation of root damage by high concentration of salts, greater convenience over handling fertilizers, contribution to the reduction of environmental pollution by NO3-, and improved ecological health to agricultural activity (less contamination of groundwater and surface water) and reduction in production costs[39-43].

Despite the increased development of controlled-release materials and their many advantages, use of them is much more regulated compared to conventional agrochemicals

Table 1. Polymers and agrochemical used into slow/controlled-release materials.

Agrochemical used

Polymer used

Reference

Chitosan Urea

Polyhydroxybutyrate (phb), ethyl cellulose

[25-27]

Polyethylene, polyvinyl acetate, polyurethane, polyacrylic, polylatic acid KH2PO4

Chitosan, gellan gum

[28]

Chitosan NPK

Cellulose, natural gum, rosin, waxes

[29-31]

Paraffins, ester copolymers, urethane composites, epoxy, alkide resins, polyolefines, CaH4P2O8

Chitosan

[32]

Chitosan KNO3

Chitosan-clay (montmorillonite)

[33-35]

Xanthan Paraquat ([(C6H7N)2]Cl2)

Alginate, chitosan

Hexazinone (C12H20N4O2)

Chitosan – clay

Clopyralid (C6H3Cl2NO2)

(montmorillonite)

2, 4-d (C8H6Cl2O3)

Polysaccharides

2-chloro-;4-chloro-

Cellulose, agarose, dextran, alginates, carrageenans, starch, chitosan, gelatin,

2,4,5-Trichloro- phenoxyacetates

Albumin

Validamycin (C20H35NO13)

Polystyrene, polyacrylamide, polymethylacrylate, polyamides, polyesters,

Bifenthrin (C23H22ClF3O2)

Polyanhydrides, polyurethanes, amino resins, polycyanoacrylates

[36]

[37]

[6,11,38]

Chlorpyriphos (C9H11Cl3NO3PS) Bifenthrin (C23H22ClF3O2) Azadirachtin (C35H44O16)

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Polymers and its applications in agriculture due to their high fabrication prices[33]. Furthermore, there are others disadvantages that need to be overcome, for example, said materials are known to release nutrients prematurely and quickly (burst effect), which can damage the plant or result in a non-releasing during high crop demand periods.

Further, controlled-release materials can be a problem because of unwanted residues of synthetic materials on the ground, which can affect the soil acidity[13].

Another disadvantage to the use of controlled-release might occur during the material development process due to the lack of standard methods to validate the nutrient release profile from laboratory scale to field conditions, and the difficulty to measure the impacts of the nutrients into the soil[10]. Various simulations and experimental tests have been widely investigated in order to model the release mechanism. It is also known that by Fick’s First Law of Diffusion, the mechanism of controlled-release of agrochemicals is governed by diffusion coefficients and coat thickness. In other words, there is a relationship between the rate of solute diffusion of the material and the rate of water/vapor permeation through the material[44,45].

There have been several studies regarding mathematical models of controlled-release since 1979 when Jarrell and Boersma modeled urea delivery from sulfur-encapsulated particles in soil conditions[46]. Their materials were porous and preserved with a polymeric coating. Upon application, however, the soil conditions and microorganisms of the ecosystem degraded the polymeric material, thereby exposing the more vulnerable layers. Still, nutrient diffusion through the pores was influenced by temperature and soil water content.

As described in the previous study, the release mechanism is frequently divided into three periods. First is the lag period or initial stage, wherein almost no release is detectable.

The second period occurs once a constant-release is achieved and the third period begins once a gradual decay or release is observed[12]. Still, despite the aforementioned disadvantages related to standard methods to evaluate release profiles, controlled-release materials are promising for improved agricultural management.

3. Superabsorbent Polymers and Their Applications in Soil Aside from fertilizers, all agricultural processes require a maintenance system of irrigation in order to provide enough nutrients to the soil. In arid areas, irrigation processes are particularly expensive, thereby making certain polymer materials unfeasible. As a result, the use of superabsorbent polymer (SAP) hydrogels has become an attractive alternative to circumvent this problem because they supply the necessary cost reductions and continuous irrigation characteristic of this procedure. In the presence of water, SAPs are able to swell and hold aqueous solutions, maintaining the humidity of the soil for longer periods. Surprisingly, the utilization of SAPs for agricultural processes has only been recently discussed in the Academy[47].

Literature extensively describes the use of SAPs for biological applications, such as contact lenses[48-50], baby diapers and female hygienic tampons[51,52]. Most recently, SAPs have appeared as an excellent alternative Polímeros, 27(3) , 256-266, 2017

to drug encapsulation for delivery systems[53-55]. Previous applications derived from the swelling property of SAPs, which guarantees large capacity of water absorption, have fortified the potential for biocompatibility[56]. Erickson[57] first reported the application of a SAP as a soil conditioner.

Redenbaugh[58] in turn studied its functionality to encapsulate and slowly release agrochemicals in the soil. Moreover, the origin of SAPs is closely related to its applicability; i.e.

SAPs’ properties deserve further research and development as a promising field of study.

As a positive contribution to this expanding field, the aim of this review is to provide an overview of SAPs, focusing in particular on their agricultural applications. For this purpose, we will briefly describe the characteristics of a SAP, and explore literature regarding its use as a soil conditioner and its application for controlled-release of fertilizers.

3.1 Superabsorbent polymer hydrogels Superabsorbent Polymer (SAP) hydrogel is defined as a lightly crosslinked three-dimensional polymer network, usually composed of ionic monomers, whose property of interest is the capacity of swelling in the presence of aqueous or biological fluids[59]. The density of chain crosslinking creates an expressive amount of free volume between the polymer network, which in turn, combined to the presence of huge number of hydrophilic groups, can absorb and hold large quantity of water, on the order of 10 to 1000 g/g (1.000-100.000%)[60].

The water holding and swelling capacity of a SAP can be explained by a multi-step mechanism. The first step comprises the hydration of hydrophilic groups present in the polymer network with strong bonds to water. Secondary weaker bonds are made with water by the interaction between water and the exposed hydrophobic groups. As a next step, physical or chemical crosslinks minimize the effect of osmotic forces that would lead to infinite dilution of the polymer network, and for this reason, the network retains additional water[59] .

SAPs are mainly classified according to their origin, crosslinking nature, and their responsive mechanism.

Considering its origin, a SAP can be natural, synthetic or a hybrid. Regarding its crosslinking structure, a SAP is chemically crosslinked, which forms covalent bonds between the polymer networks resulting in a permanent state, or physical crosslink in which polymer networks are bonded through intermolecular forces, such as ionic interactions, hydrogen bonds, or Van der Walls forces. The responsive mechanism of a SAP is either conventional or induced by an external stimulus. A conventional SAP will not have altered swelling equilibrium under environmental stimulus, for example, changes in pH, temperature, or electric field. On the other hand, Samchenko et al.[61] calls a stimulus responsive SAP a smart or intelligent hydrogel, as these materials adapt their swelling equilibrium to environmental conditions.

Applications based on reversible swelling/deswelling can be modulated based on chemical, biochemical or physical responses and are dependent on environmental conditions, such as high ionic conductivity, high permeability and sorption capacity.

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Milani, P., França, D., Balieiro, A. G., & Faez, R. 3.2 Superabsorbent polymer hydrogels as soil conditioners The appearance of SAP hydrogels as soil conditioners dates back to the 1980s when the first patents describing the application of hydrogels were published. Erickson[57] patented a method for improving the water retention and aeration capacity of soil matrixes by inserting polymeric films with water-swelling properties in soil. The films could be produced from several polymer-based materials, such as alkyl acrylate (i.e. methyl, ethyl, propyl and hexyl acrylates), alkyl methacrylates (including butyl, hexyl, octyl, and decyl methacrylate), and omega hydroxyalkyl acrylates (comprising 2-hydroxyethil acrylate and hydroxymethyl acrylate).

The discovery of this film has presented alternatives to the previous use of absorbent powder added directly to the soil, which can cause noteworthy problems as the particles shift positions with soil management, causing sealing effects in the soil. By adding the absorbent polymer in film format, it can be twisted, crumpled, or chopped and positioned in the soil in a way to avoid sealing effects and improve its water retention capacity[62].

Callanghan et al.[62] tested polyacrylamide and poly(vinyl alcohol) as superabsorbent polymers applied on drought affected soils in Sudan. They proved that poly(vinyl alcohol) was able to increase the field capacity of the soil by 22%, and both polymers increased the period of plant survival in the conducted nursey experiments, providing evidence for the economic importance of SAPs. Aside from polyacrylamide and poly(vinyl alcohol), Woodhouse and Johnson[63] tested a starch-copolymer on the growth of barley and lettuce in coarse soil. Their observations were that as the percentage of SAPs for water retention increased, the longer the period until plants reached their wilting point.

Stahl et al.[64] was the first to study the rate of biodegradability of superabsorbent copolymers made from polyacrylamide and polyacrylate as well as pure polyacrylate in soils. They estimated the degree of biodegradability based on the mineralization of the polymers, which was measured by the release of CO2 produced by biodegradation reaction. As Shuterland et al.[65] had previously proven, microorganisms naturally present in the soil were not able to biodegrade the introduced polymers. For this reason, their research was conducted in a way to investigate potential microbes in order to maximize the rate of biodegradability.

In the past few years, research has been performed in order to investigate different kinds of polymeric composites that would increase the efficiency of SAPs in improving soil conditions. Islam et al.[66] for example, used granular polyacrylamide under moderate and deficient irrigation conditions. They proved that the presence of SAPs was able to increase irrigation’s efficiency in 8.1% under moderate exposure to drought, and 15.6% under higher drought condition. Moreover, they noticed that the presence of granular polyacrylamide was able to reduce the activity of antioxidant enzymes, which under drought stress produced oxygen radicals responsible for harming the plant’s capacity of carbon fixation and, consequently, its growth. Therefore, in the presence of polyacrylamide SAPs, plant growth was satisfactory even under lower irrigation conditions[66].

260

Sharma et al.[67] described the synthesis of interpenetrating polymer networks based on poly(aspartic acid) whose characteristic of interest is the elevated capacity of biodegradation.

Furthermore, they suggest that improving the synthetic route through the incorporation of biochemical cross-linkers, or charged functional groups to hydrophilic biopolymers would increase the material’s capacity to biodegrade as well as increase its absorbency. However, neither of those synthetic routes resulted in a fully biodegradable material.

Cannazza et al.[68] in turn, applied cellulose-based SAPs as an environmentally friendly alternative to acrylate-based SAPs. They used derivatives of carboxymethyl cellulose sodium salt and hydromethyl cellulose as starting materials for the synthesis of the cellulose-based hydrogels. Posterior tests showed an efficiency of swelling, water retention and soil conditioning very similar to those obtained with acrylate-based SAPs. Dragan[69] described the studies of interpenetrating polymer networks (IPN) hydrogels based on biopolymeric materials, such as, chitosan, alginate, starch, and other polysaccharides.

Chauhan and Mahajan[70] apud Dragan[69] described the use of IPN based on a combination of biopolymers with poly(methacrylamide) for metal ion sorption in aqueous solution. They noticed that the sorption capacity of the material drastically increased after inserting cellulose derivatives in the polymer network.

Besides soil conditioning or metal ion sorption, SAPs can also be employed for slow- release of fertilizers in the soil. Zhan et al.[71] reported the synthesis of a SAP combining groups of polyvinyl alcohol with phosphoric acid (H3PO4) through esterification. In this case, the material had the property of not only absorbing and retaining water, but also of releasing phosphate fertilizer. However, the study only tested the swelling capacity and release properties in an aqueous environment, not in soil, the latter showing a release of 79% of the material in 28 days[71]. The mechanism of release could be explained for the swelling of the hydrogel material when immersed in water, such as the hydrolysis of pendant phosphate groups of the polymer network.

Consequently, the dissolved phosphate groups diffused out of the hydrogel due to a difference of concentration gradient inside and outside the material[56].

Liang et al.[72] on the other hand, tested the entrapment of NPK inside a poly (acrylic acid co-acrylamide)/kaolin clay superabsorbent composite. The addition of clay to the synthetic polymer was an attempt to reduce costs and improve the material’s swelling capacity. In the described work, Liang et al.[72] also tested the swelling capacity and the release mechanism in aqueous medium. They established a temperature dependent release mechanism showing that as temperature increased, the higher the solubility of the nutrient and the faster its release.

Although the addition of clay to a synthetic material can enhance its properties and diminish costs of production, environmental concerns caused by non-biodegradable materials added to the soil are still issues to overcome. In this context, researchers are focused on finding environmental friendly materials that can fulfill human economic necessities without amplifying waste accumulation. Jamnongkan and Kaewpirom[73] for example, synthetized and tested three Polímeros, 27(3) , 256-266, 2017

Polymers and its applications in agriculture different hydrogels: the first based on poly(vinyl alcohol), the second based on a combination of poly(vinyl alcohol) and chitosan and the last composed only of chitosan. They tested the swelling capacity of the three materials in soil and the hydrogel made of PVA alone was the one that guaranteed the best soil conditioning. With regard to the release capacity of potassium in the soil, the pure chitosan hydrogel was most promising.

At this point the reader may recognize the necessity to advance the use of hydrogels in agriculture, as well as the discovery and synthesis of environmental friendly materials, such as biopolymers, with high biodegradability in soil.

Although the necessity to maintain impressive swelling capacity with combined fertilizer release properties still constitutes diverse challenges, the new era of chemistry research guided by green chemistry principles can lead to novel discoveries in this immerging area of materials and agriculture science.

4. Metal Ions Removal from Water and Soils Metal ions can be classified into three different categories: toxic metals, such as Hg, Cr, Pb, Zn, Cu, Ni, Cd, As, Co, and Sn; precious metals, including Pd, Pt, Ag, Au, and Ru; and radioactive metals, like U, Th, Ra, and Am[74]. Some of these metals can be coupled to agrochemicals, such as to fertilizer, insecticides and herbicide, and then released into the soil. The soil itself, with no addition of agrochemicals, is an essential and natural source of nutrients for plants and other living organisms, consequently, trace amounts of metal ions occur naturally in the soil’s composition, functioning as important micronutrients. These metal ions are mainly classified into essential micronutrients (Fe, Zn, Mn, Ca, Na, K, Cu, and Mo), beneficial micronutrients (Co, Ni, and V), and non-essential micronutrients (Cd, Cr, Hg, and Pb)[75,76].

However, the enforcement of intensive agriculture has resulted in abusive techniques that are combined with the massive use of agrochemicals solely to increase profit and productivity. Due to this fact, an enormous amount of organic and inorganic toxic contaminants, namely metal ions, are deposited daily into soils, rivers, and groundwater[77]. Such problems with metal ions are threatening because these ions are cumulative and do not degrade naturally in the environment. Hence, their toxicity can cause severe ecological damage that can be harmful to human, other animals, insects, etc. through the consumption of contaminated water and/or food[78]. In this context, concern over soil and water pollution has stimulated scientists to develop techniques that are fundamentally aimed to circumvent or mitigate ecological harm as well as cleanse current ecosystems that are already threatened[79]. Perusing the literature, one might notice the abundance in the quantity of publications that address methods for the removal of heavy metal pollution in aqueous environments, compared to a few articles that address techniques for decontamination of soils.

One method for the recovery of contaminated areas involves polymeric biosorption: biosorption being a biological property characteristic of microbial biomass, which has been adopted into biomimetic-type techniques for the binding heavy metals even from very dilute solutions. Biosorption Polímeros, 27(3) , 256-266, 2017

is a technique of bioremediation whose application has been studied in both water and soil systems as an alternative for metal ions removal. This technique is able to remove either organic or inorganic pollutants through ionic exchange.

Research results highlight the high capacity of sorption and the low cost of this method[80].

Besides biosorption, polymers have also been commonly used as sorbent materials of toxic elements arising from agrochemicals as a way to decontaminate water and soils[81].

To better understand the mechanism of sorption by a polymer it is necessary first to identify the functional groups that are present in the material. Metal ions are absorbed on the surface of the polymeric material due to their chemical affinity for certain functional group[82]. Polysaccharides such as cellulose, chitin, chitosan and exobiopolymers are a group of polymers most commonly applied for the purpose of metal sorption due to their ionizable carboxylic acid, alcohol, amine and amide functional groups for ionic bonding[81,83,84].

Additionally, natural fiber extracted from sugarcane bagasse, green coconut fiber, and bamboo fiber, have also been applied as biosorbents. The main advantage of these materials is they are generally industrial and/or agricultural residues, low cost, and easily recovered. Such crude fibers are mainly composed of cellulose and hemicellulose which possess acidic groups responsible for the sorption of metal ions[85].

Focusing on the metal decontamination of soils as phytoremediation and rhizoremediation, research involving biopolymers, such as polysaccharides, as adsorbents is still ongoing and the development of materials capable of decontaminating affected areas with high efficiency and low cost, is still active[86].

Pal and Paul[87], in turn tested the use of exopolymer films for the remediation process of contaminated water.

Exopolymers are extracellular substances originating from microorganisms in the form of biofilms, which are capable of absorbing metal ions through electrostatic interactions.

Their results showed that zinc, copper, chromium, cadmium, cobalt, nickel and CrO metals were absorbed most efficiently.

In aqueous environments, the methods of metallic ions removal have thus far been based on biological treatments, such as filtration, electrochemical treatment, chemical precipitation, ion exchange, membrane technologies, and adsorption over active coal[74,87]. However, current studies are investigating novel applications for the use of recycled biomass as feedstock for the development of heavy metalabsorbing biomaterials in aqueous environments. Such agricultural residues include sugarcane bagasse, which has gained much attention for application in this field due to its abundance in several regions from Brazil[85]. Table 2 presents other commonly studied biomasses and their sorption capacity.

Besides soil, Table 2 shows the biomass and polymer constituent relationship and its relation with the metal ions in water adsorption. It can be observed that among the materials used as biosorbents, cellulose is the most commonly used metal ion “sequestrant,” because of cellulose’s abundant surface hydroxyl groups that can be functionalized to possess more polar or ionizable pendant groups for enhanced ion exchange for decontaminating water.

261

Milani, P., França, D., Balieiro, A. G., & Faez, R. Table 2. Examples of biomass feedstock and associated polymer compositions for ion removal.

Biomass Sugarcane bagasse ash Bagasse sugar beet Green coconut shell Chitosan Sawdust Microalgae: Spirulina platensis Rice straw Banana peel

Cotton fiber Scolymus hispanicus L (Plant) Fungus - Fusarium verticillioide

Polymer

Sorption capacity* 36.3 mg g-1 (Cu) 45% cellulose, 28% hemicellulose and 18% lignin 41.31 mg g-1 (Cr) Cellulose and pectin 185 mg g-1 Fibers (polysaccharide) 22.96 mg g-1 polysaccharide 36.8 mg g-1 26 mg g-1 (Pb) 22.5 mg g-1 (Cu) polysaccharide 19.75 mg g-1 (Zn) 100.39 mg g-1 Cellulose – polysaccharide 30.0 mg g-1 (Cr) Cellulose, hemicellulose and lignin 22.5 mg g-1 (Cu) 2.18 mg g-1 (Pb) Cellulose, hemicellulose, lignin and pectin 5.71 mg g-1 (Cd) 6.12 mg g-1 (Cu) 4.53 mg g-1 (Zn) Cellulose, hemicellulose and lignin 8.22 mg g-1 (Cd) 21.62 mg g-1 (Pb) Cellulose – polysaccharide 54.05 mg g-1 50.50 mg g-1 (Mg) polysaccharide 92.59 mg g-1 (Ca)

ions removed Cu(II) Cr(III) Tl(II) Cr(VI) Pb(II) Pb(II) Cu(II) Zn(II) Cr(VI) Cr(III) Cu(II) Pb(II) Cd(II) Cu(II) Zn(II) Cd(II) Pb(II) Cd(II) Mg(II) Ca(II)

Reference [88] [89] [90] [91]

[92]

[93] [94]

[95]

[96]

[97] [98]

*Maximum Amount of Solute Retained in the Adsorbent in Balance (qe).

As mentioned before, several materials are being applied as biosorbents for removal of metallic ions from aqueous environments. Hence, literature frequently presents tests of new materials for metal depollution of hydric resources whose properties are developed in way to increase efficiency and lower costs. On the other hand, publications regarding the research and development of materials for soil remediation are still scarce. The reasoning for this may be explained by the fact that removal of heavy metals from soil may also adversely impact micro or macronutrient populations for plants, which causes the vegetables’ roots to absorb or adsorb metal ions according to its necessity through a mechanism called phytoremediation.

5. Trends of Biodegradable Polymer materials in Agriculture Application The degradation of organic chemicals in the environment is known as biodegradation, in which the compounds are mainly converted into mineralized forms and redistributed through carbon, nitrogen and sulphur cycles by micro‑organism activities[99,100]. Natural and/or biodegradable polymers are still a requisite to overcome some problems concerning environmental aspects of polymer applications in agriculture.

As seen before, synthetic polymers are most commonly used for agricultural applications, either in intelligent agrochemicals or superabsorbents, which severely creates large quantities of non-biodegradable waste and soil contamination.

Polymer biodegradation is a very comprehensive and recent area of research; however, the connection with the polymers application in soil on agriculture context is still incipient in the literature. Several reviews with reference to degradation processes of polymers in soil and aqueous medium have been published[101,102]. However, few discussions 262

have accomplished polymer degradation from agrochemicals encapsulated in a polymer matrix.

Additionally, some difficulties have been pointed out to find a polymer with both long accurate release control and a high rate of biodegradation under the mild conditions of agricultural fields[100]. Some reports present the biodegradation matrix used in controlled-release of drugs[103,104]. Watanabe et al.[100] studied the influence of the Tg on release and biodegradation of urea coated with polycaprolactone based polyurethanes. According to the authors the biodegradation characteristic was affected by the glass transition temperature and the recrystallization of polyurethane membrane. Nevertheless, few works have dealt with the relationship between polymer applications in agriculture and their biodegradation in soils under natural conditions. Ge et al.[101] used biodegradable polyurethane as ammonium sulfate fertilizer coating and subsequent biodegradation studies were performed. They suggested that the contribution of the PU matrix on the initial release stage occurs by the diffusion process and an additional release is associated with degradation in the last stage. However, they did not perform the release and biodegradation tests in the same conditions to better understand the whole process.

Saruchi et al.[105] evaluated the influence of acrylic acid based hydrogel on the soil fertility. Additionally, they investigated the release properties of potassium-containing hydrogel; however, no information was discussed about the polymer biodegradation in a presence of fertilizer.

Biodegradation studies have also been performed on superabsorbent polymers. Ye et al.[106] prepared a degradable superabsorbent material based on a silicate/acrylic-based polymer. They proposed a hydrolysis process from Si‑O-C bonds in aqueous solution which degraded the hybrid hydrogels in a matter of days. Wilske et al.[107] characterized the biodegradability of SAP based on polyacrylate for Polímeros, 27(3) , 256-266, 2017

Polymers and its applications in agriculture four agricultural soils and at three temperatures. Detailed results suggest that the polyacrylate main chain degraded in the soils, if at all, at rates of 0.12-0.24% per 6 months[93].

Further, Sthal et al.[64] studied the rate of biodegradability of superabsorbent copolymers made out of polyacrylamide and polyacrylate, and pure polyacrylate in soils. They estimated the degree of biodegrability based on the mineralization of the polymers, measuring the released CO2 produced by biodegradation reactions[64]. As Shuterland et al.[65] had previously proved, the microorganisms naturally present in the soil were not able to biodegrade the added polymers.

For this reason, the following research was conducted in a way to investigate potential microbes in order to maximize the rate of biodegrability.

6. Final Consideration and Perspectives Based on the survey of topics presented in this study, we conclude that there is ample opportunity to enhance the application of polymers in agriculture. It is imperative that the use of polymers be in a way that is environmentally, technically, socially, and economically sustainable in practice.

Key opportunities exist to produce intelligent polymers with biodegradable and renewable properties for various agricultural applications.

7. Acknowledgements The authors thank FAPESP for financial support (Proc. 2014/06566-9) and CNPq and Capes for a scholarship.

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Received: July 15, 2016 Revised: Nov. 19, 2016 Accepted: Mar. 01, 2017

Polímeros, 27(3) , 256-266, 2017