Polímeros VOLUME XXVIII - Issue II - Apr./May, 2018
São Paulo 994 St. São Carlos, SP, Brazil, 13560-340 Phone: +55 16 3374-3949 Email: abpol@abpol.com.br 2018
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ISSN 0104-1428 (printed) ISSN 1678-5169 (online)
P o l í m e r o s - I ss u e I I - V o l u m e X X V I I I - 2 0 1 8 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 Antonio Aprigio S. Curvelo (USP/IQSC) - President
Editorial Committee Sebastião V. Canevarolo Jr. – Editor-in-Chief
Members 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) Edvani Curti Muniz (UEM/DQI) 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) Marco-Aurelio De Paoli (UNICAMP/IQ) Osvaldo N. Oliveira Jr. (USP/IFSC) Paula Moldenaers (KU Leuven/CIT) Raquel S. Mauler (UFRGS/IQ) Regina Célia R. Nunes (UFRJ/IMA) Richard G. Weiss (GU/DeptChemistry) Rodrigo Lambert Oréfice (UFMG/DEMET) Sadhan C. Jana (UAKRON/DPE) Sebastião V. Canevarolo Jr. (UFSCar/DEMa) Silvio Manrich (UFSCar/DEMa)
A ss o ci at e E d i t o r s Adhemar C. Ruvolo Filho Alain Dufresne Bluma G. Soares César Liberato Petzhold José António C. Gomes Covas José Carlos C. S. Pinto Paula Moldenaers Richard G. Weiss Rodrigo Lambert Oréfice
Sadhan C. Jana
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: May 2018
Financial support:
Polímeros / Associação Brasileira de Polímeros. vol. 1, nº 1 (1991) -.- São Carlos: ABPol, 1991Available online at: www.scielo.br
Quarterly v. 28, nº 2 (Apr./May 2018) ISSN 0104-1428 ISSN 1678-5169 (electronic version)
Website of the “Polímeros”: www.revistapolimeros.org.br
1. Polímeros. l. Associação Brasileira de Polímeros. Polímeros, 28(2), 2018
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I I I I I I I I I I I I I I I I I
Editorial Section News....................................................................................................................................................................................................E3 Agenda.................................................................................................................................................................................................E4 Funding Institutions.............................................................................................................................................................................E5
O r i g in a l A r t ic l e Cellulose nanomaterials: size and surface influence on the thermal and rheological behavior Marcos Mariano, Nadia El Kissi and Alain Dufresne....................................................................................................................................... 93
Effect of compatibiliser on the properties of polypropylene/glass fibre/nanoclay composites Normasmira Abd Rahman, Aziz Hassan and Javad Heidarian....................................................................................................................... 103
Hybrids membranes with potential for fuel cells – Part 3: extruded films of nanocomposites based on sepiolite and PC/sulfonated PC blends Ana Catarina de Oliveira Gomes, Eduardo Henrique Backes, Adhemar Colla Ruvolo Filho, Caio Marcio Paranhos, Fábio Roberto Passador and Luiz Antonio Pessan.......................................................................................................................................... 112
Fabrication of poly(lactic acid) incorporated chitosan nanocomposites for enhanced functional polyester fabric Zulfiqar Ali Raza and Faiza Anwar.................................................................................................................................................................. 120
Properties of barrier shrink bags made with EVOH and polyamide for fresh beef meat preservation Jose Boaventura Rodrigues, Kleber Brunelli, Claire Isabel Grígoli de Luca Sarantopoulos and Lea Mariza de Oliveira........................... 125
Potential doxorubicin delivery system based on magnetic gelatin microspheres crosslinked with sugars Josefa Souza, Manoel Silva and Marcos Costa............................................................................................................................................... 131
Ultrasound-assisted synthesis of polyacrylamide-grafted sodium alginate and its application in dye removal José Manoel Couto da Feira, Jalma Maria Klein and Maria Madalena de Camargo Forte.......................................................................... 139
Preparation and characterization of composites from plastic waste and sugar cane fiber Ricardo Yoshimitsu Miyahara, Fábio Luiz Melquiades, Ezequiel Ligowski, Andressa do Santos, Silvia Luciana Fávaro and Osmar dos Reis Antunes Junior................................................................................................................................................................ 147
Ultrasound assisted miniemulsion polymerization to prepare poly(urea-urethane) nanoparticles André Eliezer Polloni, Alexsandra Valério, Débora de Oliveira, Pedro Henrique Hermes de Araújo and Claudia Sayer............................. 155
Effect of heat cycling on melting and crystallization of PHB/TiO2 compounds
Nichollas Guimarães Jaques, Ingridy Dayane dos Santos Silva, Manoel da Cruz Barbosa Neto, Andreas Ries, Eduardo Luis Canedo and Renate Maria Ramos Wellen................................................................................................................................. 161
The effect of molecular weight and hydrolysis degree of poly(vinyl alcohol)(PVA) on the thermal and mechanical properties of poly(lactic acid)/PVA blends Iván Restrepo, Carlos Medina, Viviana Meruane, Ali Akbari-Fakhrabadi, Paulo Flores and Saddys Rodríguez-Llamazares...................... 169
Polysaccharides of red alga Gracilaria intermedia: structure, antioxidant activity and rheological behavior Joana Paula Lima de Castro, Luís Eduardo Castanheira Costa, Maísa Pessoa Pinheiro, Thiago dos Santos Francisco, Pedro Hermano Menezes de Vasconcelos, Lizandra Mistrello Funari, Renata Moschini Daudt, Gustavo Ramalho Cardoso dos Santos, Nilo Sérgio Medeiros Cardozo and Ana Lúcia Ponte Freitas.......................................................................................................................... 178
Cover: SEM micrographs of PP/Glass fibre composite (pg 108); PLA treated polyester fabrics (pg 122); Gelatin microspheres (pg 136). Arts by Editora Cubo.
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Polímeros, 28(2), 2018
Braskem’s UTEC Extends Global Ultra High Molecular Weight Polyethylene (UHMWPE) Innovation and Client Reach
New Polymer Manufacturing Process Slashes Energy Use By 90%
Braskem highlights its first full year of UTEC Ultra High Molecular Weight Polyethylene (UHMWPE) production at its La Porte, Texas production facility by introducing new market leading innovation, including its new UTEC 9540, the highest molecular weight polyethylene polymer commercially available worldwide. Braskem, the largest thermoplastics resins producer in the Americas and the worldwide leader in bio-polymers, serves UHMWPE clients globally with increased flexibility and product availability from its UTEC facilities in the U.S. and Brazil, which have a combined annual UHMWPE production capacity of 70 kilotons. Braskem sells high-performance UHMWPE under the trade name UTEC, developed and produced through Braskem’s proprietary technologies. Christopher Gee, UTEC Global Director, commented, “As we celebrate our first full-year of UTEC production at our La Porte, Texas facility, and extend our product portfolio, we have strengthened Braskem’s clear position as a market leader for UHMWPE clients worldwide. The successful UTEC technology transfer from our operations in Brazil to the U.S. has allowed Braskem to increase UTEC production capacity, expand its global footprint, and enhance client responsiveness. With global demand for UHMWPE increasing, including a rising presence in high-growth markets such as energy storage, we feel confident that we made the right decision to invest in this technology. We look forward to expanding our leadership position in the market through new product offerings and strategic partnerships. The future for UTEC has never been brighter.” UTEC 9540 has a molecular weight over 50 times higher than High Density Polyethylene (HDPE) resins and is used in applications that require extreme abrasion resistance. This latest addition to Braskem’s UTEC portfolio extends the depth and breadth of products available to a wide range of applications including RAM extruded and compression molded sheets, rods, and profiles. Braskem’s UTEC is an engineered polymer with excellent mechanical properties, such as high abrasion resistance, impact strength, and low coefficient of friction. It is a self-lubricating, high-strength, lightweight machinable product used for semi-finished goods. Eight times lighter than steel and lasting ten times longer than HDPE, UTEC is utilized across a vast array of applications in the following industries: construction, agriculture, material handling, transportation, textile, pulp and paper, food and beverage, mining, marine, porous plastics, oil and gas, high performance fibers, energy storage, and waste water treatment. UTEC has products spanning a range of molecular weights and particle sizes. The molecular weight begins in the low range (3 million g/mol) and extends to a high range (7 to 11 million g/mol).
The essence of commerce is selling stuff, but before you sell stuff you have to make stuff. Sometimes, making stuff consumes a lot of energy. Finding new ways to lower energy usage is important to reducing the carbon footprint of the products you manufacture. A case in point is manufacturing commercial aircraft and large vehicles. The process of curing just one section for such vehicles can consume over 96,000 kilowatthours of energy and produce more than 80 tons of carbon dioxide. That’s according to Scott White, one member of a team of researchers at the University of Illinois at Urbana-Champaign. That’s roughly the amount of electricity needed to supply nine single family homes for one year, according to the US Energy Information Administration. “The airliner manufacturers use a curing oven that is about 60 feet in diameter and about 40 feet long — it is an incredibly massive structure filled with heating elements, fans, cooling pipes and all sorts of other complex machinery,” White says. “The temperature is raised to about 350 degrees Fahrenheit in a series of very precise steps over a roughly 24-hour cycle. It is an incredibly energy-intensive process.” The researchers say they have found a way to make heat set polymer parts for cars, trucks, buses, and airplanes that uses one tenth as much electricity. “There is plenty of energy stored in the resin’s chemical bonds to fuel the process,” Moore said. “Learning to unleash this energy at just the right rate — not too fast, but not too slow — was key to the discovery. “By touching what is essentially a soldering iron to one corner of the polymer surface, we can start a cascading chemical-reaction wave that propagates throughout the material,” says White. “Once triggered, the reaction uses enthalpy, or the internal energy of the polymerization reaction, to push the reaction forward and cure the material, rather than an external energy source. This development marks what could be the first major advancement to the high-performance polymer and composite manufacturing industry in almost half a century.” The team has demonstrated that this reaction can produce safe, high quality polymers in a well controlled laboratory environment. Because it is compatible with commonly used fabrication techniques like molding, imprinting, 3-D printing, and resin infusion, the researchers envision the process being applicable to large scale production, according to Science Daily. The research findings have been published recently in the journal Nature. For those interested in electric cars, trucks, and buses that have a lower well-to-wheel carbon footprint and have longer range due to reduced weight, the new manufacturing process is very welcome news.
Source: Braskem - www.braskem.com/UTEC
Source: CleanTechnica - https://cleantechnica.com
Polímeros, 28(2), 2018 E3
N N N N N N N N N N N N N
July
A A A A A A A A A A A A A A A A A A A A A
IUPAC World Polymer Congress (Macro 2018) Date: July 1-5, 2018 Location: Cairns – Australia Website: www.macro18.org Polymer Physics - Connecting Fundamentals to Broad Applications in Polymer Physics Date: July 21-22, 2018 Location: South Hadley - USA Website: www.grc.org/polymer-physics-grs-conference/2018 Polymer Physics - New Developments in Hierarchical Structure and Dynamics of Polymers Date: July 22-27, 2018 Location: South Hadley - USA Website: www.grc.org/polymer-physics-conference/2018
August 6th International Conference & Exhibition on Advanced & Nano Materials (ICANM 2018) Date: August 6-8, 2018 Location: Quebec - Canada Website: icanm2018.iaemm.com 3rd International Conference on Material Engineering and Smart Materials (ICMESM 2018) Date: August 11-13, 2018 Location: Okinawa - Japan Website: www.icmesm.org Interplast Date: August 14-17, 2018 Location: Joinville – Brazil Website: www.interplast.com.br 25th Bio-Environmental Polymer Society (BEPS 2018) Date: August 15-17, 2018 Location: New York – USA Website: www.beps.org/meetings 12th European Symposium on Thermal Analysis and Calorimetry (ESTAC12) Date: August 27-30, 2018 Location: Brasov – Romania Website: estac12.org 5th International Conference and Exhibition on Polymer Chemistry Date: August 27-28, 2018 Location: Toronto – Canada Website: polymer.conferenceseries.com
September 4th International Conference on Bio-based Polymers and Composites (BiPoCo 2018) Date: September 2-6, 2018 Location: Balatonfüred - Hungary Website: bipoco2018.hu 10th Conference of Modification, Degradation and Stabilization of Polymers (MoDeSt2018) Date: September 2-6, 2018 Location: Tokyo - Japan Website: biz.knt.co.jp/tour/2018/modest/index.html 2nd Polymer Testing & Analysis Date: September 11-12, 2018 Location: Pittsburgh - USA Website: www.ami.international/events/event?Code=C0916 5th Symposium on Innovative Polymers for Controlled Delivery (SIPCD 2018) Date: September 14-17, 2018 Location: Suzhou - China Website: www.sipcd.com/C911 17th Brazilian MRS meeting (SBPMat 2018) Date: September 16-20, 2018 Location: Natal - Brazil Website: www.sbpmat.org.br/17encontro
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Polymers in Flooring Date: September 20-21, 2018 Location: Atlanta - USA Website: www.ami.international/events/event?Code=C911 WEASC International Conference on Advances in Applied Sciences, Engineering Technology, Polymers, Plastics & Robotics Date: September 22-23, 2018 Location: Athens - Greece Website: world-easc.com/aepr-sep-2018 International Conference on Engineering Polymers and Plastic, Building Design and Computer Sciences (AETL 2018) Date: September 24-25, 2018 Location: Bangkok - Thailand Website: aetleducation.com/eppbc-bangkok-event-2018 Thermosetting Resins 2018 Date: September 25-27, 2018 Location: Berlin - Germany Website: thermosetting-resins.de
October 8th International Conference and Exhibition on Biopolymers and Bioplastics Date: October 15-16, 2018 Location: Las Vegas – USA Website: biopolymers-bioplastics.conferenceseries.com 8th International Conference on Polymer Science and Engineering Date: October 15-16, 2018 Location: Las Vegas – USA Website: polymerscience.conferenceseries.com
November Polymers + 3D Date: November 1-2, 2018 Location: Houston – USA Website: www.poly3d.org 9th International Conference on Biopolymers and Polymer Sciences Date: November 1-2, 2018 Location: Bucharest - Romania Website: biopolymers.materialsconferences.com Regional Conference of the Polymer Processing Society (PPS-Americas) Date: November 5-9, 2018 Location: Boston – Massachusetts - USA Website: www.pps2018boston.com.
December 12th SPSJ International Polymer Conference (IPC 2018) Date: December 4-7, 2018 Location: Hiroshima - Japan Website: main.spsj.or.jp/ipc2018 13th European Bioplastics Conference Date: December 4-5, 2018 Location: Berlin - Germany Website: www.european-bioplastics.org/events/eubp-conference Plastics Regulations – 2018 Date: December 11-12, 2018 Location: Pittsburgh - USA Website: www.ami.international/events/event?Code=C0946
January 22nd Thermoplastic Concentrates Date: January 29-31, 2019 Location: Coral Springs - USA Website: www.ami.international/events/event?Code=C0937
Polímeros, 28(2), 2018
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
PolĂmeros, 28(2), 2018
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ABPol Associates Our welcome ... To the new Sponsor Partner Borealis Brasil S.A. We appreciate your valuable support! José Donato Ambrósio President
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
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Polímeros, 28(2), 2018
ISSN 1678-5169 (Online)
http://dx.doi.org/10.1590/0104-1428.2413
Cellulose nanomaterials: size and surface influence on the thermal and rheological behavior Marcos Mariano1,2*, Nadia El Kissi2 and Alain Dufresne1 University Grenoble Alpes, National Center for Scientific Research – CNRS, Grenoble INP, LGP2, Grenoble, France 2 University Grenoble Alpes, National Center for Scientific Research – CNRS, Grenoble INP, LRP, Grenoble, France
1
*marcos.mariano@lgp2.grenoble-inp.fr
Abstract Cellulose nanocrystals (CNC) and nanofibrils (CNF) were obtained by acid hydrolyzis and mechanical treatment, respectively, of cellulosic fibers from paper. Additionally, surface modification was performed for CNC by neutralization (NaOH) and oxidation (TEMPO). The thermal stability, surface properties and rheological behavior of these nanomaterials were compared. A clear difference in CNC surface was found upon neutralization and oxidation treatments, leading to distinct thermal behaviors. Optical and rheological properties were found to be predominantly by the particles size, being strongly affected by inertial effects. Keywords: cellulose nanocrystals, cellulose nanofibrils, thermal stability, rheology.
1. Introduction The use of natural fibers as raw materials to produce biocomposites is a tendency that fits perfectly the actual society needs. A combination of extensive research and recent technological advances allows the use of nanomaterials in new renewable material development. As the most abundant polymer on earth, cellulose studies goes toward the described scenario. Its fibers are now used to produce new composites with different matrix by different processing techniques[1-4]. The use of its natural structure to produce nanomaterials was developed during the last two decades and shows promising results. Many initial difficulties are now overcame and advances have been made concerning their obtention, surface modification and processing[5]. Derivated from natural fibers, bacteria or algaes, cellulose nanomaterials are versatiles materials that are now applied in nanopapers[6], hydrogels[7], composites[8], electronic[9], biomedicine[10], etc. Being a versatile material, cellulose processing can be adjusted to provide different materials as, for example, cellulose nanofibril (CNF) or cellulose nanocrystal (CNC). The first one can be obtained by mechanical defibrillation or enzymatic treatment of cellulose fibers[11], which leads to individualization of cellulose nanofibers that compose the intrisic structure of the polymer . The second one is obtained by the isolation of crystalline domains present in the structure by hydrolysis of amorphous part using acid or enzymatic steps[12,13]. Concerning its actual applications, CNF is used in hydrogels and high technology papers. During its production by mechanical defibrillation, it is common to use TEMPO (2,2,6,6-tetramethyl-1-piperidinyloxy radical) oxidation as pretreatment. This step facilitates the fiber individualization, being a way to save energy and make greener process. CNC is normally used in nanocomposite production to provide better barrier or mechanical properties[14,15].
Polímeros, 28(2), 93-102, 2018
To ensure the quality and explore the complete potential of this nanoparticle, its surface can be modified to improve the filler-matrix compatibility or increase the thermal stability. In the case of molecule grafting to nanocrystal surface, also a TEMPO modification has became a common process used as first stage to allow chemical modification[16]. Aiming in the production of nanocomposites at high temperatures, the CNC surface also can be modified. The presence of sulfate groups attached to its surface after hydrolysis can decrease its thermal stability. A dessulfation process can be performed by the addition of a diluted NaOH solution into the suspension. In this case, the thermal degradation seems to be postponed[17]. In this study, these different materials were obtained from paper using different methodologies. Cellulose nanocrystals were obtained by acid hydrolyzis and neutralized (CNC-n) or oxidized (CNC-t) using TEMPO reagent. The nanocrystal properties were compared to cellulose nanofibrils(CNF) obtained by mechanical treatment. Aiming to compare the properties of these different materials, obtained from the same source, the nanoparticles were characterized by different techniques.
2. Materials and Methods 2.1 Materials The paper used in this study was purchased from Whatman. NaOH and sulfuric acid were obtained from Sigma-Aldrich.
2.2 Methods In this section, the experimental process used during the preparation of the nanomaterials will be described. An illustrative scheme of the obtained materials is reported in Figure 1.
93/102 93
O O O O O O O O O O O O O O O O
Mariano, M., El Kissi, N., & Dufresne, A. filtration and washing with distilled water was performed. CNC was dialysed against distilled water for several days. Nanofibril preparation. The previrous oxidized fibers were resuspended in water and the concentration was adjusted to 1wt%. This suspension was treated with a Masuko Supermass colloider grinder for 60 passes. 2.2.2 Materials characterization
Figure 1. Illustrative scheme of the preparation methods and obtained materials.
2.2.1 Nanomaterials obtention Nanocrystal preparation. Firstly, the cellulose fibers obtained from paper were maintained under agitation in a 2% NaOH solution for 6 hours at 25 °C and after that it was washed until neutral pH and dried in ambient conditions. For hydrolysis of bleached fibers, 400 mL of a sulfuric acid solution (65 wt%) was heated up to 45 °C and 20 g of fibers was suspended in the solution under a strong mechanical stirring for 45 min. After the desired time, the reaction was stopped with the addition of some ice cubes and the suspension was centrifuged in cycles of 10 minutes under a centrifuge force of 1373.6 G. After each cycle the acid supernatant was discarded and replaced by distilled water, this process was repeated a few times to increase the pH of the suspension. The resultant water suspension was dialyzed with a cellulose membrane against distilled water until its pH was around 5. Some drops of chloroform were added to avoid any bacterial proliferation and the suspension was stored under refrigeration. Particle neutralization. After its preparation and dialysis, the pH of part of the cellulose nanocrystal suspension was adjusted to 9 using a diluted NaOH solution. This suspension was kept under agitation for 24h and again dialysed until neutral pH. Tempo oxidation. The oxidation of hydroxyl groups by TEMPO reagent was develop by Nooy[18] and is broadly used in the oxidation of cellulosic materials. Here, the generic procedure used for the modification of cellulose nanocrystals and cellulose fibers (pretreatment) is described. In an ice bath, a weight of 4 g of cellulose material was suspended in 400 mL of water to obtain a 1 wt% suspension. The suspension was homogenized with a strong mechanical agitation and 0.4g of NaBr and 0.1g of 2,2,6,6-tetramethylpiperidine (TEMPO) were added to the suspension. This agitation was maintained for almost 20 min, to ensure a complete dissolution of the reagents, and slowly 10 mL of a 12% aqueous NaOCl- solution was added to the reaction medium. The pH should be kept between 10 and 11 with an eventual addition of diluted NaOH solution. After 2 hours, 10 mL of ethanol was added to the medium and the pH adjusted to 7 with a diluted acid solution. For the cellulose fibers, 94 94/102
Atomic Force Microscopy (AFM). The images were obtained on a NanoscopeIIIa microscope from Veeco Instruments. A drop of a diluted aqueous CNC suspension with 0.01 wt% concentration was deposited on a mica substrate and dried. It was imaged in tapping mode with a Silicon cantilever. The nanoparticle dimensions were estimated from 50 measurements analyzed using the associated Nanoscope software. Zeta Potential (ξ). Nano particle suspensions with concentration around 0.01 wt% were analyzed on an equipment model DTS0230 from Malvern Instruments. To avoid the effects of ionic strength and pH during measurements, all the concentrated solutions were diluted in an aqueous standard solution with pH 10 and 180 µS/cm conductivity. This solution was prepared by the addition of diluted NaOH solution and solid NaCl into distilled water. Contact Angle (CA). Films 0.1 ± 0.01 mm thick were prepared by casting/evaporation and analyzed with an Attension Theta contact angle meter equipment by using water as liquid. The samples were cast in a Teflon mold to ensure a smooth surface and the contact angle was measured during 120 s. The final angle value before surface deformation was used as reference. Fourrier Transform Infrared Spectroscopy (FTIR). ATR mode measurements were performed on a FTIR Perkin‑Elmer Spectrum One equipment between 600 and 4000 cm-1 in 4 cm-1 intervals. All the analyses were carried out using films casted and dried at room temperature. Thermogravimetric Analysis (TGA). The analysis was carried out under Air atmosphere using Perkin-Elmer TGA‑6 equipment. The sample heating was performed from room temperature up to 600 °C with a heating of 10 °C.min-1. Rheological Measurements. The rheological behavior of the dispersions was characterized with a DHR-3 equipment from TA Instruments. A cone (2º, 50 mm) - plate (50 mm) geometry was used to study the dilute suspensions in flow and oscillatory modes at 20 °C. Degree of polymerization. The samples viscosity were measured by the TAPPI methodology T230m08 and the D.P. estimated by the Equation 1. D.P. =−449.6 + 598.4 ln [ η] + 118.02(ln [ η])² (1)
X-ray diffraction. XRD measurements for the samples were recorded on a Philips PW 1720 X-ray generator operated at 45 kV and 40 mA in a Bragg–Brentano geometry. The 2θ range was from 5º to 65º using a fixed time mode with a step interval of 0.066º and Cu Ka radiation (k = 1.5418 Å). The crystallinity index (C.I.) was obtained by the relation between crystalline phase 2θ = 22º and amorphous phase on 2θ = 18.5º following the Equation 2. Polímeros, 28(2), 93-102, 2018
Cellulose nanomaterials: size and surface influence on the thermal and rheological behavior
= C.I . ( % )
I 200 − Iam ×100 (2) I 200
Samples nomemclature and other obtained properties are on Table 1.
3. Results and Discussion 3.1 Atomic force microscopy and surface properties Different procedures were applied to obtain cellulose nanomaterials with distinct characteristics. The morphology of the nanoparticles resulting from the acid and mechanical treatments is shown in Figure 2. For the acid hydrolyzed material (CNC) rod-like structure can be observed (Figures 2a and 2b). These nanometric rods are the non-hydrolyzed crystalline domains present in the hierarchical structure of cellulose. After post-treatment performed on the pristine CNC, the dimentions (Table 1) of CNC-t and CNC-n seem to be basically the same, no visible sign of degradation caused by the basic of oxidative treatment at the surface could be found. In both cases, the particle dimentions are in agreement with dimentions reported in literature, being similar to the estimated size of nanocrystals obtained from hardwood[19]. Figure 2c shows CNF particles. These particles have very different characteristics when compared to the previous one. Longer and more flexible than CNC nanorods, CNF occurs as particles with higher aspect ratio (L/d). This characteristicis are in agreement with the expected structure of particles obtained after mechanical or enzymatic treatment. In this process, hydrolyzis of the amorphous part of cellulose does not occur, with the isolation of individual fibrils being the
principal objective. From the AFM image it is possible to observe some microscale fibers. Probably, the mechanical treatment was not able to provide particles with a narrow size distribution, were individual nanofibers are completely isolated. However, the presence of fibers with a nanometric diameter (around 6 nm) is predominant compared to bigger agglomerates. Besides its dimentions, the particle surface also are quite different. The ξ measurements shows that the CNC-n present a less negative surface charge. This is probably a consequence of the neutralization stage, where the NaOH solution can neutralize acid residues and the negative sulfate groups presents at the nanoparticle surface. As a consequence, it means that this suspension is less stable against coalescence than the others. CNC-t and CNF have similar ξ values. It can suggest that the number of oxidized –OH groups during the TEMPO reaction was not so different although the different surface areas avalaible during the reaction.
3.2 Infrared spectroscopy Figure 3a shows the FTIR spectrafor the different nanoparticles. The figure also shows the spectrum of pristine cellulose fibers (CP) to be used as reference. Firstly, all the curves show typical peaks of cellulose, i.e. C-H (2900, 1300 cm-1), C-C (1600 cm-1), C-O (3330,1000 cm-1), C-O-C (1150 cm-1) and O-H (3330 cm-1) bands. However, besides the homogeneous distribution of size and no sign of particle degradation, the FTIR analyzis shows some modifications in CNC structure. Probably, these modifications occur at the surface of the nanorods during the post-treatment of oxidation or neutralization. Surprisingly, the spectra for CNC-n and CNC-t are quite
Figure 2. AFM images of (a) CNC-n, (b) CNC-t and (c) CNF.
Table 1. Samples nomenclature and general properties. Contact
Sample
Nomenclature
D.P.
L (nm)
D (nm)
L/D
ξ (mV)
Paper NaOH-neutralized Cellulose Nanocrystals TEMPO-oxidized Cellulose Nanocrystals Cellulose Nanofibrils
CP CNC-n
2598 462
147 ± 20
8.5 ± 2.5
17.3
-26.9
Angle (°) 17
CNC-t
375
145 ± 41
8.3 ± 2.7
17.4
-32.7
33
88
CNF
485
~ 1000
6.0 ± 3.7
>100
-33.7
21
80
Polímeros, 28(2), 93-102, 2018
C.I. 87
95/102 95
Mariano, M., El Kissi, N., & Dufresne, A. similar. The presenceof a C=O peak at 1740 cm-1 for CNC-t is a consequence of the TEMPO-oxidation, that can oxidize the C-O bonds naturally present in the cellulose structure. On CNC-t also is possible to observe a 2840 cm-1 peak, that is normally attributed to C-H bonds in alkanes. This peak can corroborate the higher value of contact angle if we assume some surface degradation during the reaction. It seems possible once the C.I. for this sample had a slight decrease when compared to the CNC-n sample. The presence of the same C=O signal for CNC-n was not expected. In this sample, the neutralization process was supposed to cause just a dessulfation of the nanocrystal surface[17]. However, this step seems to also cause an oxidation of the O-H groups to form carbonyl groups, probably resulting in carboxylic acids or aldehydes formation. CNF also presents the C=O peak, a consequence of the TEMPO pre-treatment. In this sample it is worth to notice a prominent band at 1600 cm-1, that can be related to a slightly offset O-H bending peak of absorbed water (~1620 cm-1) band or a C=C aromatic band. In the last case, it can be attributed to some residual hemicelluloses. Once in this sample no acid hydrolysis was performed, the presence of residual oligomers is possible. Concerning the samples polymorphism, all the samples had shown cellulose I patterns with intense peaks on 2θ equals to 15º (110), 17º (110) and 23º (200). However, CNC-n shows a small peak on 2θ equals to 12º (101), that is characteristic of cellulose II. It suggests a cellulose conversion (of a small number of chains) from type I → II during the neutralization step. Figure 3b shows the XRD curves for the CNC-t sample in comparison to CNC-n.
3.3 Thermal behavior As a consequence of the surface modification suggested by FTIR, C.I. and zeta potential (Table 1) values, the particles could show very different properties, once these parameters are critically important to define particle behavior in suspension and its thermal behavior. The thermal stability of the samples was investigated by thermogravimetric analysis. The obtained thermograms are shown in Figure 4 and the related data reported in Table 2. Firstly, a higher amount of water present in the nanometric samples is observed. This can be explained by the higher surface area of this particles, providing more available –OH groups with which water molecules can interact. As suggested by contact angle measurements (Table 1) and FTIR, CNC-t is more hydrophobic than other nanocellulosic samples, showing a lower water content. The thermal behavior and degradation stages of cellulose materials are very well described in literature. While cellulosic fibers tend to degrade in multiple stages due to their heterogeneous composition with presence of macromolecules such as lignin and hemicelluloses, acid hydrolyzed nanocrystals tend to degrade at lower temperatures due to the presence of sulfate groups[20,21]. Here, the use of a bleached commercial paper seems to minimize the first phenomena for the CP sample. Nevertheless, it is possible to observe the degradation of cellulose over a very broad temperature range (Figure 4b), being the cellulose maximum degradation peak normally described around 360 °C[22]. Concerning the nanomaterials (CNC and CNF) the thermal degradation can be distinguish in two groups. In the first one, CNC-n seems to have a similar degradation mechanism
Figure 3. FTIR (a) spectra for paper and nanocellulosic materials and XRD (b) patterns for CNC-t and CNC-n. Table 2. Thermogravimetric analysis data. Sample CP CNC-n CNC-t CNF
96 96/102
Water % (at 150 °C) 4.5 5.8 5.0 5.8
Degradation Stages
Onset degradation
dTg peak
2 2 2 2
214 235 190 220
350 350 315 315
Char residue % at 600 °C 11.7 20.3 25.2 28.8
Polímeros, 28(2), 93-102, 2018
Cellulose nanomaterials: size and surface influence on the thermal and rheological behavior as CP. These nanoparticles presents a good thermal stability, with a dTG (Figure 4b) in the same temperature range as pristine cellulose (CP), but its degradation occurs within a narrower temperature range. This higher thermal stability is a direct consequence of the neutralization step, that avoids early degradation of the material by removing residual acids and causing dessulfatation of the particle surface[23]. As expected, the molecular weight (expressed in terms of DP, Table 1) of the pristine cellulose was almost six times higher than for other samples. The nanomaterials had present similar values of DP on a range similar to the described by Matsuoka et al.[24] as the length of cellulose crystalline region. However, it seems to not be the major influence on this property since the CP and CNC-n shows similar dTG values besides its different values of DP. In the second group, CNC-t and CNF also show similar thermal behavior. In these samples, the presence of two degradation stages is more pronounced than in the first group. These two stages are normally visible on acid hydrolyzed nanocrystals. In fact, it seems to be a consequence of the early degradation of external chains of the particle that accelerates the formation of oligomers that can cause a kind of caramelisation effect in the sample surface that can coat the material core, retarding its degradation. The similar degradation temperatures and dTG for CNC-t and CNF suggests that carboxylic groups can also produce this kind of effect. For all samples, there is a remarkable increase in char residue when comparing nanomaterials to pristine paper. Is known that the pathway of thermal degradation of the samples seems to be responsible to change the residual content and that parameters such as activation energy, heating rate, temperature of dehydratation and levoglucosan formation can strongly affect the residues content and appearance[25]. Besides the normally cited presence of sulfate groups, other surface characteristics also can lead to modify the degradation pathway. For example, an increase on the surface area or presence of grafted groups on cellulose surface also can cause variation on the thermal degradation.
Here, the higher residues content were present on TEMPO-oxidized samples. It occurs due to the easy surface oxidation of the material and is clearly observable by comparison with the other samples (presenting C=O or -OH groups, for example). This oxidation seems to accelerate the process of coating the internal cellulose chains. Due to its lower C.I. (Table 1), the core chains of CNF and CNC-t are not so protected from external heat transfer as CNC-n and these samples presents lower dTG values. However, on these two samples the final mechanism seems to occur by a pathway of smaller Ea that leads to the formation of char[24].
3.4 Optical properties and suspension behavior In water suspension, these particles also show different behaviors. The observation of nanoparticles under a polarized light is a way to observe oriented and cristalline structures. Figure 5 shows the presence of birefringence in the nanocrystal samples. The birefringence of CNC suspensions can result from two factors: (i) the structural form anisotropyof cellulose nano-domaines (anisotropic refractive index, ∆n of 0.05) and (ii) a flow anisotropy resulting from alignment of the nanorods (if long enough) under flow. In this study, this alignment was induced by the creation of a shear in the suspension using magnetic stirring. Some studies demonstrated that an imposed shear can produce planar domains of randomly oriented nanoparticles that are aligned or broken with the shear rate variation[26]. The level of organization in the nanoparticle suspension is a key factor to its rheological properties and will be discussed in sequence. This birefringence phenomenon cannot be observed in the CNF system. In addition to its big dimentions (at least in one axis), it was demonstred by Karppinen et al.[26] that CNF have a tendency to entangle (in a reversible way) when submitted to shear, making difficult their alignment with the flow.
3.5 Flow curves According to Bercea and Navard[27], the critical concentrations for a regime transition between diluted‑semidiluted (φ* ) and semidiluted-concentrated (φ** ) CNC particles suspended in water can be calculated using the rigid rod approximation:
Figure 4. (a) TGA and (b) dTG curves for: CP (▲), CNC-n (X), CNC-t (o) and CNF (▄). Polímeros, 28(2), 93-102, 2018
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Mariano, M., El Kissi, N., & Dufresne, A.
Figure 5. Birrefringence domaines of nanoparticle suspensions with concentration of 0.6wt%.
d2 φ* = 2 (3) L
d φ** = (4) L
Equations 3 and 4 can show the transitions limits calculated from the particle dimentions, where L is the length and d the diameter of the particle. In our system, the particle dimentions lead to a diluted to semi-diluted regime transition around 0.0033 vol% (or 0.05 wt%) and a semi-diluted to concentrated transition around 0.057 vol% (or 0.91 wt%) . In this study, we choose to perform all the rheological measurements in the semidiluted regime where the samples are in an isotropic-at-rest regime and no anisotropy can be found. For CNF suspensions, the concentration was chosen to be the same as for the CNC suspensions in order to facilitate the discussion and data comparison. Figure 6 presents the evolution of the viscosity as a function of the shear rate. For all samples it is possible to observe a decrease in the viscosity value with shear rate increase (shear thinning behavior). For the CNC samples, the viscosity value at lower shear rates is around 1 Pa.s, similar to values reported in the literature[27,28]. For CNC suspensions, the shear thinning behavior is normally explained by the progressive organization of the particles in suspension. At lower shear rates the nanocrystals are assumed to randomly organize. However the scenario starts to change with the shear rate increase and the rods are progressively align in the flow direction. This organization causes a decrease in the drag force, decreasing the viscosity value until a plateau. In this region, it is possible to observe an almost Newtonian flow behavior due to the maximum organization of the nanocrystals. Besides the similar size and shape, CNC-n and CNC-t viscosity curves shows some slight differences. The viscosity for the CNC-n suspension is higher at lower shear rates and this sample reaches the Newtonian plateau before CNC-t. 98 98/102
Figure 6. Steady-state viscosity versus shear rate for CNF (▄), CNC-n (X) and CNC-t (●) suspensions with concentration of 0.6 wt%.
Probably this behavior is caused by the differences in the nanocrystal surface chemistry. For the CNF suspension, the viscosity values are significantly higher than for CNC at lower shear rate values, being around 30 Pa.s. This value is also similar to previous data reported in literature[11,26]. The shear thinning behavior is even more accentuated in this system. As a diluted sample, it is not possible to observe an intermediary plateau that is described by some authors with the concentration increase. Concerning the flow kinetics and the microstructureof this system under shear, it seems to be much more complicated than for the CNF system. Inhomogeneous flow is caused by erratic mesostructural changes. Floculation, fiber entanglement and wall slippage are phenomena present in this system due to the particles characteristics, i.e. size polidispersity, flexibility and surface charge[26,29]. Polímeros, 28(2), 93-102, 2018
Cellulose nanomaterials: size and surface influence on the thermal and rheological behavior 3.6 Oscillation measurements The rheological behavior of the particle dispersions was also characterized by dynamic oscillatory measurements. The intrinsic difference between the nanorods and the nanofibers was very clear even during the oscillation strain sweeps were the CNC samples showed a critical strain around 0.5%, a much higher value than for CNF, which showed a critical strain around 0.025%. The evolution of the values of G’ and G” as afunction of angular frequency is ploted in Figure 7. Cellulose nanomaterials are known for the possibility for form gels. These structures can be applied in different areas due to their great capacity to retain water and, still, behave as a solid. The “gel” state is defined by rheology as a suspension state where it presents a storage modulus (G’) much higher than its loss modulus (G”), i.e. G’ >> G”. The storage modulus values can be regarded as the suspension ability to restore energy, and therefore, as the gel strength of the suspensions. In this study, all the samples presents gel-like behavior with G’ values higher than G”. However, the ratio between G” and G’ (called tan δ) is greater than 0.1 for all studied samples until intermediary angular frequencies (Fq). This means that the samples are not true gels, but can be considered as weak gels[30], what is reasonable considering their low concentrations. This denomination is
corroborated by the dependence of G’ with Fq, since very stiff CNC gels (arising from more concentrated CNCs suspensions) do not present this dependence, presenting an almost constant G’ value[31,32]. For the nanorod samples, the weak gel behavior is very clear until Fq values closer to 1 rad.s-1. After this Fq the values of loss and storage moduli start to be very similar (i.e. tan δ is close to 1) and increase. This behavior is also described by Karppinen[33] and Wu et al.[34] for diluted nanocellulose (fibers and rods) suspensions. The first one considers it as a consequence of the concentration proximity with the threshold concentration for gel-like behavior[33,34]. However, this phenomenon seems to be more complex. In this study, the G’ and G” valuesfor diluted suspensions start to increase and became similar at intermediary Fq. This behavior is described by Ewoldt et al.[35] as a consequence of inertial forces that can cause interferences in the measurements. It seems to be explained by the Reynolds number, Equation 5. Re =
ρ.ν.e InertialForces ~ (5) η ViscousForces
Where: ρ is the fluid viscosity, ν is the fluid medium speed and e is the characteristic length of the flow process being considered (i.e. the rheometer plates gap in our case) and η the fluid viscosity.
Figure 7. Storage, G’ (solid symbols) and loss, G” (empty symbols) moduli as a function of oscilation frequency for (a) CNC-t, (b) CNC-n and (c) CNF. Polímeros, 28(2), 93-102, 2018
99/102 99
Mariano, M., El Kissi, N., & Dufresne, A. This equation illustrates the relationship between inertial and viscous forces. When the fluid presents low viscosity or flow values, the inertial effect can be negligible due to a compensation between both forces. However, in the case of systems with very low viscosity and/or high flow values, these effects begin to be representative and can influence the measurements, leading to false data. It is possible to observe in Figures 7a and 7b an increase in G’ and G” at higher frequency values. Probably it occurs because inertial effects are causing a stress that is measured by the equipment. After a critical frequency it begins to distort and the obtained modulus values and G’ and G” seems to grow. This critical frequence (Fqc) is a limit after which it is not possible to neglect inertial distortions in the obtained data. Equations 6 and 7 were suggested by Baravian (2013) and can provide quantitative information about this frequency due to inercial effects of fluid and equipment respectively[36]. Fqc <
η (6) ρ.e²
Fqc <
G′ / a (7) 2π
Where: (Equation 6) Fqc is the critical frequence (in Hz), η is the viscosity values at 1 Hz frequency, ρ is the suspension density, e is the plates gap distance; (Equation 7) α is a 0.05 coeficient for standard geometries and G’ is the modulus values at low frequencies.The respective obtained values for the studied suspensions are collected in Table 3. For CNC samples, Fqc is quite similar, considering the fluid and equipment inertial effects. It means that most of the frequency range for the CNC suspension at this concentration is under influence of inertial effects. The inertial effect from the rheometer seems to arise at lower frequencies if compared to the fluid effects for all the samples. According to Baravian, above Fqc the rheometer error can be higher than 20%[36]. Since thes effects are viscosity-dependent, they are much less observable for CNF in the studied Fq range. This sample present higher viscosity values, the fluid inertial effects only become significant after 92.4 rad.s-1 and probably shows most visibly distorted data at higher frequencies. For lower Fq, where no inertial effects are affecting the data, the magnitude of the modulus values are clearly different between CNC an CNF samples, being much superior for the nanofibril suspension. It suggests that nanofibrils could create a stronger fibrous network at this concentration. In this suspension G’ value is higher than G” over the whole studied frequency range and a relative weak influence of the frequency was found, a typical behavior of weak gels[32]. Table 3. Samples viscosity and inertial critical frequencies. Sample MNC-t MNC-n CNF
100 100/102
Fluid Inertia Fqc
Viscosity (mPa.s) 2.1 1.7 53.0
(rad.s-1) 3.6 2.96 92.4
Equipment Inertia Fqc G’ (mPa) (rad.s-1) 47 35 3800
0.96 0.83 8.71
4. Conclusions The aim of this study was to show how different properties can be obtained from the same cellulose source due to the versatility of this material. By controlling the size/shape of the particle and its surface charge and composition it was possible to modify its thermal and rheological characteristics. The presence of charged groups for CNCs seems to be a key parameter that induces their earlier degradation during heating, an important drawback for composite processing by extrusion or injection. The particle dimentions seem to be responsible for significant modifications in the rheological parameters, such as viscosity and modulus of the nanoparticle suspensions in water. Also, we propose that diluted nanocellulose suspensions measurements can be highly influenced by inertial effects after a critical frequency, leading to false data obtainment. These effects also seem to be dependent of the sample dimensions, being more pronounced for CNC in the studied concentration and frequency range.
5. Acknowledgements The authors gratefully acknowledge the Brazilian National Council for Scientific and Technological Development (CNPq) and “Ciência Sem Fronteiras” (CsF) program for the financial support (PhD fellowship of M.M.). LGP2 and LRP are part of the LabEx Tec 21 (Investissements d’Avenir - grant agreement n° ANR-11-LABX-0030) and of the PolyNat Carnot Institut (Investissementsd’Avenir - grant agreement n° ANR-16-CARN-0025-01). This research was made possible thanks to the facilities of the TekLiCell platform funded by the Région Rhône-Alpes (ERDF: European regional development fund).
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Mariano, M., El Kissi, N., & Dufresne, A. suspension: Influence of concentration and aspect ratio. Journal of Applied Polymer Science, 131(15), 1-8. http:// dx.doi.org/10.1002/app.40525. 35. Ewoldt, R. H., Johnston, M. T., & Caretta, L. M. (2015). Experimental challenges of shear rheology : how to avoid bad data. In S. Spagnolie (Ed.), Complex fluids in biological systems (1st ed., pp. 207-241). New York: Springer-Verlag. http://dx.doi.org/10.1007/978-1-4939-2065-5.
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36. Baravian, C. (2013). Effets inertiels en rhéométrie instationnaie. In J. L. Grossiord, & A. Ponton (Eds.), La mesure en rhéologie: des avancées récentes aux perspectives (pp. 31-48). France: EDP Sciences. Received: Nov. 09, 2015 Revised: Apr. 01, 2016 Accepted: Apr. 13, 2016
Polímeros, 28(2), 93-102, 2018
ISSN 1678-5169 (Online)
http://dx.doi.org/10.1590/0104-1428.2357
Effect of compatibiliser on the properties of polypropylene/ glass fibre/nanoclay composites Normasmira Abd Rahman1*, Aziz Hassan1 and Javad Heidarian1 Polymer and Composite Materials Research Laboratory, Department of Chemistry, University of Malaya, Kuala Lumpur, Malaysia
1
*nmmira@um.edu.my; nmmira14@gmail.com
Abstract Glass fibre (GF), nanoclay (NC) and hybrid composites compatibilised with maleic anhydride polypropylene (MAPP) were prepared by extrusion and injection moulding. The fourier transform infra-red spectra revealed the characteristic absorption peaks of MAPP in the compatibilised GF and NC composites. A decrement in the peak intensity of X-ray diffraction patterns of NC composite was obtained as the MAPP content increased indicating a partial exfoliation of NC. The thermogravimetric analysis showed that the incorporation of MAPP into hybrid composites reduced the thermal stability of the material. The dynamic mechanical analysis showed an increase in the storage modulus in the hybrid composites with lower content of MAPP due to the enhancement in the interfacial adhesion between the GF, NC and PP matrix. Keywords: compatibiliser, dynamic mechanical properties, hybrid materials, thermal behavior.
1. Introduction Polypropylene (PP), one of the most exploited thermoplastic resins in the composites, alloy and blends industries[1], possesses outstanding properties like low density, good abrasion resistance and excellent electrical properties[2]. However, to cope with some limitations of PP, for example, low stiffness and low strength, and to expand their applications in different sectors, inorganic fillers, such as glass fibre (GF), carbon nanotubes and clays, are often added when processing the polymer composites, which normally combine the advantages of their constituent phases[3]. Of these composites, glass fibre reinforced PP composite is quite attractive because of their ease of fabrication, superior mechanical properties, high specific modulus and strength and low manufacturing cost[4]. Meanwhile, a relatively new development in polymer-clay nanocomposites (PCN) has attracted great interest as they exhibit remarkable improvement in the material properties. Researchers first demonstrated that the properties of nylon-6, containing a homogenous dispersion of 2-7 wt% of nanoclay (NC) showed excellent mechanical, thermal and barrier properties when compared with non-reinforced nylon[5,6]. A number of other researchers also began an investigations with various types of polymer nanocomposites[7-9]. Recently, it has been observed that by hybridising nanoparticles into the matrix of fibre-reinforced composites, synergistic effects may be achieved[10-12]. Mohan and Kanny[13] in their work, combined PP/NC composite with chopped GF. A small amount (5 wt%) of nanoscale dispersed layered silicate was shown to enhance the degree of crystallinity as well as the tensile properties. Using nano reinforcements in a GF reinforced PP matrix, Cui et al.[14] observed that the addition of 2 wt% nano-ZnO in 40 wt% GF composite gave
Polímeros, 28(2), 103-111, 2018
the optimum mechanical properties. However, the extent to which fillers can vary the properties of the resultant composite parts is depends on the nature of adhesion between the filler and the matrix. Although a number of reports have been published on the use of maleic anhydride grafted polypropylene (MAPP) as compatibiliser for PP composites, there is very little information available regarding the effect of compatibiliser in the hybrid composite systems[15]. For this reason, in this paper, the function/role of MAPP as coupling agent in PP/GF/NC hybrid composites prepared through a melt processing method is investigated. Chemical analyses and microstructure characteristics were used to evaluate the composites.
2. Experimental 2.1 Materials Chopped E-glass fibre (GF) with the density of 2,550 kg m-3, average diameter of 14 µm, and length of 6 mm was used as a principal reinforcement. It was obtained from KCC Corporation, Korea. The nanoclay (NC, type PGV), was a natural, untreated montmorillonite clay, with a particle size of about 16 µm and a density of 776 kg m–3 and manufactured by Nanocor, USA. Commercially available polypropylene (PP, Propelinas H022), with a density of 910 kg m-3 and melt flow index of 11 g/10 min (at 230 °C, 2.16 kg load), was supplied by Petronas, Malaysia and used as the matrix. Maleic anhydride grafted polypropylene (MAPP, Polybond 3200) with density of 910 kg m-3 was supplied by Chemtura Corporation, USA (formerly Crompton Corporation) and used as the compatibiliser.
103/111 103
O O O O O O O O O O O O O O O O
Rahman, N. A., Hassan, A., & Heidarian, J. composites with a nominal thickness of 200 nm were analysed with a Hitachi H-600 (Japan) TEM. Bright-field TEM images of the composites were obtained at 300 kV under low-dose conditions.
2.2 Specimen preparation In order to produce composites materials, PP, nanoclay, glass fibre and MAPP were pre-mixed in different compositions, as presented in Table 1. The materials were then compounded using the Brabender, KETSE 20/40 (Germany) twin-screw extruder with the screw diameter and screw aspect ratio of 20 mm and 40, respectively. The temperature profile was set between 185 °C and 190 °C. For the PP/GF composites, the screw speed was set to 100 rpm, whereas for the PP/NC composites, the screw speed used was 800 rpm. In order to produce PP/GF/NC composites, the different ratios of the PP/NC pellets and GF were physically mixed and re-compounded in a twin-screw extruder, using the same temperature profile and screw speed of 100 rpm. The materials extruded from both formulations were pelletized into length of about 6 mm. The pellets were injection moulded into a dumb-bell shaped tensile test specimens with geometry defined in the ASTM Standard D-638, type 1[16] using a Boy 55M (Germany), a 55 tonne clamping force injection moulding machine. The processing temperature was set between 175 °C and 185 °C and the mould temperature was set at 25 °C. The screw speed was maintained at 30-50 rpm. The list and abbreviation of specimens prepared are given in Table 1. Specimens were designated according to their composition; for example, (PP80:C5/G15)/NC6 referred to specimen with 80 wt% of PP, 5 wt% of MAPP, 15 wt% of GF and 6 phr of NC.
2.4 Thermal and dynamic mechanical analysis Thermogravimetric analysis (TGA) measurements were carried out using the Perkin Elmer TGA (USA) 6 on 5-10 mg samples of each of the composites in a ceramic crucible, over a temperature range from 50 to 850 °C at a heating rate of 10 °C min-1. The tests were conducted in a nitrogen atmosphere at a flow rate of 20 mL min-1. Test specimens for dynamic mechanical analysis were taken from the middle section of the injection moulded dumb-bell test bar with dimensions of 60.0 mm × 13.0 mm × 3.3 mm (length × width × thickness) characterised using dynamic mechanical analyser, Thermal Analysis Instrument, TAI Q800 (USA). Measurements were conducted over a temperature range from –100 °C to 110 °C with a heating rate of 3 °C min-1 at a constant frequency and amplitude of 1.0 Hz and 15 µm, respectively. Specimens were subjected to a three-point bending mode with a support span of 50 mm.
3. Results and discussion 3.1 Fourier-transform infra-red (FTIR) The FTIR spectra of PP matrix, MAPP, NC and composites in the region of 650 cm-1 to 4000 cm-1 are given in Figure 1. The bands at 1375 cm-1 and 1451 cm-1 are the characteristic of polypropylene. In the case of MAPP, absorption bands at 1700 cm-1 and 1750 cm-1 are observed, which are assigned to the absorption of carbonyl groups (C=O) of maleic anhydride (MA)[17]. Therefore, it is confirmed that MA was grafted onto the PP backbone. Meanwhile, for NC, the band at 3588 cm-1 is attributed to the hydroxyl stretching of Al-OH and Si-OH[18]. The broad band in the region of 700 cm-1 to 1100 cm-1 is mainly due to the contribution of several structural –OH groups in the clay[19]. Moreover, the absorption in the regions of 1624 cm-1 is assigned to –OH
2.3 Chemical and microstructural characterisations FTIR spectra of samples were recorded using the FTIR spectrophotometer (Spotlight 400, Perkin Elmer, USA) at a resolution of 4 cm-1 for 64 scans in the range of 650-4000 cm-1. Powdered clay and clay nanocomposites were analysed by XRD using Philips-binary diffractometer and scanned over the interval of 2Θ = 2°-30° at 40 kV and 30 mA, with CuKα radiation. The fractured surface of the various nanocomposites with a thin layer of gold coated on the surface (thickness of 0.014 μm) was examined using a SEM model Auriga Zeiss (Germany). The microstructure of the PP/NC6 and PP/GF15/NC6 Table 1. Designation and compositions of composite specimens. Sample NC powder PP (PP100:C0)/NC6 (PP98:C2)/NC6 (PP95:C5)/NC6 (PP92:C8)/NC6 (PP85:C0/G15) (PP83:C2/G15) (PP80:C5/G15) (PP77:C8/G15) (PP85:C0/G15)/NC6 (PP83:C2/G15)/NC6 (PP80:C5/G15)/NC6 (PP77:C8/G15)/NC6
104 104/111
Matrix weight fraction,
MAPP weight fraction,
Fibre weight fraction,
Clay content
WM (%)
WMAPP (%)
WF (%)
(phr)
100 100 98 95 92 85 83 80 77 85 83 80 77
0 0 2 5 8 0 2 5 8 0 2 5 8
0 15 15 15 15 15 15 15 15
0 6 6 6 6 6 6 6 6
Polímeros, 28(2), 103-111, 2018
Effect of compatibiliser on the properties of polypropylene/glass fibre/nanoclay composites bending mode in adsorbed water. Meanwhile, the band at 1092 cm-1 is attributed to Si − O stretching (in-plane) vibration for layered NC. The bands at 902 cm-1, 864 cm-1 and 826 cm-1 are attributed to Al2OH, AlFeOH and AlMgOH bending vibrations, respectively[19]. The characteristic absorption bands of carbonyl groups of MA can be seen at 1722 cm-1 for compatibilised PP/GF15 composite and at 1762 cm-1 and 1693 cm-1 for compatibilised PP/NC6 composite. Again, as mentioned before, a broad band in the region of 700 cm-1 to 1100 cm-1 observed for PP/NC6 composites are related to the characteristic absorption of NC.
3.2 X-ray diffraction properties Several researchers[20,21] have shown that X-ray diffraction method can be used to observe the distribution and dispersion of NC particles in a polymer and to characterise the degree of dispersion. In order to measure the interlayer distance using the diffraction peak and its position in an XRD pattern, the Bragg’s equation is typically employed[7]. Details of this equation have been explained by Rahman et al.[11]. Results from the XRD analyses of composite specimens are shown in Figure 2 and the data extracted from these patterns are tabulated in Table 2. Figure 2 shows the series of XRD spectra of PP/NC6 composites, in which the concentration of MAPP varied from 0 to 8 wt%. The NC concentration in the composite is constant (6 phr) and the interlayer d-spacing of clay powder is
calculated to be 1.00 nm (Table 2). For the uncompatibilised system, the XRD patterns exhibited a significant increase in interlayer d-spacing (1.25 nm after compounding). This indicates that with higher processing screw speed, the PP is able to intercalate into the NC interlayer, even with poor compatibility between PP and NC[12]. Furthermore, for compatibilised systems, XRD peaks are continually shifted to lower angles, indicating an increase in interlayer d-spacing by the diffusion of PP chains. The interlayer d-spacing increased from 1.29 nm for 2 wt% MAPP to 1.34 nm for 5 wt% MAPP. Further addition of 8 wt% of MAPP resulted in insignificant changes in this value. Nevertheless, it should be noted that even though a similar interlayer d-spacing value is obtained at higher MAPP content (8 wt%), there is a decrement in the peak intensity. The peak intensity for PP/NC6 composite with 5 wt% MAPP recorded at 370 counts s-1, reduced to 357 counts s-1 for 8 wt% MAPP. Lertwimolnun and Vergnes[21] suggested that the decrease in intensity and the broadening of peaks indicate that the stacks of layered silicates become more disordered, while maintaining a periodic distance. In addition, the decrease in intensity could be the result of a partial exfoliation of layered silicates. A higher nucleation effect with the present of MAPP in the system is suggested to be due to the higher interaction between filler and matrix. The phenomenon can be seen in the morphological analysis (Figure 3). The addition of MAPP provided better dispersion of the NC within the matrix thus reducing the particle thickness hence increasing the aspect ratio.
3.3 Thermogravimetric analysis (TGA) The thermal stability of composites was studied by the thermogravimetric analysis (TGA). Figures 4 and 5 show the TGA scans in the form of weight change and derivative weight change (DTG) versus temperature for PP/NC6 and PP/GF15 composites. Table 3 presents the quantitative values of the onset temperature, derivative peak temperature, and the temperatures at 5%, 10% and 50% of weight loss, which are referred to as: TONSET, DTP, T5%, T10% and T50%, respectively.
Figure 1. FTIR spectra of PP matrix, MAPP, nanoclay and composites.
Figure 4 illustrates the TGA and DTG curves of the PP/NC6 composites with and without compatibiliser. With the addition of 2-8 wt% MAPP in the PP/NC6 composites system, an increase of about 14.5 °C to 15.1 °C in the TONSET was observed when compared with the uncompatibilised PP/NC6 composites. This improvement is probably due to the physico-chemical adsorption of the volatile products on the clay[22,23] which indicate that the dispersion of clay is improved by the addition of the compatibiliser in the PP/NC6 composites. Table 2. X-ray diffraction data of nanoclay and clay nanocomposites.
Figure 2. The XRD patterns of nanoclay and PP/6 phr nanoclay composites compatibilised with different MAPP loadings. Polímeros, 28(2), 103-111, 2018
Sample
2θ (°)
NC powder (PP100:C0)/NC6 (PP98:C2)/NC6 (PP95:C5)/NC6 (PP92:C8)/NC6
8.86 7.04 6.85 6.60 6.61
Interlayer d-spacing (nm) 1.00 1.25 1.29 1.34 1.34
Peak intensity (counts s-1) 127 320 467 370 357
105/111 105
Rahman, N. A., Hassan, A., & Heidarian, J.
Figure 3. TEM images of PP/6 phr nanoclay with (a) 0; (b) 2; (c) 5 and (d) 8 wt% of MAPP contents.
Figure 4. TGA and DTG thermograms of PP/6 phr of nanoclay composites with 0 to 8 wt% of MAPP.
Table 3. TGA data of composites. Sample
Range of decomposition (°C)
TONSET (°C)
T5% (°C)
T10% (°C)
T50% (°C)
DTP (°C)
PP (PP100:C0)/NC6 (PP98:C2)/NC6 (PP95:C5)/NC6 (PP92:C8)/NC6 (PP85:C0/G15) (PP83:C2/G15) (PP80:C5/G15) (PP77:C8/G15)
228.3- 465.7 215.2-468.6 240.0-441.4 239.6-445.0 241.3-444.3 250.1-525.4 249.1-472.9 202.3-450.7 237.9-447.2
356.9 360.4 374.9 374.4 375.5 423.3 385.0 373.4 357.5
306.7 309.4 324.3 327.2 322.5 368.0 345.8 331.9 320.1
327.0 334.8 347.2 351.3 345.5 396.4 368.8 351.5 341.8
397.6 406.8 402.9 408.6 405.9 471.7 422.9 408.1 339.2
430.5 429.0 416.0 422.6 419.5 480.7 428.6 420.4 409.2
106 106/111
Polímeros, 28(2), 103-111, 2018
Effect of compatibiliser on the properties of polypropylene/glass fibre/nanoclay composites It can also be seen that the initial thermal decomposition temperatures are enhanced by the addition of MAPP into the PP/NC6 composites. The T5% increased from 309.4 °C for uncompatibilised PP/NC6 composite to 324.3 °C, 327.2 °C and 322.5 °C for PP/NC6 composites containing 2, 5 and 8 wt% of MAPP, respectively. The same trend was observed at 10% weight change (Table 3). It is possible that at this initial degradation event, incorporation of MAPP improved the compatibility and homogeneity between the matrix and the NC, resulting in more thermally stable nanocomposites. By contrast, the T50% and DTP values decreased slightly with the incorporation of MAPP into the PP/NC6 system. Figure 5 shows the TGA and DTG thermograms of PP/GF15 composite as a function of MAPP contents. It is
observed that the thermal stability of the composites generally reduced with increasing MAPP contents. Better compatibility between PP and GF, expected by the incorporation of MAPP into the composites, may not be the criteria for improvement in the thermal stability.
3.4 Dynamic mechanical analysis (DMA) The dynamic storage modulus (Eʹ) is analogous to the flexural modulus measured as per ASTM D-790 standard[2,24,25] and closely related to the load bearing capacity of a material. The storage modulus values at –100 °C and 25 °C are referred ′ °C , respectively. Variations of Eʹ as to as E−′ 100°C and E25 a function of temperature for composites are graphically illustrated in Figure 6a-c. A sharp rate of decrease from –25 °C to about 25 °C in Eʹ is suggested to be associated
Figure 5. TGA and DTG thermograms of PP/15 wt% glass fibre composites with 0 to 8 wt% of MAPP.
Figure 6. The DMA curves of PP/NC6, PP/GF15 and PP/GF15/NC6 composites with different MAPP contents. Polímeros, 28(2), 103-111, 2018
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Rahman, N. A., Hassan, A., & Heidarian, J. with the relaxation of the amorphous phase. The glassy state of the amorphous phase in the PP goes through its glass transition, followed by a sharp drop in the modulus. Figure 6a shows the E’ curves obtained for PP/NC6 composites with variation in MAPP loadings, from 2 to 8 wt%. Apparently, the compatibilised PP/NC6 composites displayed an improved E’ values throughout the experimental temperature range, indicating that more reinforcing effect is being induced by the compatibiliser[26]. This scenario is considered to be due to the real reinforcement effect of clays[27]. It was also suggested by Modesti et al.[28] and Lai et al.[29] that obvious increment in storage modulus values probably resulted from the better dispersion of the clay in the presence of the compatibiliser. The addition of 2 wt% of MAPP increased the E−′ 100°C of PP/NC6 composite by about 31% from 4.10 GPa to 5.38 GPa. However, further additions of 5 wt% and 8 wt% of MAPP resulted in a lower magnitude of increment to 5.14 GPa and 5.06 GPa, ′ °C values show a similar trend as respectively. The E25 observed for E−′ 100°C . However, the enhancements of the storage modulus values at this higher temperature are less marked compared with at lower temperature.
to 6.90, 6.88 and 6.86 GPa, respectively (Figure 6b). An improvement in E’ could be due to the enhancement in the interfacial adhesion between the GF and the PP by the presence of compatibiliser in the composite system[2]. Through microscopic studies, it can clearly be seen that in the compatibilised composite, some of the polymer matrix adhered to the GF surface indicating a good fibre–matrix adhesion (Figure 7a). Meanwhile, a clean GF surface is observed for uncompatibilised system (Figure 7b) probably due to the weak fibre–matrix interfacial bonding. Figure 6c shows the storage modulus (E’) curves obtained for PP/GF15/NC6 hybrid composites with variation in MAPP loadings, from 2 to 8 wt%. Generally, the compatibilised composites displayed improved E’ values throughout the experimental temperature range, following the same trend as previously observed for PP/NC6 and PP/GF15 composites. The addition of 2 wt% of MAPP into the hybrid composite increased the E−′ 100°C by about 12% from 6.65 GPa to 7.42 GPa. However, further additions of 5 and 8 wt% of MAPP resulted in a lower magnitude of increment to ′ °C values 7.21 GPa and 6.98 GPa, respectively. The E25 show a similar trend, as observed for E−′ 100°C .
The incorporation of 15 wt% of GF into PP matrix greatly increased the value of E−′ 100°C from 4.06 GPa to 6.36 GPa (Table 4). The additions of 2, 5 and 8 wt% of MAPP into the composite further increased the Eʹ value
The loss modulus (Eʺ) is most sensitive to the molecular motions[30]. In this study, Ta E " is referred to as the temperature at the maximum value of loss modulus in ′′ the α–transition region, while the E′′max and E25 °C are the
Table 4. DMA thermomechanical data of PP/NC, PP/GF and PP/GF/NC hybrid composites. Storage modulus (E’) Sample
Tα (°C)
PP (PP100:C0)/NC6 (PP98:C2)/NC6 (PP95:C5)/NC6 (PP92:C8)/NC6 (PP85:C0/G15)/NC6 (PP83:C2/G15)/NC6 (PP80:C5/G15)/NC6 (PP77:C8/G15)/NC6
-16.3 -16.1 -12.2 -11.0 -11.2 -17.1 -10.9 -10.9 -10.6
Loss modulus (E”)
E’–100°C
E’25°C
TaE”
E”max
E”25°C
(GPa) 4.06 4.10 5.38 5.15 5.06 6.65 7.42 7.21 6.98
(GPa) 1.28 1.33 1.91 1.80 1.77 2.94 3.29 2.91 3.17
(°C) -2.3 -3.5 1.7 1.7 1.7 -3.3 1.1 2.0 2.3
(MPa) 140.2 126.9 169.0 165.2 167.9 185.2 210.0 225.5 199.7
(MPa) 74.9 66.9 94.4 91.7 92.7 125.3 132.0 138.3 129.0
TG
Tan delta tan δMAX
tan δ25°C
(°C)
(x10-2)
(x10-2)
2.9 1.7 6.6 6.5 6.7 1.1 5.6 6.6 6.3
6.7 5.9 5.9 6.0 6.3 4.5 4.7 5.5 4.7
5.9 5.0 5.0 5.1 5.2 4.3 4.0 4.7 4.0
Figure 7. SEM images of tensile fracture surfaces of PP/15 wt% glass fibre composites with (a) 5 and (b) 0 wt% of MAPP contents. 108 108/111
Polímeros, 28(2), 103-111, 2018
Effect of compatibiliser on the properties of polypropylene/glass fibre/nanoclay composites magnitude of loss modulus at Ta E " and at 25 °C, respectively. The variation of Eʺ as a function of temperature for PP/NC6, PP/GF15 and PP/GF15/NC6 composites are graphically illustrated in Figure 6d-f. Figure 6d shows the Eʺ curves obtained for PP/NC6 composites with 0 to 8 wt% MAPP loadings. It is evident from this figure that there is a significant increase in the Ta E " values from –3.5 °C for PP/NC6 to 1.7 °C with the addition of 2 wt% MAPP. No significant changes are observed with further additions of 5 and 8 wt% MAPP. The increment in Ta E " may be attributed to the presence of MAPP which improved the interfacial adhesion between PP and NC, thus resulting in a reduction in the mobility of the polymer chains in the amorphous phase of the PP matrix. Figure 6e shows the effect of MAPP content on the Eʺ of PP/GF15 composites. The Ta E " values of the compatibilised composites shifted to higher temperature as compared to the uncompatibilised composite and PP matrix, which is probably due to restricted of the segmental motion in the amorphous PP chains at the fibre–matrix interface[31]. This suggests that the more restricted motion of the polymer molecules due to the increased fibre–matrix adhesion in the presence of MAPP resulted in less distinct and broader transition peaks[2]. The E′′max value decreased generally with MAPP loading. Meanwhile, the values ′′ °C for compatibilised composites showed only a of E25 slight decrement as the MAPP content increased when compared with uncompatibilised ones. The same trend was observed by Nayak and Mohanty[31]. It can be suggested that transition at higher temperature do not change irrespective of the interface modification and hybridisation. Figure 6f shows the Eʺ curves obtained for PP/G15/NCUT6 hybrid composites with variation in MAPP loading, from 2 to 8 wt%. Generally, the compatibilised hybrid composites displayed improved Eʺ values throughout the experimental temperature range. The addition of 2 wt% of MAPP into PP/GF15/NC6 hybrid composite shifted the Ta E " of hybrid composites from –3.3 °C to 1.1 °C. No significant changes are observed with further additions of 5 and 8 wt% MAPP. The increment in Ta E " may be attributed to the presence of MAPP which improved the interfacial adhesion between PP and nanoclay. This observation is also similar with the trend observed for PP/NC6 and PP/GF15 composites. The maximum E′′max value is obtained with the addition of 5 wt% MAPP. The ratio of the loss modulus to the storage modulus is measured as the damping factor or mechanical loss (tan δ)[2,12]. Since the damping peak occurred in the region of the glass transition where the material changed from a rigid to a more elastic state, it is associated with the movement of small groups and chains of molecules within the polymer structure[30]. In a composite system, damping is affected by the incorporation of fibres. This is due mainly to stress concentration at the fibre ends in association with the additional viscoelastic energy dissipation in the matrix material. Chen et al.[32] reported that the tan δ curve of pure PP is generally related to three relaxations localised in the neighbourhood of –50 °C, 10 °C and 100 °C. The variations of tan δ as a function of temperature are represented in Figure 6g-i. In this work, TG is referred to as the temperature at the maximum value of tan δ, while the tan δMAX and tan δ25°C are the magnitude of tan δ at TG and at 25 °C, respectively. Figure 6g shows the effects of the compatibiliser loadings on the tan δ curves for PP/NC6 composites. Polímeros, 28(2), 103-111, 2018
A remarkable shift in TG to a higher temperature is recorded with the presence of MAPP. The TG is shifted from 1.7 °C for uncompatibilised PP/NC6 composite to 6.6 °C with the addition of 2 wt% MAPP into the system. Further additions of 5 and 8 wt% of MAPP did not change the TG significantly, but only stabilised the temperature. The addition of 6 phr NC in PP matrix resulted in a decrement in the tan δ max value relative to PP, which indicates that the materials experienced a strengthening effect[2,12]. No changes in tan δMAX has been observed with the presence of low MAPP loading (2 wt%). However, the additions of 5 and 8 wt% MAPP increased this value. Although the presence of compatibiliser is expected to increase the interfacial bonding between PP and NC, hence reduced the tan δ max , it is suggested by Lee et al.[26] that higher content of MAPP will act as a lubricating modifier at the glass transition temperature region and results in an increment in this value. The presence of GF in the system has lowered the value of tan δMAX relative to PP matrix. Even though the GF content is suggested to be the major factor in determining the tan δMAX, the interaction between GF and PP is also expected to affect the damping properties of composites. The incorporation of 2 to 8 wt% compatibiliser in PP/GF15 composites resulted in a further reduction in this value (Figure 6h). The tan δMAX for PP/GF15 is recorded at 0.061 and this value dropped to 0.041 as 8 wt% of MAPP is added. An improvement in the interfacial bonding in the composites is shown by a decrement in tan δMAX value with the addition of MAPP[2]. Nayak and Mohanty[31] have also suggested that a better adhesion between the filler and matrix will result in a lower damping property. Figure 6i shows the tan δ curves obtained for PP/GF15/NC6 composites with variation in MAPP loadings, from 2-8 wt%. A remarkable shift of TG to a higher temperature is recorded with the presence of MAPP. The TG is shifted from 1.1 °C in the uncompatibilised hybrid composite to 5.6 °C with the addition of 2 wt% MAPP into the system. Further additions of 5 and 8 wt% of MAPP only stabilised the TG value. This appreciable change implied that an improvement in the interfacial adhesion between the GF, NC and PP matrix itself has been achieved. However, the presence of the MAPP slightly increased the magnitude of tan δMAX values. This phenomenon could be related to the lubricating effect of the compatibiliser at higher MAPP content. The PP/GF15/NC6 hybrid composites with the addition of 5 wt% MAPP shows the highest value of tan δMAX (0.055) when compared with other MAPP loadings. The tan δ25°C values also show a similar trend.
4. Conclusions The FTIR spectroscopic investigations confirmed that the maleic anhydride was present in the compatibilised composites. TGA demonstrated that the thermal stability of nanocomposites was enhanced by the addition of MAPP in the system. Dynamic mechanical analysis showed that the compatibilised composites possessed higher storage modulus than the uncompatibilised systems. In addition, a remarkable shift of TG to a higher temperature was recorded with the presence of MAPP. However, the incorporation of higher content of MAPP in the PP/NC6 and PP/GF15/NC6 composites increased the tan δMAX values due to lubricating effect by the compatibiliser. On the other hand, the addition of MAPP into PP/GF15/NC6 hybrid composites enhanced 109/111 109
Rahman, N. A., Hassan, A., & Heidarian, J. better interfacial adhesion between the components, leading to enhanced thermal and dynamic mechanical properties of the resultant composites.
5. Acknowledgements The authors wish to thank the University of Malaya who supported the work reported in this paper with grants UMRG (RG343-15AFR) and BKP (BK009-2014).
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Effect of compatibiliser on the properties of polypropylene/glass fibre/nanoclay composites 27. Wang, Y., Chen, F. B., Li, Y. C., & Wu, K. C. (2004). Melt processing of polypropylene/clay nanocomposites modified with maleated polypropylene compatibilizers. Composites. Part B, Engineering, 35(2), 111-124. http://dx.doi.org/10.1016/ S1359-8368(03)00049-0. 28. Modesti, M., Lorenzetti, A., Bon, D., & Besco, S. (2006). Thermal behaviour of compatibilised polypropylene nanocomposite: effect of processing conditions. Polymer Degradation & Stability, 91(4), 672-680. http://dx.doi.org/10.1016/j. polymdegradstab.2005.05.018. 29. Lai, S. M., Chen, W. C., & Zhu, X. S. (2009). Melt mixed compatibilized polypropylene/clay nanocomposites. Part 1: the effect of compatibilizers on optical transmittance and mechanical properties. Composites. Part A, Applied Science and Manufacturing, 40(6-7), 754-765. http://dx.doi.org/10.1016/j. compositesa.2009.03.006. 30. Mandal, S., & Alam, S. (2012). Dynamic mechanical analysis and morphological studies of glass/bamboo fiber reinforced
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ISSN 1678-5169 (Online)
http://dx.doi.org/10.1590/0104-1428.02616
O O O O O O O O O O O O O O O O
Hybrids membranes with potential for fuel cells – Part 3: extruded films of nanocomposites based on sepiolite and PC/sulfonated PC blends Ana Catarina de Oliveira Gomes1*, Eduardo Henrique Backes1, Adhemar Colla Ruvolo Filho1, Caio Marcio Paranhos2, Fábio Roberto Passador3 and Luiz Antonio Pessan1 Laboratório de Permeação e Sorção – LAPS, Departamento de Engenharia de Materiais, Universidade Federal de São Carlos – UFSCar, São Carlos, SP, Brasil 2 Laboratório de Polímeros – LabPol, Departamento de Química, Universidade Federal de São Carlos – UFSCar, São Carlos, SP, Brasil 3 Instituto de Ciência e Tecnologia, Universidade Federal de São Paulo – UNIFESP, São José dos Campos, SP, Brasil 1
*acogomes@gmail.com
Abstract Fuel Cells based in polymers are an alternative for the conventional energetic matrices. However, materials currently available still present disadvantages to overcome. Membranes of polycarbonate (PC)/sulfonated polycarbonate (PCs) blend/sepiolite nanocomposites have previously been studied by the authors, resulting in good mechanical properties and promising properties of vapor transmission and ionic migration resistance. However, their production in large scale is still a challenge. The aim of this work was the development further the formulation and processing of the previously studied material. Films of PC/PCs blends (50/50 wt%) with different content of sepiolite clay, with and without chemical modification, have been prepared in an extruder and evaluated by FTIR, XRD, DSC, TGA, DMA, tension strength and water vapor transmission (WVT). Even after two processing steps, the blend-based nanocomposites keep good thermal and mechanical properties. However, changes in WVT were observed with respect to data obtained in previous studies. Keywords: polymeric membrane, nanocomposite, fuel cell, polymeric electrolyte.
1. Introduction Fuel cells are considered the future of clean sustainable energy, besides being considered the step to be overcome for the development and application of all new Technologies. The search field is powered by chronic and imminent problem of depletion of energy matrix based on fossil fuels, in addition to environmental factors that are well known and discussed in the media. Among several advantages of fuel cells compared to conventional power technologies, stand out the reduction of pollutant emissions, the higher efficiency and its simplicity[1,2]. In brief, fuel cells work oxidizing fuel at the anode, releasing electrons that are transported through an electrolyte to the cathode, where they react generating a byproduct and releasing heat. This electrolyte can be a polymeric membrane, which has as advantages the possibility of lower temperature applications, leading to a faster start of the system, smaller thickness for mounting and less weight of device, implicating in smaller and portable fuel cells, among other improvements. As disadvantages, there are the mechanical, thermal and electrochemical degradation of the polymeric membranes[1-3]. Nanocomposites of sulfonated polymers has been studied for applications on fuel cells[4,5]. The use of polycarbonate (PC) chemically modified with sulfone groups and
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incorporation of sepiolite, though, is not found in literature, to our knowledge. The bisphenol-A PC is an engineering thermoplastic, amorphous, tough, with impact resistance very superior to others thermoplastics at low temperatures[6]. Meanwhile the sepiolite is a natural hydrated magnesium silicate with a significant number of silanol groups (Si-OH) present on the surface of the mineral[7]. More information about system components can be found in our previous publications[8,9]. In previous works of our group[8,9], the sulfonation of the PC was successfully obtained, despite of generating small difference between the various contents of sulfonation agent tested. In fact, the sulfonation reaction of substituted aromatic rings have several issues regarding the extension of reaction, reversibility, and main chain scission as reported[4,10,11]. The mixture of chemically modified and not modified polymeric matrix generated a phase separation which can be beneficial to the desired application. The unmodified PC matrix is able to form a chemically and mechanical resistant structure while dispersed domains of the modified polymer can generate a percolated network to conduct protons in a more efficient and quick way. Even at low levels, the sulfonation was sufficient to promote phase separation and change the interaction between the matrix and the nanoparticles of
Polímeros, 28(2), 112-119, 2018
Hybrids membranes with potential for fuel cells – Part 3: extruded films of nanocomposites based on sepiolite and PC/sulfonated PC blends sepiolite. The chemical modification decreased thermal resistance, as expected. This effect, however, was offset by the incorporation of nanoclay, although not fully solve the problem. In the case of sepiolite, the presence of the nanoparticle assisted in conducting ionic species across the membrane, unlike the observed in the literature for other clay minerals. Sanchez et al.[12] pointed to the importance of producing conductive polymer membrane using more applied processing methods, which, in addition to using an existing structure, can be considered an ecologically friendly process by not using large amounts of solvent, such as casting or phase inversion processes. The processing in the molten state, however, has its own challenges. The thermal stability of the polymer matrix, the chemically modified matrix and clay minerals, the influence of the presence of degradation products of these chemical modifiers, the shear strength, the development and stability of the morphology of the polymer blend during processing steps, among others. Previous results[8,9] showed great potential of these nanocomposites for use as solid electrolyte and the present study aimed to test its preparation in real processing conditions in which the incorporation of the clay mineral was carried out by mixing in the molten state, using a twin screw extruder and the molding was carried out by film extrusion. The goal was to achieve an appropriate balance among processing-structure-property for an optimized performance of the resulting membranes and to verify the possibility of industrial scale production, adapting the method of production of the membranes to the currently available technological park.
2. Materials and Methods 2.1 Materials Polycarbonate (Lexan 103) used was donated by Sabic Corporation (density 1.20 g/cm3; flow index 7.0 g/10min). Dimethyldichlorosilane (DMDCS) was purchased from Sigma-Aldrich Inc, sepiolite clay from Flucka Inc and N-methyl-pyrrolidone (NMP), sulfuric acid (H2SO4), acetic anhydride (AcAn), ethyl alcohol (EtOH), toluene and dichloromethane (DCM) from Vetec Química S.A. All reagents have analytical grade and were used as received.
2.2 Methods 2.2.1 Sulfonation of PC (PCs) The chemical modification was performed using “acetyl sulfate” as sulfonation agent with 75% (H2SO4/AcAn in volume) of concentration. The polymer was dissolved in NMP and the reaction was carried at 80°C, in nitrogen atmosphere, drip addition of sulfonation agent through 1h, and another hour for reaction. The product was precipitated in EtOH and vacuum drying for 48h[8-10]. 2.2.2 Sepiolite modification The modification was carried out DMDCS 30% (volume) in toluene, refluxed in soxhlet apparel for 6 h over neat sepiolite. The modified clay was washed and filtered with Polímeros, 28(2), 112-119, 2018
chloroform and methanol, and dried at vacuum at 120°C through 24h[8,9,13]. 2.2.3 Blend and nanocomposites preparation The PC/PCs blend (50/50% in weight) and nanocomposites were prepared at molten state blend in a co-rotational twin-screw extruder MT19TC from B&P Process Equipment and Systems with diameter of 19 mm e L/D = 25, screw speed of 120 rpm, flow rate of 1 kg/h and temperature profile of 200 °C/215 °C/225 °C/230 °C/240 °C from barrel to matrix. The produced samples are described in Table 1. The sample “PC w/ Proc” is a blank reference were PC was extruded into films directly from pellets, without previous processing, and the sample “Proc. PC” is another blank reference where neat PC was extruded in the twin screw extruder and them conformed into films. 2.2.4 Films extrusion Films were produced in a flat film single screw extruder from AX Plastics, LAB16 model, with three heating zones (215/225/235 °C) from barrel to matrix. The drawing speed was 1 rpm.
2.3 Characterizations The elemental analysis was used to evaluate the sulfonation degree (DS) in PCs samples and the test was performed in Eager 200 equipment from CE Intruments. The DS content, e.g. sulfur content in polymer matrix, it is obtained from Equation 1, where MMpol is the molecular weight of Polycarbonate monomeric unit, S is the sulfur content (weight percent), 32 and 80 are the molecular weight of Sulfur and sulfonic group, respectively[8,9,14]. DS =
MM pol X 100 X S (1) 32 X (100 − 80S )
Infrared spectroscopy analyses were performed on film samples, with a FTIR Nicolet 6700 from Thermo Scientific, from 4500 to 400 cm-1 in absorbance mode. The X-ray diffraction measurements were performed using a Rigaku Rotaflex Ru 200B diffractometer, using Cu K(α) radiation (λ=1.5406 Å), at a rate of 1º/min, operating at 40 kV and 80 mA, over 2θ range of 2º-40º. The curves were submitted to a smoothing process using Savitzky-Golay method with a data window of 300 points. The tensile tests were performed with an Instron universal testing machine, 5569 model, as ASTM D882-10 (5 mm
Table 1. Samples formulation (contents in phr). Sample PC w/ Proc. Proc. PC PCs PC/PCs Nano1 Nano2 Nano3 Nano4
PC 100 100 50 50 50 50 50
PCs 100 50 50 50 50 50
Sepiolite 1.5 3.0 -
Modified sepiolite 1.5 3.0
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Gomes, A. C. O., Backes, E. H., Ruvolo Filho, A. C., Paranhos, C. M., Passador, F. R., & Pessan, L. A. / min until 4% strain, following 50 mm/min). The testing samples were cut from films (thickness from 50 - 250μm) by Franktseet cutter, in dimensions of 24 x 200 mm (width x length). The differential scanning calorimetry (DSC) was measured using a Q2000 TA Instruments equipment, nitrogen atmosphere with flux of 50 ml/min, temperature interval of 30 to 250 °C and heat rate of 10 °C / min, with two heating cycles. The thermogravimetric analysis (TGA) were performed using Q50 TA Instruments equipment, under nitrogen atmosphere, up to 800°C and heating rate of 20°C/min. The dynamic-mechanical analysis (DMA) were performed in a Q800 TA Instruments equipment, in tension film accessory, temperature ramp of 3 °C/min and deformation of 25µm at a frequency of 1 Hz. The water vapor uptake analyses were realized according to ASTM E96, with conditioning at 23 °C and humidity of 16%, during five days.
two processing steps (twin screw extrusion and flat film extrusion) the structural modification with the sulfone group insertion was maintained. It is also noted that the nanocomposite with higher clay content shows lower peak intensity (in an approximated and naked eye comparing) than other samples. This difference can be attributed to a higher interaction of the non-modified clay with the sulfone groups of the modified matrix. This same effect was observed in previous works[8,9]. For the nanocomposites with modified clay it is possible to observe a less significant effect in the band related to sulfone group, but there is a split in a band near to it (1666 cm-1) related to the main chain of polycarbonate. This split is probably related with a stronger interaction of the modified clay which is less polar, with the non-modified polycarbonate chain portion. Figure 3a shows that samples without clay present a characteristic profile of amorphous materials as polycarbonate. The sulfonation, blending without clay (not shown) and processing in molten state do not alter this pattern in X-ray diffraction.
3. Results and Discussion The Elemental Analysis of sulfonated polycarbonate returns an average content of sulfur of 1.54 wt%, corresponding to a sulfonation degree of approximately 12.7% for PCs. Figure 1 shows infrared spectra of neat PC and PCs. The characteristics absorptions bands of the sulfone group (S=O) occur at 1684, 1149 and 779 cm-1[8,9,11]. The band at 1690 cm-1 is observed in the PCs spectra indicating its sulfonation. The other bands are overlapped and cannot be differentiated. There is no difference between proc. PC and PC w/ proc. (not shown), indicating no significant changes in chemical properties due the additional processing step. Figures 2a and b shows FTIR spectra for the nanocomposites and PC/PCs blend prepared. It can be observed the characteristic peak of the sulfone group, as previously described. These results prove that even after
Figure 1. Infrared spectra of processed PC and PCs.
Figure 2. Infrared spectra of (a) PC/PCs, Nano1 and Nano2; (b) PC/PCs, Nano3 and Nano4. 114 114/119
Polímeros, 28(2), 112-119, 2018
Hybrids membranes with potential for fuel cells – Part 3: extruded films of nanocomposites based on sepiolite and PC/sulfonated PC blends Diffratograms in Figure 3b also show that nanocomposites differ from the neat matrix samples by a peak in the region of 7º, which is related to the sepiolite incorporation. This reflection is associated to spacing between crystalline planes that form the zeolitic pores of sepiolite[8]. The amorphous band referent to polymeric matrix does not alter.
3.1 Tensile test Table 2 presents the stress (δ) versus strain (ε) properties for the samples tested. The yield stress (δesc) and rupture stress (δrup) are not significantly affected by sulfonation and sepiolite insertion. The elastic modulus (E) was determined from the linear portion of the stress versus strain curve, where the Hooke Law is valid. It can be noticed an increase in elastic modulus when comparing PC w/ Proc. to Proc. PC and PC/PCs blend. This result can be correlated to the insertion of polar groups of PCs in the system that increases intersegmental interactions and results in a small increase of elastic modulus in PC/PCs blend. However this increase in stiffness led to a reduction in the elongation at break of approximately 6%. According Smitha et al.[15], the decreasing in elongation at break can also be associated to the increasing in free volume due sulfonation chemical modification. Comparing the nanocomposites to PC/PCs blend, it can be noticed an increase in the values of elastic modulus,
except for Nano3. The superior elastic modulus of Nano 1, Nano 2 and Nano 4 indicates that the sepiolite is acting as a reinforcement increasing the rigidity of the material. The smaller amount (1.5 wt%) of modified sepiolite in Nano3 is possibly located in the non-modified component of the system, due to the higher hydrophilicity in this region reducing the rigidity of the sample. When the modified clay content increases (Nano 4), the elastic modulus becomes higher than the neat blend. These suppositions are supported by DMA results, presented ahead[8,9]. It is observed an increasing in modulus in the order of 9.7% for Nano1, 11.5% for Nano2 and 7.8% for Nano4.
3.2 Diferencial Scanning Calorimetry (DSC) The glass transition temperatures (Tg) for the samples prepared were determined by DSC and the values obtained are presented in Table 3. The neat PC show a Tg value of 149.7 °C at the first heating cycle, and 149.2 °C at the second one. These values are close of the 150 °C reported by literature[16] and supplier. There is no variation of Tg for the PC with (PC Proc) and without (PC w/ proc) processing. It can be noticed a reduction of Tg in PC/PCs blend due sulfonation, which can be related to the higher free volume among polymer chains due to the high volume of the functional group of the sulfone group. The nanocomposites blends have no significant variation in Tg with the incorporation of increasing amount of clay
Figure 3. X-ray Difratograms of (a) neat polymers and PC/PCs blends; (b) nanocomposites. Table 2. Mechanical properties in tensile strength. Sample
δesc (MPa)
E (GPa)
δrup (MPa)
εrup (%)
PC w/ Proc. Proc. PC PC/PCs
53.9 ± 4.4 52.1 ± 4.4 54.3 ± 4.0
2.01 ± 0.06 2.10 ± 0.05 2.17 ± 0.06
55.8 ± 4.5 53.9 ± 4.9 54.0 ± 3.7
80.4 ± 19.2 91.6 ± 15.9 5.7 ± 0.6
Nano1 Nano2 Nano3 Nano4
55.4 ± 3.5 51.8 ± 2.3 53.5 ± 2.7 55.1 ± 1.8
2.38 ± 0.07 2.42 ± 0.06 2.03 ± 0.09 2.34 ± 0.04
55.1 ± 3.4 51.3 ± 2.8 54.7 ± 3.4 54.7 ± 2.0
10.2 ± 3.5 10.5 ± 4.4 5.7 ± 1.9 6.4 ± 0.6
Polímeros, 28(2), 112-119, 2018
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Gomes, A. C. O., Backes, E. H., Ruvolo Filho, A. C., Paranhos, C. M., Passador, F. R., & Pessan, L. A. without or with modification. However, the variation of Tg is significant when comparing nanocomposites and PC/PCs blend. The more significant difference occurs for Nano2, with 3.0% of non-modified sepiolite, showing an increase of approximately 6 °C when compared to PC/PCs blend. There is no significant variation of Tg observed in the first and second heating cycles, indicating that there was no stress relaxation from processing in the first heating cycle (as seen for the samples PC w/ Proc and Proc. PC).
3.3 Thermogravimetric Analysis (TGA) Figure 4a presents TGA curves for neat and modified clays. The neat clay shows a significant reduction of mass (approximately 8%) at 75 °C, follow by another mass loss (2%) at 250 °C. From 450 °C there is a constant reduction of mass until 600 °C, where mass variation stabilizes, resulting in a residue of 85.3% at 800 °C. These transitions agree with the loss of water described for sepiolite where four types of interactions between the mineral structure and the water molecules results in four different transitions: hygroscopic Table 3. Tg set by DSC. Sample
Tg (ºC) 1st heating
Tg (ºC) 2nd heating.
PC w/Proc. Pc Proc. PC/PCs Nano1 Nano2 Nano3 Nano4
149.7 149.9 132.7 136.0 138.3 133.0 135.2
149.2 148.4 132.0 135.8 138.1 133.5 135.5
adsorption (first transition), zeolitic adsorption (second transition), chemically adsorbed water (or linked – third transition) and structural water (a transition that would appear after 800°C, a temperature higher than reached in our test)[17,18]. The TGA curve of the modified clay shows an initial humidity loss of 4% at 65 °C, what can indicate a decrease in polarity due chemical modification. The observed behavior was similar to non-modified clay, except for an 8.5% loss at 540 °C, which can be related to the degradation of the functional groups added during chemical modification of the clay surface. This difference was also observed by Tuhan et al.[18]. Finally, a residue of 85.3% is obtained at 800 °C of the modified clay analysis. The PC/PCs blends showed similar behavior as compared to the nanocomposites (Figure 5b), only differing in the residue (Table 4) and maximum temperature of loss. There is a gradual loss of mass from 200 °C (humidity), followed by a significant mass loss in its main degradation step, around 500 °C (sulfone groups and main chain). The obtained dada for all samples are presented at Tables 4, for flakes and films forms. From the data in Table 4, it is noticed that the insertion of sulfone groups in PC chains increased slightly the degradation temperature for samples in film form. Similar behavior is found in literature[19], but this difference can be insignificant, once at this point a first degradation step (of the sulfone groups) already occurred in the molecule[9]. The clay addition reduced the maximum degradation temperature for both cases, indicating an acceleration of degradation process. In the case of processed samples (film samples),
Figure 4. TGA curves of (a) pure and modified clay; (b) PC/PCs and nanocomposites samples.
Figure 5. Loss modulus curves versus temperature for (a) PC Proc., PC s/ Proc. and PC/PCs blend; (b) PC/PCs blend and nanocomposites. 116 116/119
Polímeros, 28(2), 112-119, 2018
Hybrids membranes with potential for fuel cells – Part 3: extruded films of nanocomposites based on sepiolite and PC/sulfonated PC blends Table 4. TGA data. FLAKES Sample Pc w/Proc PC Proc. PC/PCs Nano1 Nano2 Nano3 Nano4
Mass Loss (%) 79.3 79.6 78.3 4.9 72.1 78.7
FILMS Mass Loss
Max. Loss Temp. (ºC)
Residue (%)
538.4 514.3 521.9 463.9 518.2 519.0
21.0 20.6 21.8
(%) 78.3 78.5 77.8 79.7 77.9
20.7 21.5
Max. Loss Temp. (ºC)
Residue
523.5 529.9 526.2 514.7 517.8
(%) 22.1 21.7 22.4 20.3 22.1
77.9
520.4
21.9
77.8
515.3
22.0
Table 5. Water permeation data. Sample PC w/ Proc. Proc. PC PC/PCs Nano1 Nano2 Nano3 Nano4
WVT (g.day-1.m-2) Extrusion Casting[9] 28.0 ± 1.0 54 ± 3 18.0 ± 2.0 16.6 ± 0.8 194 ± 5 15.2 ± 0.7 180 ± 36 15.2 ± 0.9 49 ± 23 17.0 ± 1.0 75 ± 59 13.6 ± 0.5 200 ± 84
this reduction is smaller. There are reports in the literature of the occurrence of degradation due to the addition of clay[20]. The degradation temperature decreased with the increase of the clay content, and increased again for the sample with a higher content of modified clay (Nano 4). This behavior suggests a saturation limit for the influence of the presence of the clay[9]. Flakes were the sample form before film extrusion, and its analysis is important to states if there is a difference in thermal properties after and before the last processing step. It was observed that the sample residue is kept practically constant for the both sample types. It is important to recall that is not characteristic of PC to degrade with complete conversion into volatiles products in the case of thermogravimetrics tests performed in inert atmosphere, having always a level of residual ashes. This fact explains the high residue content, which corresponds not only to the inorganic material present in the sample, but also the ash content inherent to the PC. The differences among temperatures of maximum mass loss between samples forms are due the geometric factor of the samples, where more compact samples, as films, have a heat transfer more efficient and present lower temperatures of maximum loss. This fact can be considered also as an indication that the last processing step does not significantly alter the sample, i.e., does not causes significant degradation.
3.4 Dynamic-Mechanical Analysis (DMA) The most common use of DMA technique is the determination of Tg and secondary transitions related with small chain portions or side groups of the polymeric chain. The technique, however, is generally underestimated. DMA results are sensible enough to be also useful to determine Polímeros, 28(2), 112-119, 2018
P (E10) (g.Pa-1.s-1.m-1) Extrusion Casting[9] 9.4 ± 0.6 8.1 ± 0.4 9.0 ± 2.0 7.0 ± 0.4 14 ± 5 7.7 ± 0.3 44 ± 24 7.7 ± 0.4 12 ± 5 8.0 ± 1.0 7.0 ± 2 7.3 ± 0.4 24 ± 2
phase separation and interaction among components of a blend[21,22] and therefore can be used in morphology discussions. Figure 5a presents the loss modulus for samples of processed PC, PC without processing and PC/PCs blend. It can be noticed that the Tg of both neat PC are close to 150 °C, as expected and corroborated by DSC tests. The neat samples present sub-Tg transition around 80-90 °C. The PC/PCs blend presented two peaks below Tg, being the first around 30 °C, and the second around 80 °C. These peaks are related to intersegment motion of PC chain and formation of PCs phase, respectively. These results indicated that the sulfonation changes significantly the polymer chains mobility, leading to a reduction of the Tg, as well as phase separation between sulfonated polycarbonate and non-modified polycarbonate[8,9]. Figure 5b shows loss modulus for PC/PCs blend and the nanocomposites. It is noticed that the nanocomposites do not show a peak in 91 °C, differently from PC/PCs blend. A possible explanation is the clay addition shift the peak for higher temperatures values, where it overlaps to the Tg peak of the PC. All nanocomposites presented a secondary transition around 25 °C, related with relaxation of chain segments of the PC non-modified main chain. Samples Nano3 and Nano4 showed peaks of higher intensity indicating higher interaction of the modified Clay with the non-sulfonated phase[8,9].
3.5 Water vapor uptake The transport of small molecules through a homogeneous membrane is usually described by a process where the permeant molecules adsorb at the surface, migrate to the opposite surface due to a concentration gradient and evaporate from the posterior surface to the environment. Therefore, the permeation of vapor through polymeric materials involves distinct processes of sorption, diffusion and desorption of the penetrant[23]. 117/119 117
Gomes, A. C. O., Backes, E. H., Ruvolo Filho, A. C., Paranhos, C. M., Passador, F. R., & Pessan, L. A. The transport properties of a gas, vapor or liquid in a polymer material are ruled by several factors, some of them depending of the properties of the penetrating species (in this case, the gas or vapor), the properties of the polymer and the interaction degree between them. The environment conditions are important as well. Among the influencing factors, stand out the chemical nature of the polymer and the permeant molecules, that influences in the diffusion and solubility parameters and consequently in permeation, and also the polymer phase morphology. Table 5 presents the average values of water vapor transmission (WVT) and permeability (P). For comparing, data for casting films obtained in previous work[9] are presented too. The first difference to be discussed is the higher permeation of water for casting films than processed films. It is common knowledge that solvent evaporating process can generate more voids into membranes than the molten state processing. Firstly, the slow exit of solvent permits molecular mobility through longer times, what can lead to ordering of molecules, crystallization and phase separation. Secondly, the exit of solvent can lead defects to film core and surface, depending on drying velocity of the casting solution[24,25]. All these factors can be responsible by higher permeation of the casting films. The insertion of bulky groups in polymeric chain by sulfonation increases the free volume, what usually facilitates the transport of permeant through the membrane. The results obtained, however, show a reduction in WVT and P with the sulfonation. We believe that the increase in polarity is acting toward to retain the humidity into membrane phase[8,9]. This effect is observed in both processed and casting films. The sepiolite addition is believed to have the same effect of sulfonation in nanocomposites which is to facilitate water transport through membranes by forming preferential channels. However, the data shows that the addition of clays reduces slightly the WVT values in all cases for processed films. A probable explanation for this behavior is that the sepiolite is well dispersed and distributed into polymeric matrix, increasing the tortuosity of the path the water vapor molecules need to take to pass through the film sample. These effects can be associated to the processing method, which do not permit molecular rearrangement and migration through phases, freezing a less ordered structure of the blend. In general, the results are very different between processed and cast films, indicating that the material processing has a great role in blend morphology, resulting in very different transport properties, by reasons already discussed. Besides, the difference in small molecules transmission properties among nanocomposites are higher for samples prepared through casting process, indicating a reduced influence of the chemical nature of the clay surface into the blend morphology development, what is in agreement with the rapid establishment of film structure in processed films.
4. Conclusions Membranes of nanocomposites blends of PC/PCs with sepiolite were produced with good optical properties, good mechanical resistance and with maintenance of the sulfone sites, as detected by FTIR and thermal analysis. Phase separation in the systems studied was confirmed by DMA. 118 118/119
The WVT tests showed that the melt state processing play an important role in the clay-matrix interaction as compared to samples prepared by casting procedure, decreasing water permeation and indicating that the transposition of casting process for the preparation of polymer membranes to a melt extrusion process needs more detailed and specific development.
5. Acknowledgements The authors are grateful to the funding agencies Coordination for the Improvement of Higher Education Personnel (CAPES), National Council for Scientific and Technological Development (CNPq) and São Paulo Research Foundation (FAPESP) by financial support, and SABIC by donation of PC.
6. References 1. Larminie, J., & Dicks, A. (2003). Fuel cell systems explained. Chichester: John Wiley. http://dx.doi.org/10.1002/9781118878330. 2. O’Hare, R. P., Cha, S. W., Colella, W., & Prinz, F. B. (2006). Fuel cells fundamentals. New York: John Wiley. 3. Wu, J., Yuan, X. Z., Martin, J. J., Wang, H., Zhang, J., Shen, J., Wu, S., & Merida, W. (2008). A review of PEM fuel cell durability: Degradation mechanisms and mitigation strategies. Journal of Power Sources, 184(1), 104-119. http://dx.doi. org/10.1016/j.jpowsour.2008.06.006. 4. Pinto, B. P., Santa Maria, L. C., & Sena, M. E. (2007). Sulfonated poly(ether imide): A versatile route to prepare functionalized polymers by homogenous sulfonation. Materials Letters, 61(1112), 2540-2543. http://dx.doi.org/10.1016/j.matlet.2006.09.060. 5. Park, C. H., Lee, C. H., Guiver, M. D., & Lee, Y. M. (2011). Sulfonated hydrocarbon membranes for medium-temperature and low-humidity proton exchange membrane fuel cells (PEMFCs). Progress in Polymer Science, 36(11), 1443-1498. http://dx.doi.org/10.1016/j.progpolymsci.2011.06.001. 6. Ehrenstein, G. W., & Kabelka, J. F. (1992). Reinforced plastics. In F. Ullmann (Ed.), Ullmann’s encyclopedia of industrial chemistry (Vol. 28, Cap. 8, pp. 603-612). Berlin: VCH Publishers. 7. Ruiz-Hitzky, E. (2001). Molecular access to intracrystalline tunnels of sepiolite. Journal of Materials Chemistry, 11(1), 86-91. http://dx.doi.org/10.1039/b003197f. 8. Gomes, A. C. O., Uieda, B., Tamashiro, A. A., Ruvolo Filho, A. C., Pessan, L. A., & Paranhos, C. M. (2014). Membranas híbridas com potencial uso em células a combustível - parte 1: nanocompósitos de poli(eterimida) sulfonada. Polímeros: Ciência e Tecnologia, 24(4), 464-473. http://dx.doi.org/10.1590/01041428.1131. 9. Gomes, A. C. O., Machado, I. M. M., Ruvolo, A. C., Fo., Pessan, L. A., & Paranhos, C. M. (2014). Membranas híbridas com potencial uso em células a combustível - parte 2: nanocompósitos de poli(carbonato) sulfonado. Polímeros: Ciência e Tecnologia, 24(3), 402-410. http://dx.doi.org/10.4322/polimeros.2013.049. 10. Lakshmi, R. T. P. S., Bhattacharya, S., & Varma, I. K. (2006). Effect of sulfonation on thermal properties of poly (ether imide). High Performance Polymers, 18(2), 115-126. http:// dx.doi.org/10.1177/0954008306056503. 11. Pinto, B. P., Santa Maria, L. C., & Sena, M. E. (2007). Sulfonated poly(ether imide): a versatile route to prepare functionalized polymers by homogenous sulfonation. Materials Letters, 61(3), 2540-2543. http://dx.doi.org/10.1016/j.matlet.2006.09.060. 12. Sanchez, J.-Y., Chabert, F., Iojoiu, C., Salomon, J., El Kissi, N., Piffard, Y., Marechal, M., Galiano, H., & Mercier, R. (2007). Extrusion: an environmentally friendly process for Polímeros, 28(2), 112-119, 2018
Hybrids membranes with potential for fuel cells – Part 3: extruded films of nanocomposites based on sepiolite and PC/sulfonated PC blends PEMFC membrane elaboration. Electrochimica Acta, 53(4), 1584-1595. http://dx.doi.org/10.1016/j.electacta.2007.04.022. 13. Alkan, M., Tekin, G., & Namli, H. (2005). FTIR and zeta potential measurements of sepiolite treated with some organosilanes. Microporous and Mesoporous Materials, 84(1-3), 75-83. http:// dx.doi.org/10.1016/j.micromeso.2005.05.016. 14. Genies, C., Mercier, R., Sillion, B., Cornet, N., Gebel, G., & Pineri, M. (2001). Soluble sulfonated naphthalenic polyimides as materials for proton exchange membranes. Polymer, 42(2), 359-373. http://dx.doi.org/10.1016/S0032-3861(00)00384-0. 15. Smitha, B., Sridhar, S., & Khan, A. A. (2003). Synthesis and characterization of proton conducting polymer membranes for fuel cells. Journal of Membrane Science, 225(1-2), 63-76. http://dx.doi.org/10.1016/S0376-7388(03)00343-0. 16. Abts, G., Eckel, T., & Wehrmann, R. (1992). Polycarbonates. In F. Ullmann (Ed.), Ullmann’s encyclopedia of industrial chemistry (Vol. 21, Cap. 2, pp. 207-214). Berlin: VCH Publishers. 17. Hevesut, H., Otsuka, H., & Imai, N. (1969). Infrared study of sepiolite and palygorskite on heating. The American Mineralogist, 53(nov-dec), 1613-1624. 18. Turhan, Y., Turan, P., Doĝan, M., Alkan, M., Namli, H., & Demirbas, O. (2008). Characterization and adsorption properties of chemically modified sepiolite. Industrial & Engineering Chemistry Research, 47(6), 1883-1895. http:// dx.doi.org/10.1021/ie070506r. 19. Hande, V. R., Rath, S. K., Rao, S., & Patri, M. (2011). Crosslinked sulfonated poly (ether ether ketone) (SPEEK)/reactive organoclay nanocomposite proton exchange membranes (PEM).
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Journal of Membrane Science, 372(1-2), 40-48. http://dx.doi. org/10.1016/j.memsci.2011.01.042. 20. de la Orden, M. U., Pascual, D., Antelo, A., Arranz-Andrés, J., Lorenzo, V., & Martínez Urreaga, J. (2013). Polymer degradation during the melt processing of clay reinforced polycarbonate nanocomposites. Polymer Degradation & Stability, 98(5), 11101117. http://dx.doi.org/10.1016/j.polymdegradstab.2013.03.024. 21. Lucas, E. F., Soares, B. G., & Monteiro, E. E. C. (2001). Caracterização de polímeros - determinação de peso molecular e análise térmica. Rio de Janeiro: e-Papers. 22. Sepe, M. P. (1998). Dynamic mechanical analysis for plastics engineering. New York: Plastics Design Library. 23. Chinellato, A. C., Vidotti, S. E., Moraes, M. B., & Pessan, L. A. (2007). Effects of plasma etching on surface modification and gas permeability of bisphenol-a polycarbonate films. Journal of Macromolecular Science, Part B: Physics, 46(6), 1165-1177. http://dx.doi.org/10.1080/00222340701582928. 24. Qipeng, G., editor (2016). Polymer morpohology. principles, characterization, and processing. New Jersey: John Wiley & Sons Inc. 25. Baker, R. W. (1991). Membrane separation systems: recent developments and future directions. Michigan: Noyes Data Corporation. Received: Mar. 03, 2016 Revised: Oct. 04, 2016 Accepted: Nov. 07, 2016
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ISSN 1678-5169 (Online)
http://dx.doi.org/10.1590/0104-1428.11216
O O O O O O O O O O O O O O O O
Fabrication of poly(lactic acid) incorporated chitosan nanocomposites for enhanced functional polyester fabric Zulfiqar Ali Raza1* and Faiza Anwar1,2 Department of Applied Sciences, National Textile University, Faisalabad, Pakistan School of Textile and Design, University of Management and Technology, Lahore, Pakistan 1
2
*zarazapk@yahoo.com
Abstract This study dealt with the fabrication and analysis of poly(lactic acid) (PLA) mediated chitosan nanocomposite. Such a novel nanobiocomposite may get future applications in drug delivery, and nanofinising of textile and polymer surfaces. Herein, this nanocomposite has been considered as an antibacterial finishing agent for a hydrophobic textile fabric like polyester. The prepared nanocomposite was characterized for zeta size and morphology, and subsequently applied on a woven polyester fabric though a coross linker. The treated polyester fabric was analyzed for textile functional characteristics as well asantibacterial activity. The spectral and optical properties demonstrated that the nanocomposite developed exhibited spherical morphologies with a mean nano particle size of ca. 88 nm. The treated fabric projected satisfactory antibacterial and fair fabric attributes. Hence, the nanofinished polyester fabric is a potential biocompatible candidate as medical and antibacterial textiles in addition to be used in antibacterial water filtration and materials packing. Keywords: antibacterial, chitosan, cotton fabric, nanocomposite, PLA.
1. Introduction Biodegradable polymers, like poly(lactic acid) (PLA), have much attentions in recent years because of their excellent biocompatibility and low toxicity even after degradation. PLA is finding application as drug carrier due to its good bioresorptivity and elimination on drug depletion[1,2]. PLA is an example of bio-polyesters i.e., bio-based thermoplastic which in addition biodegradable are eco-compatible. Such bio-derived polymers reduce the carbon cycle by repaying plant-based carbon to the soil via biodegradation and alternatively reduce the environmental impact; hence, they reduce carbon dioxide emission during their life cycle. The PLA is extensively used in tissue engineering, bioscaffolds, drug delivery systems, membranes, cosmetic, biomedical implants and son on[3]. Chitosan (CS) is an innate biodegradable polymer attained by alkaline N-deacetylation of chitin and it is the 2nd most available natural polymer after cellulose. It finds diverse applications in medical sctor. Chitosan nanoparticles are prefer over powder or bulk because nanoparticles have potential of slow/controlled release of drug and other textile finishe[4]. The chitosan based nanoparticles (NPs) have gained enormous attention because of their good environmental compatibility and prompt modifiability[5,6]. It has broad spectrum antibacterial activity[7]. Some studies have suggested that the degree of deacetylation and molecular mass of chitosan can be altered to obtain desirable charcteristics. On the other hand, chitosan can also be employed to coat different NPs of other materials to enhance their effectiveness for the body and increase their bioavailability[8,9]. In recent decades, chitosan based composite namomaterails have been explored. Chitosan is versatile material that can be
120 120/124
used to prepared chitosan based different NPs likemediation chitosan with zinc oxide and other biopolymer for instance polylactic acid (PLA) to obtain different properties that can potentially be used in biomedical and drug delivery applications[10]. Bsiides PLA, other candidate mey be poly(glycolic acid) (PGA)[11-13]. Textile materials have specific mechanical properties which make them a unique material for society. Nanoscience has recently introduced some nanofinishes for textile substrates. Textile finishes based on nanotechnology provides oil and water repellency, antibacterial activity and UV protection[14,15]. The PLA nanofinishes are gaining importance in recent years because of fact that PLA is obtained from renewable resources and have ability to incorporate mechanical strength, antibacterial and ultraviolet protection to textiles[16]. The PLA is one of the promising polymers in which stereochemical structure can easily be modified through the process of polymerization in order to form homo- and co-polymers. It is biocompatible and biodegradable. The biocompatibility of PLA is a useful property in biomedical application while biodegradability is adapted for packing purposes. The aim of this study was to prepare the PLA‑CS based NPs and their application on polyester fabric. The characterization of composite NPs was done by using the advanced analytical techniques like scanning electron microscopy, FTIR, zeta size and zeta potential whereas the treated polyester fabric was evaluated for its functional and textile performance properties.
Polímeros, 28(2), 120-124, 2018
Fabrication of poly(lactic acid) incorporated chitosan nanocomposites for enhanced functional polyester fabric
2. Materials and Methods
2.4 Fabric impregnation of nanocomposites
2.1 Materials
The PLA/CS nanocomposite (as 1 to 3% w/v) was ultrasonicated for 10 min at room temperature to get uniform aqueous suspension. A low temperature cross-linking agent of Knittex RCT was used as 80 g/L and a catalyst Knittex MO as 20 g/L to bind the nanocomposite to the polyester fabric. To this suspension, the fabric specimens (at a fabric to liquor ratio of 0.05) were padded in two dips and nips at 100% wet pick up. The specimens were dried at 110°C for 1 min and cured at 150°C for 3 min.
Chitosan, with >90% purity and degree of deacetylation of 90% and viscosity 200-800 cps, was procured from Bio Basic Inc., Canada and acetic acid (congealing temperature ≥15.6 °C, limit of nonvolatile residue ≤1.0 mg and heavy metal ≤5 ppm) from Acros Organics. Polylactic acid (melting point 150-160 °C and density 1 .21-1.43 g·cm−3) and sodium hydroxide (molecular mass 39.997109 g/mol)from Sigma‑Aldrich. Sodium tripolyphosphate (STPP, Mol. mass 367.86 amu) was purchased from BioM Lab, USA. Nutrient agar was procured from Merck and nutrient broth from Lab. M. Ltd., UK. Knittex RCT (modified dihydroxy ethylene urea) and Knittex MO (magnesium chloride basaed) were kindly donated by the local office of Huntsman. The properties of polyester fabric are shown in Table 1.
2.2 Fabrication of nanocomposite PLA/CS NPs were synthesized by ionic gelation method. Briefly; CS was dissolved in 0.5% acetic acid. Different concentrations of PLA from 1-3% were dissolved in chloroform, which was then added into the 1-2% CS solution with different volumes. 1% STPP solution was then added after ultrasonication for 30 min at 30°C. Blank PLA/CS NPs were gradually formed under vigorous stirring condition. After that the sample was centrifuged at 6500 rpm at 4°C for 30 min and resultant pallets were freeze-dried to obtain the required composite NPs. The design of experiments (DOE) of preparation of PLA incorporated CS NPs is shown in Table 2.
2.3 Size determination of nanocomposite The average diameter of PLA/CS NPs was measured using dynamic light scattering technique using a Zetasizer Nano (Malvern ZEM-3600, Malvern instruments Ltd., UK) at 25°C. Table 1. Properties of used polyester fabric. Standard method Fabric construction Plain weave (1×1) AATCC-20 Surface density, g/m2 97 ASTM D-3776 Ends/inch 130 ASTM D-3775 Picks/inch 70 ASTM D-3775 Warp count 32 ASTM D-1059 Weft count 32 ASTM D-1059 Tensile strength (warp), KgF 34 ASTM D-5034 Tensile strength (weft), KgF 32 ASTM D-5034 CIE whiteness 70°CIE AATCC-110 Bending length (warp), cm 0.71 ASTM D-1388 Property
Values
2.5 Characterization of nanocomposites treated fabrics The prepared nanocomposite particles were spherical in shape with the mean size of 88 nm measured as number distribution (Figure 1). The surface morphology of the nanocomposite applied polyester fabic samples were examined under an FEI Quanta 250 scanning electron microscope (SEM) at an accelerating voltage of 20 kV. The fabric impregnation of NPs on fabric was examined using Attenuated Total Reflectance-Fourier transform infrared spectroscopy (ATR-FTIR, Bruker Tensor 27).
2.6 Textile performance properties Bnding length of fabric samples was measured as per standard protocol of ASTM D-1388 on a cantilever, crease recovery angle (CRA) as per standard protocol of ISO 2313 and tensile strength as per standard protocol of ASTM D-5034. The bacterial strains of Escherichia coli and Staphylococcus aureus were employed to estimate the antibacterial activity of treated samples as per documented earlier[17].
3. Results and Discussions The morphology of the control fabric sample and after implication of PLA-CS NPs was detected under SEM. The SEM images revealed the fact that the control fabric has smooth surface (Figure 2a); while, the treated one indicated the impregnation of particles on fabric surface (Figure 2b). The FTIR results (Figure 3a) show that the peak at 1701 cm-1 indicated the presence of C=O, an another peak in the range 1016 to 1239 cm-1 corresponds to C-O-C band stretching; hence confirm the ester linkage in polyester fabric. Whereas, the treated fabric showed the peaks at 1407 cm-1 regarding to C-NH2amine band (Figure 3b). The peak at 1071-1016 cm-1 is the indication of C-O band stretching. While the peaks appearance at 1710 cm-1 is assign to the C=O.
Table 2. DOE for PLA incorporated chitosan nanoparticles preparation. Chitosan (%w/v) 1.0 1.5 2.0
TPP
PLA
(Sodium tri polyphosphate) (%) 1.0 1.5 2.0
(Poly lactic acid) (%) 1 2 3
Polímeros, 28(2), 120-124, 2018
Figure 1. Particle size distributions of PLA-CS nanocomposite. 121/124 121
Raza, Z. A., & Anwar, F.
Figure 2. SEM images of (a) untreated and (b) PLA-CS nanocomposite treated polyester fabrics.
Figure 3. FTIR spectra of (a) untreated and (b) PLA-CS nanocomposite treated polyester fabrics.
The bending length of polyester sample after treatment with PLA/CS NPs showed an overall increment. Figure 4a shows that in the case of control sample, warp and weft bending lengths were lower while after treatment there observed a significant increase in warp and weft bending lengths. The warp bending length was higher than weft due to presence of size material. The overall increment was due to the presence of chitosan which enhance the stiffness of fabric and increase stiffness of fabric[18]. 122 122/124
Figure 4. Effect of PLA-CS nanocomposite on (a) bending length, (b) CRA and (c) tensile strength of polyester fabric.
The minimal increase in CRA values, provide improved easy care properties. Figure 4b shows an increase in CRA as compare to untreated sample the reason behind that the PLA-CS nanocomposite fills the amorphous regions and causes to improve the CRA[19]. Figure 4c shows that after the application with PLA-CS nanocomposite, the tensile strength of the treated fabric Polímeros, 28(2), 120-124, 2018
Fabrication of poly(lactic acid) incorporated chitosan nanocomposites for enhanced functional polyester fabric
Figure 5. Antibacterial activity of PLA-CS nanocomposite treated fabric against (a) S. aureus and (b) E. coli, and (c) respective quantitative antibacterial activities.
improved on increasing the nanocomposite contents (1-3% w/v). The overall increment as compare to control fabric is due to the binding of PLA nanocomposite between fiber and yarn[18]. Chitosan is being used worldwide because of its antimicrobial propertie. The zone of inhibition could not be expected if the antibacterial agent were firmly attached to the fabric surface (e.g. covalently) which could prevent the diffusion into the nutrient agar. Nevertheless, if the antibacterial agent could diffuse into the nutrient agar, a zone of inhibition became evident and its size indicated the potency of the antibacterial agent. The antibacterial action of polyester fabric after implication of finish was checked and Figure 5a, b shows that there is clear zone of inhibition against the tested strains of S. aureus and E. coli. The zone of inhibition was higher against E. coli. The results show that the antibacterial activity was higher in case of 1 and 3% (w/v) PLA-CS nanocomposite because the higher concentration of antibacterial agent (i.e., chitosan and PLA) adsorb on the cell wall of bacteria and ultimately cause the destruction of cell membrane and death of bacteria.
4. Conclusions The presented study under took finishing of 100% polyester fabrics with some combinations of chitosan and polylactic acid (PLA) to impart functional attribtes in treated fabrics. The impregnation of PLA incorporated chitosan nanocomposite on the polyester fabric was authenticated by, SEM and FTIR spectroscopy. The particle size was confirmed by using zeta sizer. All combinations of PLA-CS showed inhibition against bacterial strains, viz. E. coli and S. aureaus but zone of inhibition increased by increasing nanobiocomposite. On combining PLA with chitosan, the crease recovery properties of treated polyester fabric improved Polímeros, 28(2), 120-124, 2018
in addition to the antibacterial activity. The CRA of treated polyester fabrics was almost comparable to control sample.
5. Acknowledgements The authors credit the financial support by Higher Education Commission of Islamabad.
6. References 1. Colon, G., Ward, B. C., & Webster, T. J. (2006). Increased osteoblast and decreased Staphylococcus epidermidis functions on nanophase ZnO and TiO2. Journal of Biomedical Materials Research. Part A, 78(3), 595-604. PMid:16752397. http:// dx.doi.org/10.1002/jbm.a.30789. 2. Reddy, K. M., Feris, K., Bell, J., Wingett, D. G., Hanley, C., & Punnoose, A. (2007). Selective toxicity of zinc oxide nanoparticles to prokaryotic and eukaryotic systems. Applied Physics Letters, 90(213902), 2139021-2139023. http://dx.doi. org/10.1063/1.2742324. PMid:18160973. 3. Dhandayuthapani, B., Yoshida, Y., Maekawa, T., & Kumar, D. S. (2011). Polymeric scaffolds in tissue engineering application: A review. International Journal of Polymer Science, 2011, 1-19. http://dx.doi.org/10.1155/2011/290602. 4. Wang, J. J., Zeng, Z. W., Xiao, R. Z., Xie, T., Zhou, G. L., & Wang, S. L. (2011). Recent advances of chitosan nanoparticles as drug carriers. International Journal of Nanomedicine, 6, 765774. http://dx.doi.org/10.2147/IJN.S17296. PMid:21589644. 5. Dong, Y., Ng, W. K., Shen, S., Kima, S., & Tana, R. B. H. (2013). Scalable ionic gelation synthesis of chitosan nanoparticles for drug delivery in static mixers. Carbohydrate Polymers, 94(2), 940-945. PMid:23544653. http://dx.doi.org/10.1016/j. carbpol.2013.02.013. 6. Sousa, F., Guebitz, G. M., & Kokol, V. (2009). Antimicrobial and antioxidant properties of chitosan enzymatically functionalized with flavonoids. Process Biochemistry, 44(7), 749-756. http:// dx.doi.org/10.1016/j.procbio.2009.03.009. 123/124 123
Raza, Z. A., & Anwar, F. 7. Goy, R. C., Britto, D. D., & Assis, O. B. G. (2009). A review of the antimicrobial activity of chitosan. Polímeros: Ciência e Tecnologia, 19(3), 241-247. http://dx.doi.org/10.1590/S010414282009000300013. 8. Janaa, S., Majia, N., Nayakb, A. K., Sena, K. K., & Basua, S. K. (2013). Development of chitosan-based nanoparticles through inter-polymeric complexation for oral drug delivery. Carbohydrate Polymers, 98(1), 870-876. PMid:23987423. http://dx.doi.org/10.1016/j.carbpol.2013.06.064. 9. Alishahi, A., Mirvaghefi, A., Tehrani, M. R., Farahmand, H., Shojaosadati, S. A., Dorkoosh, F. A., & Elsabee, M. Z. (2011). Shelf life and delivery enhancement of vitamin C using chitosan nanoparticles. Food Chemistry, 126(3), 935-940. http://dx.doi. org/10.1016/j.foodchem.2010.11.086. 10. Jiang, T., Khan, Y., Nair, L. S., Abdel-Fattah, W. I., & Laurencin, C. T. (2010). Functionalization of chitosan/poly(lactic acidglycolic acid) sintered microsphere scaffolds via surface heparinization for bone tissue engineering. Journal of Biomedical Materials Research. Part A, 93(3), 1193-1208. http://dx.doi. org/10.1002/jbm.a.32615. PMid:19777575. 11. Braz, J. (2011). Effect of the molecular weight on the physicochemical properties of poly(lactic acid) nanoparticles and on the amount of ovalbumin adsorption. Journal of the Brazilian Chemical Society, 22(12), 2304-2311. http://dx.doi. org/10.1590/S0103-50532011001200010. 12. Huang, L., Hu, J., Lang, L., Wang, X., Zhang, P., Jing, X., Wang, X., Chen, X., Lelkes, P. I., Macdiarmid, A. G., & Wei, Y. (2007). Synthesis and characterization of electroactive and biodegradable ABA block copolymer of polylactide and aniline pentamer. Biomaterials, 28(10), 741-751. PMid:17218007. http://dx.doi.org/10.1016/j.biomaterials.2006.12.007. 13. Li, L., Ding, S., & Zhou, C. (2003). Preparation and degradation of PLA/chitosan composite materials. Journal of Applied Polymer, 91(1), 274-277. http://dx.doi.org/10.1002/app.12954. 14. Dev, A., Binulal, N. S., Anitha, A., Nair, S. V., Furuike, T., Tamura, H., & Jayakumar, R. (2010). Preparation of poly(lactic
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acid)/chitosan nanoparticles for anti-HIV drug delivery applications. Carbohydrate Polymers, 80(3), 833-838. http:// dx.doi.org/10.1016/j.carbpol.2009.12.040. 15. Raza, Z. A., Rehman, A., Mohsin, M., Bajwa, S. Z., Anwar, F., Naeem, A., & Ahmad, N. (2015). Development of antibacterial cellulosic fabric via clean impregnation of silver nanoparticles. Journal of Cleaner Production, 101, 377-386. http://dx.doi. org/10.1016/j.jclepro.2015.03.091. 16. Babu, A., Wang, Q., Muralidharan, R., Shanker, M., Munshi, A., & Ramesh, R. (2014). Chitosan coated polylactic acid nanoparticle-mediated combinatorial delivery of cisplatin and siRNA/plasmid DNA chemosensitizes cisplatin-resistant human ovarian cancer cells. Molecular Pharmaceutics, 11(8), 27202733. PMid:24922589. http://dx.doi.org/10.1021/mp500259e. 17. Mohsin, M., Farooq, U., Raza, Z. A., Ahsan, M., Afzal, A., & Nazir, A. (2014). Performance enhancement of wool fabric with environmentally-friendly bio cross-linker. Journal of Cleaner Production, 68, 130-134. http://dx.doi.org/10.1016/j. jclepro.2013.12.083. 18. Arain, R. A., Khatri, Z., Memon, M. H., & Kim, I. S. (2013). Antibacterial property and characterization of cotton fabric treated with chitosan/AgCl-TiO2 colloid. Carbohydrate Polymers, 96(1), 326-331. PMid:23688488. http://dx.doi. org/10.1016/j.carbpol.2013.04.004. 19. Sunder, A. E., Nalankilli, G., & Swamy, N. K. P. (2014). Multifunctional finishes on cotton textiles using combination of chitosan and polycarboxylic acids. Indian Journal of Fibre and Textile Research, 39, 418-424. Received: Sept. 16, 2016 Revised: Nov. 07, 2016 Accepted: Dec. 01, 2016
Polímeros, 28(2), 120-124, 2018
ISSN 1678-5169 (Online)
http://dx.doi.org/10.1590/0104-1428.04516
Properties of barrier shrink bags made with EVOH and polyamide for fresh beef meat preservation Jose Boaventura Rodrigues1, Kleber Brunelli1, Claire Isabel Grígoli de Luca Sarantopoulos2* and Lea Mariza de Oliveira2 DuPont Packaging & Industrial Polymers, Barueri, SP, Brazil Instituto de Tecnologia de Alimentos – ITAL, Centro de Tecnologia de Embalagem – CETEA, Campinas, SP, Brazil 1
2
*claire@ital.sp.gov.br
Abstract The objective of this work was to compare the barrier and mechanical properties and shrinkability of coextruded films chlorine-free, with high barrier chlorine multilayer films traditionally used to preserve fresh beef. Four 9-layer barrier‑shrink films containing PET, ethylene ionomers, polyamide PA66/6 modified with amorphous PA, 32 or 44 mol% EVOH and PE were produced in a commercial scale triple bubble co-extrusion line. Seal strength, puncture resistance, oxygen and water vapor permeability and film shrink were measured for the four films and compared to the EVA/PVDC/PE film properties. The results obtained under controlled laboratory conditions show that films made with one layer of EVOH 32 mol% of ethylene encapsulated between two layers of PA66/6 modified with amorphous PA have gas barrier properties and puncture resistance better than a typical EVA/PVDC/PE, seal strength and shrinkability comparable to this film and therefore have potential to preserve fresh beef. Keywords: chlorine free, oxygen transmission rate, puncture resistance, shrink film, vaccum packaging.
1. Introduction Trends in the food market clearly show an increasing demand for healthier and safer food and the need for packages with lower environmental impact[1]. In regards to health the food industry is striving to provide fresh products with reduced amounts or no preservatives that meet stringent safety requirements in the globalized market. Beef is considered fresh if it is recently processed, vacuum-packed or packed in modified atmospheric gases, and has not undergone any treatment other than chilling to ensure preservation[2]. To avoid undesired changes in appearance, odor, texture, and flavor due to microbial activity or interaction with the environment, the packaging material used must be able to enclose the meat cuts and maintain the ideal atmospheric environment inside the package. Therefore the packaging must provide a hermetic and reliable closure, must have the ability to retain vacuum and to minimize gas transfer through the film surface to maintain a low oxygen partial pressure to reduce oxidative reactions and aerobic bacteria growth[3]. Plastic films with structural strength and shrink ability are used for wrapping uneven cuts of fresh meat to achieve a skin-tight and compact pack. The skin-tight feature is also effective to prevent liquid purge from inside the muscle tissue. The advantages of plastic shrinkable films include ease handling, a contour fit and neat appearance[4]. Traditionally vacuum packaging bags are designed to optimize gas barrier, shrinking properties, toughness and sealing characteristics, among other features[5]. Those properties are highly dependent on the resins used, the
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manufacturing technology and the actual structure of the multilayer film[6]. According to Zhou et al.[2], vacuum packages for fresh beef are usually coextruded multilayer films composed of ethylene-vinyl acetate copolymer (EVA) and polyvinylidene chloride-methyl acrylate copolymer (PVDC) which have oxygen permeability lower than 15.5 mL (STP).m-2.day-1. In Brazil, the market for barrier shrink film for vaccum beef package is dominated by two global manufactures, that sells films with OTR lower than 25 mL (STP).m-2.day-1. Those films are treated with radiation along the conversion process so they become temperature sensitive and shrink when subjected to temperatures ranging between 80 ºC and 90 oC. Although shrinkable high barrier EVA/PVDC/PE (polyethylene) films are widely used and very effective to preserve fresh beef, they are considered not eco-friendly because they contain chlorine which produces dioxin during combustion and require controled atmospheric emissions in case they are submitted to energetic recycling after disposal. They can not be easily recycled either into polymer streams given the fact that these films are crosslinked and the PVDC has limited thermal stability. In regards to sustainability, the impact of packaging to the environment can be minimized by following some criteria such as: (i) it has to be beneficial, safe and healthy for individuals and communities throughout its lifecycle; (ii) meet the designed performance and cost; (iii) maximize the use of renewable or recycled materials; (iv) manufactured using clean production technologies and best practices; (v) made from materials healthy in all probable end-of-life
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Rodrigues, J. B., Brunelli, K., Sarantopoulos, C. I. G. L., & Oliveira, L. M. scenarios and; (vi) can be recovered effectively and used in biological and/or industrial cradle-to-cradle cycles[7]. To mitigate the environmental problems associated with PVDC films and to offer beef processors a more sustainable packaging alternative, multilayer coextruded films containing EVOH, PA and PE have been tested to package fresh meat products and are regarded as a valid alternative to traditional PVDC films used in wholesale distribution of chilled meat[8]. Multilayer barrier shrinkable films made with polyethylene terephthalate (PET), ionomeric ethylene copolymers, ethylene-vinyl alcohol copolymer (EVOH), polyamide (PA) and polyethylene (PE) are chlorine free and do not require radiation crosslinking to become thermally shrinkable. These features are of particular interest to recycling plants as non‑crosslinked materials can be more easily merged into regular recycling streams and also to film converters that do not need to operate gamma radiation units and therefore avoid high energy radiation risks in the working environment. Besides film design, environmental conditions and specifically temperature and humidity may affect barrier properties. A possible drawback of films containing EVOH and PA is that these materials are hygroscopic and regarded to be sensitive to moisture in a way that when exposed to high moisture environment they may have limited ability to provide adequate oygen barrier[9]. According to McKeen [9] the permeability of EVOH and PA films is depending on the relative humidity. The oxygen permeability coefficient of a 15 µm thick film made of EVOH with 32 mol % of ethylene varies from 0.01 to 0.05 mL.mm.m-2.day-1.atm-1 at 20 oC as the relative humidity (RH) varies from 0% to 85%, while for 20 µm thick films made of EVOH with 44 mol% of ethylene, the oxygen permeability coefficient varies from 0.04 to 0.08 mL.mm.m-2.day-1.atm-1 at 20 oC as RH increases from 65% to 85%. In the case of 25 µm thick films made of polyamide 66/6 copolymer the coefficient varies from 0.94 to 5.91 mL.mm.m-2.day-1.atm-1 at 23 oC as the RH varies from 0% to 90%. It has been reported that multilayer PET/PVDC/PE films with layer thicknesses of 12/4/50 μm respectively have oxygen transmission rate as low as 5 mL.m-2.day-1.atm-1 (23 oC, 50% RH) and water vapor barrier of 2 g.m-2.day-1 (23 oC, 85% RH), whereas a PET/EVOH/PE film (12/5/50 μm) has oxygen and water vapor transmission rates of 1 and 2 to 4 in similar units and testing conditions[10]. Therefore, and to optimize the oxygen barrier of multilayer films containing EVOH and/or PA, the structure must contain other outer layers, typically polyolefins, that can minimize water vapor transmission and can protect the inner EVOH and PA layers so the film as a whole can perform as an effective gas barrier. The ethylene content in the EVOH copolymer also has influence on the oxygen permeability of the polymer[9]. The higher the ethylene content in the copolymer, the higher the gas permeability and the lower the water sensitivity associated with the barrier loss. The permeability of EVAL F Series (EVOH 32 mol% of ethylene) is 0.4 mL.20µm.m-2. day-1.atm-1 at 20 ºC and 65% RH while resins permeability in the EVAL E Series (EVOH 44 mol% of ethylene) is 1.5 in the same units and conditions[11]. 126 126/130
Differently than polyamide 66/6 or EVOH resins, amorphous polyamides have excellent oxygen, carbon dioxide and water vapor barrier, even at extremely wet conditions such as 95% to 100% RH. The oxygen permeability coefficient of amorphous PA is reduced at higher humidity, and has values of 1.50 and 0.59 mL.mm.m-2.day-1.atm-1 at RH of 0% and 95% respectively (both at 30oC). Blending amorphous PA (aPA) into PA 6 or PA 66/6 polymers results in a product that behaves like an amorphous polymer with enhanced barrier properties at high humidity as well as toughness, strength and flexibility[12]. Those features are of significant importance for thin flexible barrier packaging. Foreseeing an increased interest from fresh beef producers to adopt more environmentally friendly packaging and given the sensitivity of EVOH and PA to high humidity, four different nine layer shrink-barrier film structures were developed and manufactured in a commercial scale triple bubble film co-extrusion line. The objective of this study was to compare the gas barrier and mechanical properties and shrink performance of films made with PET, ionomer, EVOH, PA and polyethylene with those of commercially available EVA/PVDC/PE films used to preserve the quality of fresh beef in the Brazilian market.
2. Materials and Methods 2.1 Packaging materials and structure Four different high barrier shrinkable bi-oriented tubular films containing one PET outer layer; one ionomeric ethylene copolymer layer; one EVOH layer; one or two layers of PA 66/6 blended with amorphous PA; one or two layers of PE blends made by blending pellets of Ziegler-Natta and metallocene catalysts polyethlenes and two layers of polyethylene grafted with maleic anhydride resins used as co-extrusion adhesive. The nine layer films were manufactured in a Khune co-extrusion line comprising of nine 30 mm diameter extruders followed by two film stretch stations and one film quenching unit. A five layer film, made with EVA (outerlayer), PVDC and PE (selant layer) reticulated with gamma-ray, commonly used in the Brazilian market was used as the “control” to which the four competing alternatives were compared. Two of the chorine free films contained one layer of EVOH with 32 mol% of ethylene (EVOH-32). In one case the EVOH layer was encapsulated between a layer of modified PA and a PE layer, while in the other case the EVOH layer was encapsulated in-between two modified PA layers. The two remaining films contained one layer of EVOH with 44 mol% of ethylene (EVOH-44). Like in the EVOH-32 case, the EVOH-44 layer was encapsulated between a layer modified PA layer and a PE layer in one case and in-between two modified PA layers in the other. In all the chlorine free film structures just described, the modified PA layers contained PA 66/6 and aPA (DuPont™ Selar PA) blends. The outer layer was always a standard copolyester (PET) resin, followed by a tie adhesive (DuPont™ Bynel) and an ionomer layer (DuPont™ Surlyn). The sealing layer was always a blend of polyethylene resins. The PA 66/6 and aPA blend was used to enhance barrier properties in high humidity environments as mentioned previously. Polímeros, 28(2), 125-130, 2018
Properties of barrier shrink bags made with EVOH and polyamide for fresh beef meat preservation The films structure of the five samples compared are summarized as follows: - Control: EVA/tie/PVDC/tie/LLDPE; - EVOH32-1: PET/tie/Ionomer/tie/PA6/66+aPA/EVOH 32 mol%/tie/LMDPE/LLDPE (LMDPE: linear medium density polyethylene/ LLDPE- linear low density polyethylene); - EVOH32-2: PET/tie/Ionomer/tie/ PA6/66+aPA/EVOH 32 mol%/ PA6/66+aPA/tie/LLDPE; - EVOH44-1: PET/tie/Ionomer/tie/ PA6/66+aPA/EVOH 44 mol%/tie/LMDPE/LLDPE; - EVOH44-2: PET/tie/Ionomer/tie/ PA6/66+aPA/EVOH 44 mol%/ PA6/66+aPA/tie/LLDPE.
Seal strength, puncture resistance, oxygen and water vapor transmission rates and film shrink were measured for the four films and compared to those obtained for a typical EVA/PVDC/PE film.
2.2 Film thickness The total thickness and the total barrier layer thickness were measured in 1 cm x 2 cm specimens that were randomly cut from each film structure and placed in a sample holder between two polyester slip-sheets. Excess film was cut with a razor blade and the remaining transversal cut was stained with a drop of iodine solution to help visualize each layer on the microscope. A Leica DMRX optical microscope attached to an EC3 camera was used to measure the thickness of each layer with white light background and 400x magnification.
2.3 Seal strength The maximum bottom heat seal strength was determined according to ASTM F 88/88M[13] standard procedure with an Instron universal testing machine model 5500R. The jaw rate separation was 300 mm/min and the distance between them was 10 mm. Ten 25.4 mm wide specimens per film sample were tested. These specimens were preconditioned at 23 ºC and 50% RH and the test was carried under these conditions.
2.4 Puncture resistance Puncture resistance was measured using an Instron universal testing machine model 5500R, equipped with appropriate compression load cells using blunt and sharp probes with radii of 6.35 mm and 0.79 mm respectively. Circular specimens of 95.25 ± 0.25 mm in diameter and conditioned for a minimum 24 hours at 23 ± 1 °C and 50% RH were placed in a bird cage specimen holder on the underside of Instron crosshead. The probe speed was 51 mm/min. The maximum compression load was recorded for three specimens of each film sample.
2.5 Water vapour transmission rate (WVTR) WVTR was determined by a gravimetric method according to ASTM E 96/E 96[14]. This standard procedure is based on the weight gain of anhydrous calcium chloride placed inside an aluminum capsule that is isolated from Polímeros, 28(2), 125-130, 2018
room atmosphere by the specimen. The effective permeation area for each specimen was 50 cm2. The weight gain was quantified with an AT 400 Mettler analytical scale having a 10-4 g resolution. The test was made in a Vötsch – VC 0057 chamber at 38.0 ± 0.1 ºC and 90.0 ± 0.5% RH. Five specimens of each film sample were tested.
2.6 Oxygen transmission rate (OTR) OTR was determined by coulometry method according to ASTM F1927[15] using a MOCON OXTRAN equipment model 2/20, operating with pure oxygen as permeating at 23 °C and 75% RH. Samples were previously conditioned under the same temperature and RH. The effective permeating area for each specimen was 50 cm2. Results for two specimens of each film sample were adjusted for 1 atm parcial pressure gradient of oxygen.
2.7 Film shrink The free linear thermal shrinkage of the films was determined according to ASTM D 2732[16] standard procedure. The initial test specimens dimension was 100 mm x 100 mm. Specimens were placed inside a hot water bath at 85.5 ± 0.5 ºC for 5 seconds. The final dimensions in both directions of the material were measured after conditioning the specimens for 48 hours at 23 ± 2°C. Five specimens were tested for each film.
3. Results and Discussions 3.1 Film thickness The total film thickness and the gas barrier layer thickness for all film structures are reported on Table 1. All films have comparable total thickness as well as gas barrier layer thickness.
3.2 Seal strength Fresh beef is typically packed under vacuum to remove as much oxygen as possible from the inside of the package. After it is sealed, a package must provide a hermetic closure to prevent oxygen ingress into the package allowing spoilage bacteria growth. Therefore meat packages must have adequate seal strength to allow for a tight closure of Table 1. Total film thickness and gas barrier layer thickness of EVOH and PVDC (control) multilayer shrinkable films. Treatment Control EVOH32-1 EVOH32-2 EVOH44-1 EVOH44-2
Average Total Film Thickness (µm) 59.2 69.3 67.9 56.5 62.1
Average Barrier Layer Barrier Layer Material Thickness (µm) PVDC 4.6 EVOH 4.6 EVOH 4.6 EVOH 3.5 EVOH 4.1
Control: EVA/tie/PVDC/tie/LLDPE; EVOH32-1: PET/tie/Ionomer/ tie/PA6/66+aPA/EVOH 32 mol%/tie/LMDPE/LLDPE; EVOH32-2: PET/tie/Ionomer/tie/PA6/66+aPA/EVOH 32 mol%/PA6/66+aPA/ tie/LLDPE; EVOH44-1: PET/tie/Ionomer/tie/PA6/66+aPA/EVOH 44 mol%/tie/LMDPE/LLDPE; EVOH44-2: PET/tie/Ionomer/tie/ PA6/66+aPA/EVOH 44 mol%/PA 6/66+aPA/tie/LLDPE.
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Rodrigues, J. B., Brunelli, K., Sarantopoulos, C. I. G. L., & Oliveira, L. M. packages. Seal strength results for all film samples varied from 1.5 to 2.6 kN/m (Figure 1) and no significant difference can be assigned to any of the mean values obtained for the five film samples.
result obtained for the control sample is smaller than the lowest value obtained for the films containing two modified PA layers (EVOH32-2 and EVOH44-2).
3.3 Puncture resistance
Although the WVTR is not a primary concern in fresh beef packages, water vapour barrier must be enough to protect the inner EVOH layer against moisture absorption to prevent oxygen permeation into the package. The results on Figure 3 show that all samples containing EVOH layers in their structures have significantly higher WVTR values than the control. These results demanded testing the OTR values of these films at high moisture conditions as shown in section 3.5.
Puncture resistance measures the ability of films to resist pinholes caused by rough handling or sharp objects, such as meat bones. Similarly to seal failure, pinholes must be avoided in order to maintain vacuum inside the package. The results in Figure 2 show that film samples containing EVOH have similar or better puncture resistance than the control EVA/PVDC/PE film. They also show that samples containing two modified PA layers (EVOH 32-2 and EVOH44‑2) have superior puncture resistance in comparison to the control film for either the blunt or sharp probes. In fact, the best
3.4 Water vapour transmission rate (WVTR)
3.5 Oxygen transmission rate (OTR) The OTR results at high humidity conditions (75% RH) compared in Figure 4 show that films containing EVOH 44% mol of ethylene (EVOH44-1 and EVOH44-2) have higher permeation rates compared to films containing EVOH 32% mol (EVOH32-1, EVOH32-2) and to the control samples. Additionally the EVOH32-2 film, containing two PA blend layers shows the best oxygen barrier results of the films tested suggesting that the amorphous PA in the PA layer protects the EVOH layer against moisture gain. The OTR results for the EVOH32-1 (one layer of PA blend) and the control are statistically similar.
3.6 Film shrink Figure 1. Bottom seal strength at yield of EVOH and PVDC (control) multilayer shrinkable films. Control: EVA/tie/PVDC/ tie/LLDPE. EVOH32-1: PET/tie/Ionomer/tie/PA6/66+aPA/ EVOH 32 mol%/tie/LMDPE/LLDPE; EVOH32-2: PET/tie/ Ionomer/tie/PA6/66+aPA/EVOH 32 mol%/PA6/66+aPA/tie/ LLDPE; EVOH44‑1: PET/tie/Ionomer/tie/PA6/66+aPA/EVOH 44 mol%/tie/LMDPE/LLDPE; EVOH44-2: PET/tie/Ionomer/tie/ PA6/66+aPA/EVOH 44 mol%/PA6/66+aPA/tie/LLDPE.
Figure 2. Puncture resistance using sharp and blunt probes of EVOH and PVDC (control) multilayer shrinkable films. Control: EVA/tie/PVDC/tie/LLDPE; EVOH32-1: PET/tie/Ionomer/tie/ PA6/66+aPA/EVOH 32 mol%/tie/LMDPE/LLDPE; EVOH32-2: PET/tie/Ionomer/tie/PA6/66+aPA/EVOH 32 mol%/PA6/66+aPA/ tie/LLDPE; EVOH44-1: PET/tie/Ionomer/tie/PA6/66+aPA/EVOH 44 mol%/tie/LMDPE/LLDPE; EVOH44-2: PET/tie/Ionomer/tie/ PA6/66+aPA/EVOH 44 mol%/PA6/66+aPA/tie/LLDPE. 128 128/130
The shrink results (Figure 5) indicate that all films made with EVOH have balanced shrink ratios in the machine (MD) and transversal directions (TD). This indicates that for round or cubic shape beef cuts those films may provide a more homogeneous wrapping, retaining liquid inside the muscle tissue tough rendering better product retail display. On the other hand, the control sample shows higher shrinking in the TD than in the MD. In this case for cuts that are much longer than wider the package may not wrap
Figure 3. Water vapour transmission rate of EVOH and PVDC (control) multilayer shrinkable films. Control: EVA/tie/PVDC/ tie/LLDPE; EVOH32-1: PET/tie/Ionomer/tie/PA 6/66+aPA/ EVOH 32 mol%/tie/LMDPE/LLDPE; EVOH32-2: PET/tie/ Ionomer/tie/PA6/66+aPA/EVOH 32 mol%/PA6/66+aPA/tie/ LLDPE; EVOH44‑1: PET/tie/Ionomer/tie/PA6/66+aPA/EVOH 44 mol%/tie/LMDPE/LLDPE; EVOH44-2: PET/tie/Ionomer/tie/ PA6/66+aPA/EVOH 44 mol%/PA6/66+aPA/tie/LLDPE. Polímeros, 28(2), 125-130, 2018
Properties of barrier shrink bags made with EVOH and polyamide for fresh beef meat preservation
Figure 4. Oxygen transmission rate at 23 ºC and 75% RH of EVOH and PVDC (control) multilayer shrinkable films. Control: EVA/tie/PVDC/tie/LLDPE; EVOH32-1: PET/tie/Ionomer/tie/ PA6/66+aPA/EVOH 32 mol%/tie/LMDPE/LLDPE; EVOH32-2: PET/tie/Ionomer/tie/PA6/66+aPA/EVOH 32 mol%/PA6/66+aPA/ tie/LLDPE; EVOH44-1: PET/tie/Ionomer/tie/PA6/66+aPA/EVOH 44 mol%/tie/LMDPE/LLDPE; EVOH44-2: PET/tie/Ionomer/tie/ PA6/66+aPA/EVOH 44 mol%/PA6/66+aPA/tie/LLDPE.
The combination of triple bubble blow film technology to produce these films and the selected resins resulted in more even TD and MD shrink ability of all the films containing EVOH without affecting significantly the seal strength results as compared to the control sample. In summary, the results obtained under controlled conditions indicate that the nine layer films containing EVOH 32% mol of ethylene, ionomer and PA 66/6+aPA blend in its composition, and especially the structure with EVOH 32% mol encapsulated between two layers of PA blend, have comparable or even slightly better performance features comparied to the control, in addition to be chlorine free and not requiring the radiation crosslinking used in the production of films made with EVA, PVDC and PE. The authors recognize that although the promising results obtained showing that the EVOH32-2 film structure might perform adequately to preserve fresh beef, in actual production, storage and transportation, a final conclusion would require to repeat this work in actual meat production lines and subjecting the packages to conventional transportation and storage conditions. Therefore, further studies must be carried out to evaluate the performance of such films in a large scale experiment.
5. References
Figure 5. Machine (MD) and Transverse Directions (TD) shrink of EVOH and PVDC (control) multilayer shrinkable films. Control: EVA/tie/PVDC/tie/LLDPE; EVOH32-1: PET/tie/Ionomer/tie/ PA6/66+aPA/EVOH 32 mol%/tie/LMDPE/LLDPE; EVOH32-2: PET/tie/Ionomer/tie/PA 6/66+aPA/EVOH 32 mol%/PA6/66+aPA/ tie/LLDPE; EVOH44-1: PET/tie/Ionomer/tie/PA6/66+aPA/EVOH 44 mol%/tie/LMDPE/LLDPE; EVOH44-2: PET/tie/Ionomer/tie/ PA 6/66+aPA/EVOH 44 mol%/PA6/66+aPA/tie/LLDPE.
the cut evenly allowing more fluid to exudate the beef tissue and rendering a loose appearance.
4. Conclusions We concluded that film samples containing EVOH 32% mol of ethylene are more effective to prevent oxygen permeation even under high moisture conditions than films made with EVOH 44% mol. Furthermore, the EVOH32-2, film structure with an EVOH layer encapsulated between two layers of PA 66/6+aPA blend, shows the lowest OTR values, supporting the positive roll played by the aPA in the blend to improve the gas and moisture barrier under high moisture. Samples containing two layers of PA 66/6+aPA blends (EVOH32-2 and EVOH44-2) offered the best puncture resistance, clearly showing the contribution of the PA blend in improving film toughness. Polímeros, 28(2), 125-130, 2018
1. Sarantopoulos, C. I. G. L., Rego, R. A., Dantas, T. B. H., Dantas, F. H., Jaime, S. B. M., Mourad, A. L., & Padula, M. (2012). As tendências de embalagem. In C. I. G. L. Sarantopoulos & R. A. Rego (Eds.), Brasil Pack Trends 2020 (pp. 67-83). Campinas: ITAL. 2. Zhou, G. H., Xu, X. L., & Liu, Y. (2010). Preservation technologies for fresh meat – a review. Meat Science, 86(1), 119-128. PMid:20605688. http://dx.doi.org/10.1016/j. meatsci.2010.04.033. 3. Robertson, G. L. (2013). Food packaging principles and practice. Boca Raton: CRC Press. 4. Scetar, M., Kurek, M., & Galic, K. (2010). Trends in meat and meat products packaging – a review. Croatian Journal of Food Science Technology, 2(1), 32-48. Retrieved in 2016, May 25, from http://hrcak.srce.hr/file/89744 5. Gazalli, H., Malik, A. H., Jalal, H., Afshan, S., Mir, A., & Ashraf, H. (2013). Packaging of meat. International Journal of Food Nutrition and Safety, 4(2), 70-80. Retrieved in 2016, May 25, from http://www.modernscientificpress.com/journals/ ViewArticle.aspx?6ZIT7oAL6Lqarm6Ljqm1ABuLMes5oQ LKKUOK5VwlHTsOPWlBdz6tl1E+5TyCVfuK 6. Morris, B. A. (2016). The science and technology of flexible packaging: multilayer films from resin and process to end use. Boston: Elsevier. 7. Lewis, H., Fitzpatrick, L., Verghese, K., Sonneveld K., Jordon, R. (2010). Sustainable packaging redefined: draft. Sustainable Packaging Alliance – SPA. 26 p. 8. Lee, K. T. (2010). Quality and safety aspects of meat products as affected by various physical manipulations of packaging materials. Meat Science, 86(1), 138-150. PMid:20510533. http://dx.doi.org/10.1016/j.meatsci.2010.04.035. 9. Mckeen, L. W. (2012). Permeability properties of plastics and elastomers. Kidlington: Elsevier Inc. 10. Lange, J., & Wyser, Y. (2003). Recent innovations in barrier technologies for plastic packaging —a review. Packaging Technology & Science, 16(4), 149-158. http://dx.doi.org/10.1002/ pts.621. 129/130 129
Rodrigues, J. B., Brunelli, K., Sarantopoulos, C. I. G. L., & Oliveira, L. M. 11. Kuraray. (2012). EVAL resins. Kuraray: Houston. 9 p. 12. DuPont™, Selar. (2005). PA3426 blends with nylon. Barueri: Dupont. 11 pp. 13. American Society for Testing and Materials – ASTM. (2009). ASTM F 88/F 88M: standard test method for seal strength of flexible barrier materials. West Conshohocken: ASTM. 14. American Society for Testing and Materials – ASTM. (2007). ASTM F 1927: standard test method for determination of oxygen gas transmission rate, permeability and permeance at controlled relative humidity through barrier materials using a coulometric detector. West Conshohocken: ASTM.
130 130/130
15. American Society for Testing and Materials – ASTM. (2015). ASTM E 96/E96M: standard test methods for water vapor transmission of materials. West Conshohocken: ASTM. 16. American Society for Testing and Materials – ASTM. (2012). ASTM D 2732: standard test method for unrestrained linear thermal shrinkage of plastic film and sheeting. West Conshohocken: ASTM. Received: May 25, 2016 Revised: Dec. 09, 2016 Accepted: Feb. 06, 2017
Polímeros, 28(2), 125-130, 2018
ISSN 1678-5169 (Online)
http://dx.doi.org/10.1590/0104-1428.01816
Potential doxorubicin delivery system based on magnetic gelatin microspheres crosslinked with sugars Josefa Souza1, Manoel Silva2 and Marcos Costa1* Laboratório de Química de Polímeros, Instituto de Química, Universidade do Estado do Rio de Janeiro – UERJ, Rio de Janeiro, RJ, Brazil 2 Laboratório de Física, Departamento de Física e Química, Universidade Federal de Itajubá – UNIFEI, Itajubá, MG, Brazil
1
*marcos.costa@uerj.br
Abstract The preparation and characterization of magnetic microspheres based on gelatin for use in drug delivery systems are reported. Sugars were employed as crosslinking agents and type A gelatin and type B gelatin were compared to prepare microspheres by water-in-oil emulsion. The influence of gelatin and sucrose concentration, temperature and stirring speed on microbeads’ characteristics was studied. The gelatin concentration and stirring speed were the parameters directly associated with the particle sizes. We found no relevant difference between the use of type A and type B gelatin. In addition, the gelatin crosslinking study revealed that sucrose is not a crosslinking agent but fructose can crosslink the protein chains when the reaction medium has pH 9. The size of the microspheres varied from 5 to 60 μm as measured by optical microscopic images. Doxorubicin adsorption and release were successfully performed using the microspheres crosslinked with fructose under the action of an external magnetic field. It was observed that the microspheres absorbed 69% of the doxorubicin that was in solution. After 24 h, about 45% of the DOX was displaced from microspheres to saline medium in the free form in the solution. Keywords: gelatin microspheres, magnetic properties, sugar crosslinking.
1. Introduction Gelatin is a mixture of water-soluble proteins obtained by hydrolysis of collagen from the skin, bones and connective tissues of animals[1]. There are two types of gelatins and they are characterized by their mode of manufacture. The Type A gelatin (pH 3.8-6.0; iso- electric point 6-8) is obtained from acidic hydrolysis of pork skin and the Type B gelatin (pH 5.0-7.4; isoelectric point 4.7-5.3) is obtained from basic hydrolysis of bones and animal skin[1]. Attributable to the excellent biocompatibility and biodegradability[2,3], gelatin has been widely used in biomedical materials for controlled drug release. In this application, can be found gelatin in different forms: films[4-6], disks[7], hydrogels[8,9], sponges[10] and frequently microspheres[7,9,11-14]. Microspheres are usually prepared by water-in-oil emulsion. However, the main preparation parameters vary widely in the literature. When microspheres are produced, these parameters can influence particle size and the microsphere’s size is very important to define the administration route[15,16] and the liberation rates[16]. Because of this, in this work, we designed experiments to determine the most important parameters that can influence particle size. Because gelatin is a water-soluble polymer, its must be modified for application in the human body (where the medium is aqueous). Thus, gelatin hydrogels can be prepared as three-dimensional hydrophilic networks that are able to release drugs at the controlled rates. Such networks can be physical as those obtained by gelatin mixed with other polymers such as sodium carboxymethyl cellulose[17], hydroxyethyl cellulose[18] and carboxymethyl
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guar gum[19] to form interpenetrating polymer networks (IPNs) or chemical as those obtained by using chemical crosslinking agents[2]. The chemical crosslinking agents are bifunctional or polyfunctional compounds that act by binding to carboxylic or amino groups of adjacent molecules of gelatin. Examples of this type of crosslinker include formaldehyde, glutaraldehyde, glyceraldehyde, imines, ketones, saccharides, dyes, calcium carbonate, carbodiimides, genipin and other bifunctional compounds[1]. There are many chemicals that can be used for gelatin crosslinking, but the crosslinking process in biomedical materials must be done with reagents that, like the polymer, are biocompatible and biodegradable. Most of these crosslinking agents can cause some cytotoxic effects because of unreacted fractions[7]. To avoid undesirable reactions, some studies have investigated the use of sugars as crosslinking agents[7,10,11,20]. Among the studied sugars, calls our attention the fact that the researchs conclude that sucrose is a crosslinking agent able to significantly reduce the gelatin water solubility. Additionally, sucrose is biocompatible, easy to obtain and inexpensive, making it a good candidate for use in controlled drug release. For these reasons, we will use it as a crosslinking agent in this work. Besides the main features of biocompatibility, biodegradability and low water solubility, the device designed by us should possess the ability to be transported inside the human body directly to target cells. To this end, in one of the phases of this study, magnetite is incorporated in the microspheres produced. Magnetite will confer
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O O O O O O O O O O O O O O O O
Souza, J., Silva, M., & Costa, M. magnetic properties to the device. Thus, it can be injected into the patient’s circulatory system and, with the aid of an external magnetic field, it is possible to concentrate the drug/biocomposite complexes at a specific target site in the body where the particles have entered the bloodstream[21]. Once the biocomposite is concentrated at the target, the drug can be released to act on the target cells by enzymatic activity or changes in pH, temperature or magnetic field. These magnetic drug delivery systems have many advantages over normal, non-targeted methods, such as: ability to target specific locations in the body; reduction of the drug quantity needed to attain a particular concentration in the vicinity of the target and reduction of the drug’s concentration at non-target sites, minimizing side effects[22]. The above described characteristics are essential for obtaining optimum system for controlled drug release. Despite its potential applications, few studies have investigated the use of sugars as a crosslinking agent to obtain gelatin microspheres and there is no literature about magnetic gelatin microspheres sugar crosslinked. Thus, in order to obtain a similar device, we decided to evaluate the effect of gelatin type, sucrose concentration, magnetite concentration and crosslinking time on the physical properties of the microspheres based on gelatin, sucrose and magnetite.
2. Materials and Methods 2.1 Materials Type B gelatin (225 bloom), type A gelatin (300 bloom), sucrose, fructose, corn oil and doxorubicin were purchased from Sigma-Aldrich Co. Acetone, and sodium hydroxide were acquired from B. Herzog Varejo de Produtos Químicos Ltda. Ferric chloride, sodium chloride and ferrous sulfate were purchased from Proquimios Comércio e Indústria Ltda. All chemicals were analytical grade and used as received.
2.2 Preparation of gelatin microspheres Microspheres were produced by thermal gelation. Briefly, 10 mL of 10% w/v gelatin solution preheated to 60 °C containing 40% w/w of sucrose was added dropwise to 40 mL of corn oil to form an emulsion by stirring with a two-paddle stirrer (1000 rpm). As the emulsion was obtained, the temperature was kept at 60 °C for different time periods and then lowered to 5 °C by rapid cooling in an ice bath. The microspheres formed were maintained in this condition for 30 minutes. Then, to completely solidify the droplets of the dispersed phase, 50 mL of precooled (5 °C) acetone was added and the mixture was stirred for another hour. The microspheres were filtered, washed with cool acetone (5 °C) and rapidly dried.
2.3 Preparation of magnetite
in the formation of a black colloidal magnetite solution. Subsequently, the dispersion was cooled to room temperature and was washed several times with distilled water until neutral pH. The magnetite formed was separated by magnetic decantation/separation and was dried in an oven at 60 °C for 24 h. The Fe3O4 nanoparticles’ precipitation happened according to the Equation 1 below: Fe2 + + 2Fe3+ + 8OH − → Fe3O 4 + 4H 2O (1)
2.4 Preparation of the magnetic gelatin microspheres Magnetic microspheres were produced by the same method described before. The magnetite was added in the gelatin solution and this mixture was then added dropwise to corn oil in order to form an emulsion.
2.5 Size particle distribution and average diameter These analyses were performed using a method described by Allen[24]. In this method, optical microscopic images were used to measure the diameter of 625 microspheres of each sample.
2.6 Magnetic properties The magnetic properties (saturation magnetization, residual magnetization and coercivity) were analyzed by using a Lake Shore series 7400 vibrating sample magnetometer (VSM).
2.7 DSC and FTIR analysis The thermal properties of the gelatin microspheres were analyzed by using a Perkin-Elmer Pyris 1 differential scanning calorimeter. The melting temperature (Tm) of the microspheres was determined under nitrogen atmosphere. Samples were scanned in aluminum pans, under static air atmosphere, at a heating rate of 20 °C/min in the temperature range of 50-200 °C. FTIR spectra of microspheres were measured by the KBr pellet method using a Perkin Elmer Spectrum One spectrophotometer.
2.8 Morphological analysis The gelatin microspheres’ morphology was determined by observation of the samples with a FEI Inspect 550 scanning electron microscope. The samples were coated with gold in an argon atmosphere for 120 s and the images were captured using acceleration voltages of 5 kV and 20 kV.
2.9 Swelling ratio Gelatin microspheres in the dry state were put on filter paper and weighed. Then the microspheres were immersed in distilled water at room temperature. Subsequently, the weight of the swollen microspheres was determined after 60 minutes. The swelling ratio (Rsw) of each test sample was calculated as follows (Equation 2):
Magnetite nanoparticles were synthesized using an adaptation of a previously described co-precipitation method[23]. This involved adding 100 mL of an aqueous solution of sodium hydroxide (concentration of 10 mols/L) dropwise to a mixture of iron salts with Fe2+/Fe3+ molar ratio = R sw ( ( Ws – Wd ) / Wd ) ×100 (2) of 1/2, forming an immediate dark brown/black solution. where Ws denotes the weight of the test sample after swelling The solution was stirred for 1 h at room temperature and and Wd is its initial weight in the dry state. then was heated at 90 °C for another 1 h, which resulted 132 132/138
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Potential doxorubicin delivery system based on magnetic gelatin microspheres crosslinked with sugars
Table 1. Experimental conditions in the 23 factorial design to evaluate their effects on particle size of gelatin microspheres.
Total amount of DOX − Free amount of DOX ×100 (3) Total amount of DOX
The in vitro DOX release experiments were carried out in phosphate buffer saline (PBS) (pH 7.4, 1.2 mM KH2PO4, 1.15 mM Na2HPO4, 2.7 mM KCl, 1.38 mM NaCl) in presence of constant magnetic field (6000 gauss) using a magnet. In order to determine the released amount of the DOX, 0.1 g of DOX-loaded magnetic gelatin microspheres was added to 8 mL of PBS (release medium, pH 7.4). The resulting suspension was gently shaken under a constant magnetic field of 6000 Gauss for predetermined time period. After shaking, 3 mL of supernatant was withdrawn and assayed for DOX spectrophotometrically (Fentom 600S at 480 nm). Each experiment was carried out in triplicate.
Stirring speed (rpm)
DL ( % )
The analysis of variance (ANOVA) was used to analyze the effect of gelatin concentration, sucrose concentration, stirring speed and temperature on gelatin microspheres’ particles size
Temperature (°C)
The doxorubicin (DOX) loading and in vitro DOX release were determined using gelatin magnetic microspheres made with 50% magnetite and fructose crosslinked at pH = 9. The loading of DOX was performed by allowing the magnetic gelatin microspheres (50 mg) to contact a freshly prepared DOX solution (200 ppm) for 1 hour. Then, the amount of free DOX in the solution was quantified by UV-Vis spectroscopy (Fentom 600S) at 480 nm. The DOX loading efficiency (DL)(%) was calculated using the following Equation 3:
Sucrose concentration (%)
2.11 Doxorubicin loading and in vitro doxorubicin release
The particle size distribution showed in Figure 1 revealed that the diameter of the microspheres produced in all experiments ranged from 5 to 60 µm. However, there was predominance in the range from 11 to 30 µm. For application in drug delivery systems, gelatin microspheres should have sizes below 5 µm for intravenous administration and should be smaller than 125 µm for arterial administration[15]. Thus, the particles obtained in all experiments were adequate for use in drug delivery by the arterial route.
Gelatin concentration (%)
The iron concentration of the gelatin microspheres was determined by atomic absorption spectroscopy. About 10 mg of each sample was heated in a flat-bottomed flask with 20 mL of aqua regia at reflux temperature for 24 h. Then the solution was cooled to room temperature, filtered into a 100 mL volumetric flask and the volume was completed with distilled water. The solution was analyzed by a Perkin Elmer Analyst 300 spectrometer.
(30 minutes) and acetone cooling time (1 hour). In order to replace cytotoxic crosslinkers, sucrose was chosen as crosslinking agent because it is a well-known biocompatible reagent.
Sample
2.10 Atomic absorption spectroscopy
P1 P2 P3 P4 P5 P6 P7 P8
10 20 10 20 10 20 10 20
0 0 40 40 0 0 40 40
40 60 60 40 60 40 40 60
500 500 500 500 1000 1000 1000 1000
3. Results and Discussions 3.1 Influence of gelatin and sucrose concentration, temperature and stirring speed on particle size of materials obtained A full factorial design at two levels, 23, was applied to evaluate the main effects. The variables considered and the levels studied are shown in Table 1. The experiments involved fixing the oil phase (corn oil), aqueous phase/oil phase ratio (1/4), heating time (10 minutes), cooling time
Figure 1. Particles size distribution of gelatin microspheres obtained according to Table 1.
Table 2. Analysis of the effects of variables on particles size by ANOVA. SS MS F p Gelatin concentration (1) 229.523 229.523 44.534 0.00016 Sucrose concentration (2) 22.9441 22.9441 4.45182 0.06788 Stirring speed (3) 59.213 59.213 11.489 0.00951 Temperature (4) 0.714 0.714 0.13854 0.7194 1*2 1.2544 1.2544 0.24339 0.63503 1*3 0.093 0.093 0.01805 0.89645 1*4 5.5932 5.5932 1.08525 0.32798 SS = square sum; MS = mean square; F = F-test; p = significance level; 1*2 = gelatin concentration and sucrose concentration interaction; 1*3 = gelatin concentration and stirring speed interaction; 1*4 = sucrose concentration and temperature interaction.
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Souza, J., Silva, M., & Costa, M. As can be seen in Table 2, with a 95% confidence level, the parameters that are directly associated with the particle sizes are the gelatin concentration and stirring speed. The data evaluation shows that smaller diameters are obtained with solutions of low gelatin concentration and higher stirring speeds. Because of this, the microspheres made afterward were prepared with a 10% gelatin solution and 1000 rpm stirring speed.
3.2 Gelatin type influence Manufacturers offer a wide variety of gelatins that are simple combinations of type A and type B gelatins. Thus, it is important to know if the gelatin type will influence the properties of the final particles. In this study, microspheres were prepared with both gelatin types and their magnetic and solubility properties were compared. The magnetic properties are extremely important for application in the device proposed in this paper. The saturation magnetization of the particles should be known in order to calculate the magnetic field strength that must be applied externally. Another very important characteristic is superparamagnetism. For application in the human body, this property is essential to prevent the particles’ agglomeration, which can lead to clogging of blood vessels. For this reason, we assessed the effects of the magnetite concentration on the magnetic properties of the microspheres. Table 3 shows the magnetic properties obtained. As might be expected, the saturation magnetization increased with rising magnetite concentration. Higher magnetization was observed for the microspheres obtained with type A gelatin when 50% was added during preparation, but this behavior was not observed for the other magnetite concentrations. All microspheres prepared in this experimental series had superparamagnetic behavior because remnant magnetization close to zero was observed. The data show that the experimental values are all larger than the theoretical values. The most probable hypothesis for these experimental results is that magnetite in the pure state forms clusters but is evenly dispersed when placed in a gelatin matrix. The formation of clusters tends to decrease the saturation magnetization of the particles while homogeneous dispersion has the opposite effect[25]. FTIR experiments were performed to find evidence of sugar-mediated crosslinking. Figure 2a shows the spectra of raw materials used to produce gelatin microspheres
(type A gelatin, type B gelatin and sucrose). According to Cortesi et al.[7], the absorption band located at 1450 cm-1 is characteristic of an aldimine stretching vibration, which provides evidence of the crosslinking of gelatin. However, this band is already present in the microspheres’ raw material. Figure 2b shows the FTIR spectra of type A gelatin and type B gelatin microspheres with no sucrose and with 40% (w/w) of sucrose. All microspheres showed the same peaks with similar intensities in the infrared region and no difference between type A and Type B gelatin was noted.
Figure 2. FTIR spectra (a) raw materials (type B gelatin, type A gelatin and sucrose) and (b) gelatin microspheres.
Table 3. Influence of gelatin type and magnetite concentration on magnetic properties of materials. Sample
Gelatin type
Magnetite GAM10 GAM20 GAM50 GBM10 GBM20 GBM50
A A A B B B
Magnetite Iron Magnetite concentration1 (%) concentration2 (%) concentration3 (%) 10 4.34 5.99 20 6.32 8.73 50 12.2 16.86 10 3.65 5.05 20 5.57 7.71 50 10.4 14.38
Theoretical Ms (emu/g) 2.54 3.7 7.14 2.14 3.26 6.09
Experimental Ms (emu/g) 42.33 3.27 4.42 9.24 3.64 4.80 8.18
Magnetite concentration usen on microsphere fabrication; 2Iron concentration analysed by atomic absorption analysis; 3Magnetite concentration calculated by atomic absorption analysis. 1
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Potential doxorubicin delivery system based on magnetic gelatin microspheres crosslinked with sugars Therefore, these spectra present no evidence of sucrose crosslinked gelatin. In addition, water solubility tests were conducted to verify how the solubility was affected by addition of sugar. All the samples dissolved completely in under 3 hours. These results indicate that the microspheres have a low level of crosslinking. Type A and type B gelatin had similar properties with respect to solubility in water and minor differences in relation to magnetic properties. Thus, we decided to use only type B gelatin to produce microspheres from this point on.
3.3 Gelatin microsphere crosslink study In order to optimize the efficiency of gelatin microspheres as drug carriers, their number of crosslinks in the polymer matrix should be evaluated. Because of this, we decided to observe how the heating time influences protein crosslinking in a series of experiments using type B gelatin. The obtained microspheres were characterized by differential scanning calorimetry (DSC), swelling analysis and scanning electronic micrography. The microspheres’ composition, as well as swelling and DSC results are in Table 4. As can be seen in Table 4, the melting point of gelatin microspheres increased when the heating time increased. The only difference between sample GB and GBt10 is that in the last one, sucrose was added during the preparation. The melting points of samples GB (no sucrose) and GBt10 (40% sucrose) did not differ greatly and the sample without sucrose showed a slightly higher melting point than that in the presence of sugar. We carried out swelling tests because they are a relatively easy way to measure the ability of gelatin microspheres to retain water. The results showed a significant decrease in the water retention with longer heating time. The values of Tm and swelling ratio corroborate each other, confirming that longer heating time increases the number of crosslinks in the polymer matrix. In contrast, the presence or absence of sucrose in the microspheres had little influence on the data analyzed, which leads us to believe that sugar had little influence on the crosslinks formed. Figure 3 shows scanning electron micrographs of microspheres of gelatin obtained at different heating times. As can be seen in Figure 3 (left side), no significant differences were observed. The particles present spherical morphology, but there are many agglomerates. Probably, these agglomerates are formed because it was not added a surfactant agent during emulsion preparation. Figure 3 (right side) shows the
difference on particles’ surface according the heating times. It can observed that the surfaces became smoother when the heating time increased. With a higher magnification, this difference can clearly be seen when comparing the heating times of 10 and 2880 min (Figure 4). Based on these results, there are two possible explanations for these experimental observations. The first one is based on Russo’s[26] paper. According to him, crosslinks can occur by intermolecular bonds (interstrand), which occur between arginine-lysine or arginine-arginine within the same strand, while amino acid residues from two neighboring strands can also interact and form intramolecular (intrastrand) crosslinked strands, providing strength to the gelatin. The second theoretical explanation is based on carbohydrate chemistry. Sugars commonly exist as cyclic molecules because alcohols react reversibly with aldehydes and ketones to give hemiacetals and hemicetals, respectively. However, in the equilibrium state, there is a mixture of carbohydrate isomers and a small fraction of aldehyde or ketone source. Although small, the fractions of aldehyde and ketone allow the occurrence of common reactions of these organic functions[27]. Sucrose is a disaccharide composed of one glucose and one fructose molecule, both reducing sugars. The link between the two monosaccharides (glycosidic bond) forming disaccharide prevents the opening of the cyclic-form portions of fructose and glucose, resulting in the absence of aldehydic and ketonic forms in equilibrium, so the common reactions of these functions do not occur[27]. Sucrose is liable only if there is a hydrolysis reaction of the molecule to form the start of monosaccharides, which can only happen with a strongly acidic medium or under the influence of catalysts or enzymes. Since the reaction medium for preparation of gelatin microspheres here did not provide the main conditions for hydrolysis of sucrose, the gelatin remained in its original form, i.e., unable to form crosslinking reactions. Considering the theoretical foundations presented and the results of thermal analysis and swelling, we assume that sucrose does not react with the gelatin chains, so the increase in melting point of the microspheres was only due to the crosslinks formed by intermolecular and intramolecular bonds, which were favored by increasing the heating time. Although some authors[7,11,20] have indicated the use of sucrose as a biocompatible and biodegradable alternative to crosslink gelatin, we found no evidence of chemical reaction between the gelatin amino groups and sucrose. Thus, we decided to test fructose as crosslinking agent because the ketone functional group of fructose is more reactive than the aldehyde functional group of glucose. In order
Table 4. Melting temperature and swelling ratio of the gelatin microspheres as a function of heating time and sucrose concentration. Sample Sucrose Type B gelatin GB GBt10 GBt30 GBt1440 GBt2880
Heating time (min) 10 10 30 1440 2880
Sucrose concentration (%) 0 40 40 40 40
Tm (°C) 191 161 166 165 175 197
Rsw 570 443 441 365 225
Tm = melting temperature; Rsw = swelling ratio.
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Figure 3. Scanning electronic micrographs of the samples GB (a), GBt10 (b) and GBt2880 (c) with 1,000X (1) and 15,000X (2).
Figure 4. Scanning electronic micrographs of the samples GB (a) and GBt2880 (b) with 20,000X. 136 136/138
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Potential doxorubicin delivery system based on magnetic gelatin microspheres crosslinked with sugars to compare this substance with a traditional crosslinking agent, microspheres crosslinked with glutaraldehyde were also prepared. Preliminary solubility tests were performed comparing gelatin microspheres crosslinked with fructose (GBF sample) with gelatin microspheres crosslinked with glutaraldehyde (GBG sample). In these tests, a few milligrams of sample were left in contact with water for 24 h. The GBG sample was insoluble while the GBF was soluble. This result showed that the crosslinking of the protein chains’ gelatin using sugar as crosslinking agent does not occur as easily under normal conditions, so more factors should be investigated. Therefore, we decided to modify the pH in order to change the equilibrium between cyclic and open fructose forms and thus provide more ketone available for the formation of crosslinks. According to the literature, the kinetics of bond formation in chemical crosslinking of gelatin solutions is strongly affected by the solution’s pH[28], but at pH values higher than 9 and lower than 5 the denaturation enthalpy decreases, indicating that the triple helix amount is reduced[29]. Thus, gelatin microspheres were produced by varying the pH. Preliminary solubility tests were performed and since the aim was to reduce solubility, the gelatin microspheres made with pH 9 solution (GBF9) were chosen. The thermogravimetric analysis revealed an initial degradation temperature (Tonset) of 288 °C for the GBG sample and 294 °C for the GBF9 sample. This analysis also showed that the GBF sample had a residual 10 percentage points higher than the GBG sample. Higher degradation temperatures indicate higher crosslinking degree because more energy is required to break chemical bonds. Likewise, a larger amount of residue confirms that the particle has more strongly linked protein chains. These results show that the fructose crosslinking was successful.
3.4 Preliminary drug release tests We observed that the microspheres absorbed 69% of the doxorubicin that was in solution. If this value is compared in the literature for drug absorption by gelatin microspheres[13,14,30], one can considerer that a satisfactory amount of the drug was incorporated into the gelatinous matrix. Figure 5 shows the results of in vitro DOX release tests. The saline solution mimics the biological environment because it has similar pH and osmotic pressure. In these conditions, the gelatinous support gradually increased DOX
release over the time. After 24 h, about 45% of the drug was displaced from microspheres to saline medium in the free form in the solution. In this way, these preliminary release tests show that the method described in this study can be successfully used for the magnetic gelatin microspheres obtainment to incorporation and controlled release of doxorubicin.
4. Conclusions With the aim of obtaining gelatin microspheres with suitable properties for use in drug delivery systems, we evaluated the experimental parameters using a set of experiments. The statistical results showed that smaller particles can be prepared when low gelatin concentration and high stirring speed are used. By applying these parameters we obtained microspheres with appropriate size to use in arterial drug delivery systems. Because of the large variety of types available in the market, we decided to investigate whether there are significant differences between the use of type A gelatin and type B gelatin. The analyses showed no difference between the two types regarding crosslinking or adsorption of magnetic material in the gelatinous matrix. Furthermore, superparamagnetic samples were obtained with both gelatin types. With respect to crosslinking of the protein chains, we analyzed whether use of sucrose is effective to make the beads more biocompatible. The microspheres obtained remained very soluble in aqueous media and, so sucrose is not a suitable sugar to crosslink gelatin. Nevertheless, the extent of crosslinking increased as a function of heating time periods. Because of this, we analyzed the use fructose in place of sucrose. Taken together the results obtained indicate that crosslinked gelatin microspheres can be prepared using fructose when the reaction pH is 9. The microspheres crosslinked with fructose were successfully used in preliminary tests of adsorption and release of doxorubicin (a drug that is widely used in the treatment of cancer patients). Thus, the material prepared in this paper has great potential for use in drug delivery systems.
5. Acknowledgements We gratefully acknowledge FAPERJ for a scholarship to J.V.S. Souza, and FAPERJ and CNPq for financial support.
6. References
Figure 5. Release of doxorubicin from gelatin microspheres crosslinked by fructose. Polímeros, 28(2), 131-138, 2018
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17. Rokhade, A. P., Agnihotri, S. A., Patil, S. A., Mallikarjuna, N. N., Kulkarni, P. V., & Aminabhavi, T. M. (2006). Semiinterpenetrating polymer network microspheres of gelatin and sodium carboxymethyl cellulose for controlled release of ketorolac tromethamine. Carbohydrate Polymers, 65(3), 243-252. http://dx.doi.org/10.1016/j.carbpol.2006.01.013. 18. Kajjari, P. B., Manjeshwar, L. S., & Aminabhavi, T. M. (2011). Semi-interpenetrating polymer network hydrogel blend microspheres of gelatin and hydroxyethyl cellulose for controlled release of Theophylline. Industrial & Engineering Chemistry Research, 50(13), 7833-7840. http://dx.doi. org/10.1021/ie200516k. 19. Phadke, K. V., Manjeshwar, L. S., & Aminabhavi, T. M. (2014). Biodegradable polymeric microspheres of gelatin and carboxymethyl guar gum for controlled release of theophylline. Polymer Bulletin, 71(7), 1625-1643. http://dx.doi.org/10.1007/ s00289-014-1145-y. 20. Schuler, B. J. (2004). Evaluation of novel cross-linking agents for gelatin/collagen matrices (Doctoral thesis). School of Pharmacy, West Virginia University, West Virginia, USA. 21. Pankhurst, Q. A., Connolly, J., Jones, S. K., & Dobson, J. (2003). Applications of magnetic nanoparticles in biomedicine. Journal of Physics D: Applied Physics, 36(13), R167-R181. http://dx.doi.org/10.1088/0022-3727/36/13/201. 22. Arruebo, M., Fernández-Pacheco, R., Ibarra, M. R., & Santamaría, J. (2007). Magnetic nanoparticles for drug delivery. Nano Today, 2(3), 22-32. http://dx.doi.org/10.1016/ S1748-0132(07)70084-1. 23. Amali, A. J., & Rana, R. K. (2009). Stabilisation of Pd(0) on surface functionalised Fe3O4 nanoparticles: magnetically recoverable and stable recyclable catalyst for hydrogenation and Suzuki–Miyaura reactions. Green Chemistry, 11(11), 1781-1786. http://dx.doi.org/10.1039/b916261p. 24. Allen, T. (1997). Particle size measurement: powder sampling and particle size measurement. 5th ed. London: Chapman & Hall. 25. Santa-Maria, L. C., Costa, M. A. S., Hui, W. S., Santos, F. A. M., & Silva, M. R. (2006). Preparation and characterization of polymer metal composite microspheres. Materials Letters, 60(2), 270-273. http://dx.doi.org/10.1016/j.matlet.2005.08.033. 26. Russo, P. S. (1987). A perspective on reversible gels and related systems (ACS Symposium Series, Vol. 350, pp. 1-21). Washington: American Chemistry Society Symposium. 27. Allinger, N. L., Cava, M. P., Jongh, D. C., Johnson, C. R., Lebel, N. A., & Stevens, C. L. (1976). Organic chemistry. 2nd ed. New York: Worth Publishers. 28. Abete, T., Del Gado, E., Arcangelis, L., Serughetti, D. H., & Djabourov, M. (2008). Re-entrant phase diagram and pH effects in cross-linked gelatin gels. The Journal of Chemical Physics, 129(13), 134902. PMid:19045122. http://dx.doi. org/10.1063/1.2985655. 29. Gioffrè, M., Torricelli, P., Panzavolta, S., Rubini, K., & Bigi, A. (2012). Role of pH on stability and mechanical properties of gelatin films. Journal of Bioactive and Compatible Polymers, 27(1), 67-77. http://dx.doi.org/10.1177/0883911511431484. 30. Gaihre, B., Khil, M. S., Lee, D. R., & Kim, H. Y. (2009). Gelatincoated magnetic iron oxide nanoparticles as carrier system: drug loading and in vitro drug release study. International Journal of Pharmaceutics, 365(1-2), 180-189. PMid:18790029. http://dx.doi.org/10.1016/j.ijpharm.2008.08.020. Received: Mar. 11, 2016 Revised: May 17, 2017 Accepted: May 22, 2017 Polímeros, 28(2), 131-138, 2018
ISSN 1678-5169 (Online)
http://dx.doi.org/10.1590/0104-1428.11316
Ultrasound-assisted synthesis of polyacrylamide-grafted sodium alginate and its application in dye removal José Manoel Couto da Feira1*, Jalma Maria Klein1 and Maria Madalena de Camargo Forte1 Laboratório de Materiais Poliméricos, Escola de Engenharia, Departamento de Materiais, Universidade Federal do Rio Grande do Sul – UFRGS, Porto Alegre, RS, Brasil
1
*jose.feira@ufrgs.br
Abstract A polymeric adsorbent based on sodium alginate (SAG) grafted with polyacrylamide (PAM) (SAG-g-PAM) was synthesized using an ultrasound-assisted method. The addition polymerization was carried out with ammonium persulfate as the initiator, at different acrylamide (AM) concentrations. The SAG-g-PAM copolymers were evaluated by FTIR and 13C NMR spectroscopies, thermogravimetric analysis, grafting efficiency (%GE) and intrinsic viscosity in NaCl solution at 25 °C. Graft copolymers could be obtained in reaction lasting until 10 min by using ultrasound energy with grafting efficiency above 75%. The decolorization efficiency and adsorption capacity of the SAG-g-PAM copolymers were investigated in the adsorption of methylene blue (MB). The dye adsorption was pH dependent, and adsorption capacity (69.13 mg/g) maxima was at pH 10. All the graft copolymers have shown the same decolorization efficiency (99%), and the best one for MB removing is the SAG-g-PAM6 (%GE = 75%), since lower acrylamide content is required in the synthesis. Keywords: graft copolymer, alginate, acrylamide, ultrasound, adsorbents, methylene blue.
1. Introduction Polysaccharides derivatives have many potential applications such as hydrogel, tissue engineering, cell immobilization, food applications and have also been used extensively as adsorbents for contaminants removal from polluted water[1,2]. The solubility, hydrophobicity, physicochemical properties, and biological characteristics of the polysaccharides can be modified by reacting the hydroxyl and carboxyl groups with suitable compounds through derivatization of the functional groups[3-5], grafting[6,7], oxidative reactions[8], and/or hydrolytic degradation[9]. A convenient method to produce polysaccharide derivatives is the grafting of synthetic polymers onto the polysaccharide backbone[6,7,10,11]. Alginate grafted with acrylamide obtained by conventional, gamma rays and microwaves methods has frequently been used as polyelectrolyte in drug delivery[1,12,13], flocculant[14-16] and adsorbent in dyes removal[17]. Ultrasound-assisted method was used to produce polysaccharides modified with acrylamide, and it promoted shorter graft copolymerization time and higher grafting efficiency[18]. Ultrasound has been used to enhance the performance of a wide range of chemical reactions by cavitation process within a liquid, in which energy is introduced in a short period of time producing large number of microbubbles that collapse in few microseconds, providing extra energy to the system[19-24]. High-intensity ultrasound is also used to accelerate mass transport in mixing, drying, and extraction processes and other applications[25]. Environmental regulations are becoming stricter in what concerns the discharge and removal of dyes from aqueous effluents[26], since dyes are widely used in the textile, leather tanning, paper, plastic, food, cosmetic, and printing industries, for the coloration of their respective products[27].
Polímeros, 28(2), 139-146, 2018
The dyes are commonly synthetic and based on complex aromatic structures highly stable not biodegradable[28] that pollute water if not treated properly before discharge into the environment[29]. Methylene blue (MB), a common dye used in colouring cotton, silk, and wool, has adverse effects on human health such as eye irritations, breathing problems, diarrhea, etc. Not only photosynthetic activity but also aquatic biota[27,30,31] is affected by MB, and 15% of this is discharged into rivers by the textile industry[32]. It is well known that polyacrylamide homopolymer is used as an efficient adsorbent[2,7,11] or flocculant [7,14-16,18]. However, the high cost in addition the presence of unreacted acrylamide, which is toxic and carcinogenic, limits its application in water treatment. The new regulations and the growing of environmental awareness around the world have triggered the search for more environmentally friendly materials and processes. Low cost adsorbents based on agricultural wastes (industrial solid wastes and biomass versus clays minerals and zeolites) have being highly efficient, a renewable biomass to be exploited for MB remediation. It is estimated that over 7×105 tonnes of coloured wastewater are produced annually for the more than 100,000 commercially available dyes[33]. Different methods have been investigated for the removal of dyes from water and wastewaters including biological, physical (membrane filtration, adsorption, coagulation, flocculation, precipitation, reverse osmosis, ion exchange, etc.), and chemical (oxidation, ozonation, etc.) processes[32]. Adsorption is one of the most efficient methods to remove pollutants from effluents because of its simplicity of design, ease of operation, and insensitivity to toxic substances[34]. Acrylamide-grafted polysaccharides as a polymeric adsorbent or flocculant instead of polyacrylamide can be attributed
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O O O O O O O O O O O O O O O O
Feira, J. M. C., Klein, J. M., & Forte, M. M. C. to its biodegradability, less-toxic characteristic, abundance in nature, and its economic attractiveness. Alginate grafted with acrylamide has been used as adsorbents and flocculants in dye removal due to its superior performance compared with other grafted polymers[16,17]. In this paper, sodium alginate (SAG) modified with acrylamide (AM) was obtained by an ultrasound-assisted method aiming to get an improved route of synthesis, and evaluate the effect of the ultrasound energy on the grafting reaction and efficiency of the copolymers SAG-g-PAM. The graft copolymers were synthesized at different monomer concentrations, and evaluated as methylene blue adsorbent in a MB 2.2 X 10-6 mmolL−1 aqueous solution.
2. Materials and Methods 2.1 Materials SAG (viscosity, 15-20 cP) and AM (99%) were purchased from Sigma-Aldrich. Ammonium persulfate PA (APS, 98.6%) and MB were purchased from NEON. Glacial acetic acid PA (99.7%), and acetone PA (99.5%) were purchased from ACS. Formamide PA (99.5%) was purchased from Vetec Chemical Industries (Brazil). All the chemicals were used without further purification.
2.2 Synthesis of the graft copolymer by ultrasound-assisted method (SAG-g-PAM) The graft copolymers of sodium alginate and acrylamide were synthesized following previous work on modification of polysaccharide by using ultrasound carried out in our lab[18]. One gram of SAG (0.0062 mol of anhydroglucose unit) was dissolved in 40 mL of distilled water, specific concentrations of AM from 6 to 23 mol were diluted in 10 mL of water, and 1.75 x 10-3 mol of APS was dissolved in 5 mL of water. The three solutions were poured into a 100 mL three-neck round-bottom flask fitted with an ultrasound 13 mm stainless steel sonic wave emission probe, and the reactional mixture was sonicated using an ultrasonic generator (VCX 750, Sonics & Materials) operating at an output power of 750 W and 20 kHz frequency, under N2. The dissipation power of the probe was determined to be 19.6 W at 30% amplitude, using a calorimetric method reported by Margulis & Margulis[35]. The sonication has lasted between 7 and 10 min when the solution reached the gel point, and thus, the reaction time was dependent on the acrylamide concentration. In meanwhile, the reaction temperature increased from room temperature to 65±5 ºC since the reaction is exothermic and due to ultrasound energy. At the end of the reaction, the flask was cooled by immersing in ice water, and the gel was precipitated by adding 50 to 100 mL of acetone into the flask. The resulting precipitate, graft copolymer and polyacrylamide as sub-product, was filtered and washed with 50-100 mL of acetone, dried in an oven at 60 ºC, and pulverized. The pulverized product was purified with a formamide/acetic acid mixture (1:1 v/v) to remove the homopolymer[36], and then washed with acetone[37]. The grafting efficiency (%GE)[38] of the SAG-g-PAM copolymers was evaluated according to Equation (1): %GE
Wt. copolymer − Wt. polysaccharide ×100 (1) Wt. monomer
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2.3 Polymer characterization 2.3.1 Fourier transform infrared (FTIR) spectroscopy FTIR spectra of SAG, PAM, and SAG-g-PAM were recorded on a spectrometer (Perkin Elmer, Spectrum 1000) in the 4000 to 400 cm-1 region, using a KBr disk. 2.3.2 13C NMR spectroscopy The 13C NMR spectra of the SAG, AM, PAM, and SAG-g-PAM were recorded at 125 MHz in an NMR spectrometer (Agilent 500-MHz) with standard pulse programs in deutered water at 50 ºC. 2.3.3 Thermogravimetric analysis (TGA) The thermal stability of SAG, PAM, and SAG-g-PAM (~20 mg) was evaluated in a thermogravimetric analyzer TGA 2050 (TA Instrument), at a heating rate of 20 °C/min from 20 °C to 600 °C, under nitrogen flux (100 mL/min). 2.3.4 Intrinsic viscosity measurement Viscosity measurements of the polymer solutions were carried out with an Ubbelohde viscometer (viscosity constant: 0.004925 mm2/s2; capillary diameter: 0.46 mm) at 25±0.1 °C in a 0.1 M NaCl aqueous solution. The flow time of duplicate solutions was taken at four different concentrations (0.1, 0.05, 0.025, and 0.0125 g/dL). The relative viscosity (ηrel = t/to) was obtained using the flow time of the polymer solutions (t) and that of the solvent (to). Specific viscosity (ηsp = ηrel -1), reduced viscosity (ηred = ηsp/C), and inherent viscosity (ηinh = ln ηrel/C) (C = polymer concentration in g/dL) were mathematically calculated, and ηred and ηinh were simultaneously plotted against concentration. The intrinsic viscosity was obtained from the point of intersection after the extrapolation of the two plots (i.e., ηred versus C and ηinh versus C) to zero concentration.
2.4 Dye removal from aqueous media using copolymer adsorbent The removal of MB dye by adsorption on the SAG-g-PAM copolymer was performed in a batch system. In a 250 mL vessel filled with 50 mL of dye solution (2.2 X 10-6 mmol/L) was added 500 mg of SAG-g-PAM copolymer, and the mixture agitated at 130 rpm for 8 h at 25 ± 2 ºC. The mixture pH has varied from 4.0 to 10 by adding the required amounts of 1.0 M HCl or NaOH solutions. The adsorbed dye was separated from the leftover dye in solution by centrifugation at 6000 rpm for 40 min. Dye concentration in solution before and after centrifugation process was analyzed using a UV-visible spectrophotometer (PG Instruments, T80+), at a wavelength of 650 nm. A linear relationship between absorbance and dye concentration was determined at this wavelength and used as the calibration curve to determine the dye concentration at each equilibrium. Deionized water was used as reference and the percentage of decolorization efficiency (%DE) was calculated according to Equation (2): % = DE
C0 − C ×100 (2) C0
Where C0 and C are the concentrations of the dye solution, respectively, before and after treatment with the adsorbent Polímeros, 28(2), 139-146, 2018
Ultrasound-assisted synthesis of polyacrylamide-grafted sodium alginate and its application in dye removal polymer. The amount of dye adsorbed by mass unit of the copolymer (mg/g) was calculated according to Equation (3): q ( mg / g ) =
( C0 − C ) × V (3) m
where C0 and C are the concentrations of the dye solution (mg/L), respectively; before and after treatment with the adsorbent polymer; V is the volume of the aqueous phase (L); and m is the amount of dry adsorbent (g).
3. Results and Discussions The SAG-g-PAM copolymers were synthesized by changing the AM concentration and keeping constant the initiator (APS) concentration and the US power. The chemical modification of sodium alginate with acrylamide occurs at the hydroxyl groups of the D-glucopyranosyl unit, and grafting reaction mechanism of PAM into SAG backbone is depicted in Figure 1. According to Wang and Wang[39], the free radicals of the APS decomposition attack the D-glucopyranosyl hydroxyl groups and generated macro-radicals (SAG─O•) that attack the acrylamide producing PAM grafts into SAG backbone. Chain termination occurs as in a typical radical polymerization reaction. The ultrasound cavitation process also promotes the grafting reaction by generating extra energy and heat in the medium. We believe that the ultrasound energy favors instantly the production of both radicals, since the monomer incorporation occur in its totality at a very short period compared to conventional process reaction time. In the absence of initiator in the reaction medium the gel point was not reached, and thus no product. Table 1 shows the reaction yield, %GE, and intrinsic viscosity (η) of the SAG-g-PAM copolymers synthesized as a function of the
monomer (AM) concentrations. The intrinsic viscosity of the pristine alginate (SAG) and of the sonicated alginate with (SAG-US-APS) and without (SAG-US) the initiator was comparatively determined. The reaction yield, the grafting efficiency and intrinsic viscosity of the SAG-g-PAM copolymers increased as a function of the AM/SAG molar ratio. On the other hand, reaction time decreased with the monomer concentration’s increasing, since the gel point occurred faster due to the high acrylamide incorporation into alginate backbone, with consequent increasing of the polymer molecular weight. Under the reactional conditions, the reaction yield has changed from 83% to 99% and was of the same order of those reported using conventional method[18,40]. The great advantage of using ultrasound or sonication process is the time reduction, which was much short (< 10 min) in comparison to 24 h needed in the conventional grafting reaction process; in addition, same yield and grafting efficiency were obtained. As the AM/SAG molar ratio increased from 6 to 23 mol, %GE increased reaching almost 100%. A high molar ratio provides more monomer molecules per polysaccharide macro-radical in the reaction medium, leading to a high acrylamide incorporation in the polysaccharide backbone or higher %GE[41]. The graft copolymers showed intrinsic viscosity between 2.2 and 5.9 dL/g due to the acrylamide concentration, since the initiator concentration was constant. The alginate sonicated only with the initiator, or in the absence of the monomer (SAG-US-APS), undergoes chain cleavage and thus a molecular weight reduction, since the intrinsic viscosity decreased dramatically to 1.0 dL/g, if compared with viscosities of both alginate sonicated (SAG-US) and pristine polymer (SAG). The degradation of a polymer in a solution depends on the ultrasound power generated in the medium[24]. This is in agreement with works by
Figure 1. Grafting reaction mechanism of SAG D-glucopyranosyl units with acrylamide. Table 1. Reaction yield (%), grafting efficiency (GE) and intrinsic viscosity of the SAG-g-PAM copolymers as functions of acrylamide concentration*. Sample SAG-g-PAM6 SAG-g-PAM11 SAG-g-PAM23 SAG-US-APS SAG-US SAG
AM/SAG
Time
Yield**
GE
η
(m.r.) 6/1 11/1 23/1 -
(min) 10.1 9.4 7.5 12.4 14.2 -
(%) 83.0±0.4 83.0±0.8 99.0±0.1 -
(%) 75 79 98 -
(dL/g) 2.24±0.02 3.00±0.03 5.86±0.04 1.00±0.00 3.44±0.04 3.90±0.05
*APS amount: 1,75 mol x 10-3. **From duplicate reactions.
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Feira, J. M. C., Klein, J. M., & Forte, M. M. C. Iida et al.[42] and Hosseini et al.[43] that observed a decreasing of intrinsic viscosity of starch and sodium alginate solutions after sonication. The same intrinsic viscosity values for the alginate sonicated and non-sonicated shows that the ultrasound energy does not degrade the polysaccharide molecules under the conditions used in this work. On the other hand, the low viscosity of the alginate sonicated with the APS, shows that the initiator radicals attack the alginate molecules producing polysaccharide macro-radicals that in the absence of monomer molecules degrade easily other alginate macromolecules. Thus, the viscosity of the graft copolymer depends on the monomer concentration, and there is a synergic effect between the ultrasound energy and the initiator. The highest intrinsic viscosity of the SAG-g-PAM23 copolymer is because there is high concentration of PAM short chain branches into the alginate backbone, or instead it a low concentration of longer chain branches. All graft copolymers evaluated by FTIR, NMR and thermogravimetry have shown the same curve profiles. To follow the chemical modification in the graft copolymer, only the spectra and thermogram of the SAG-g-PAM11 copolymer, obtained with the medium acrylamide concentration, have been discussed compared with those of the neat polymers SAG and PAM.
The resonance peak at δ 177.9 ppm is assigned to the carbon of the ─COO- Na+ group, located at C-6. The overlapping resonance peaks in the range 60-85 ppm are assigned to the carbon atoms (C-2, C-3, C-4, and C-5) of the glucopyranosyl units. The 13C NMR spectrum of acrylamide (Figure 3b) has three distinct resonance peaks. The peak at δ 171.19 ppm is due to the amide carbonyl carbon (C=O), and the peaks at δ 128.01 and δ 129.03 ppm signs the two sp2 hybridized carbon atoms (i.e. CH2=CH). In 13C NMR spectrum of the polyacrylamide (Figure 3C), the peaks at δ 179.6, 35.1 and 42.1 ppm corresponding to the carbonyl group of its amide (CONH2) and sp3 hybridized carbon atoms, respectively. The 13C NMR spectrum of the SAG-g-PAM11 copolymer (Figure 3d) shows three additional resonance signals, compared to SAG. The stronger peak at δ 179.6 ppm is due to the amide carbonyl groups. The presence of the two additional peaks at δ 35.1 ppm and 42.1 ppm with sp3 hybridized carbon atoms, due to the presence of groups CH and CH2, respectively, confirm that the PAM chains (-[CH2-CH]n-) are grafted onto the SAG backbone. The chemical shifts shown in Figure 3 are comparable with those reported in the literature[45]. The peaks at δ 175, 166.9 and 20.9 ppm are due to solvents trace (formamide and acetic acid) used in the purification of the graft copolymer.
3.1 FTIR spectroscopy
3.3 Thermal stability of the graft copolymers
The FTIR spectra of the neat polymers SAG and PAM, and the SAG-g-PAM11 copolymer are shown comparatively in Figure 2. The SAG spectrum at wavenumber of 818 cm-1 shows the characteristic absorption bands of Na−O stretching, at 3442 cm-1 shows a broad peak due to the stretching frequency of the –OH group, at 1610 and 1416 cm-1 two absorption peaks relative to COO- group, and at 1031 cm-1 one sharp peak due to the C−O group[44]. Meanwhile, the PAM FTIR spectrum exhibites broad absorption bands at 3400 and 3180 cm-1 assigned to the stretching vibrations of N-H, at 1664 cm-1 band attributed to C=O stretching, and at 1612 cm-1 band of the N−H bending. In the SAG-g-PAM11 spectrum, all the absorption peaks mentioned above have shifted indicating that some interaction among the functional groups of sodium alginate and polyacrylamide has occured. Accordingly, the double N-H stretching vibration peak of polyacrylamide at 3400 cm-1 (Figure 2 PAM) and the OH stretching vibration peak of sodium alginate at 3442 cm-1 (Figure 2 SAG) is overlapped in SAG-g-PAM11 copolymer at 3410 cm-1. Some additional peaks in the copolymer spectrum are the peaks at 1682 and 1615 cm-1 that correspond, respectively, to the carbonyl amide and N-H stretching vibrations, and the peak at 1402 cm-1 due to C-N stretching vibrations. This last peak signes grafted PAM chains in the copolymer and confirms the intended grafting. This result is in good agreement with the results reported in the literature[45].
Figure 4 shows the mass loss (a) and derivative (b) TGA curves of the neat polymers SAG and PAM, and of the SAG-g-PAM11 copolymer. The sodium alginate decomposition (Figure 4a) showed three distinct stages of mass loss. The initial event at 25-125 ºC is due to the moisture desorption from the polysaccharide, with a mass loss of approximately 11%. The second and third events from 170 ºC to 300 ºC result of the decomposition of the carbohydrate backbone, followed by a mass loss of approximately 40%. The mass loss with maximum degradation rate at 221 ºC (Tmax) (Figure 4b) is due to groups COO- decomposition, with a subsequent decarboxylation of CO2. The mass loss at Tmax of 239 ºC (Figure 4b) is due to the SAG backbone decomposition[46] with a subsequent chain cyclization. Above 330 ºC the organic material undergoes carbonization remaining at 575 ºC a residue of 37%. The PAM degradation
3.2 13C NMR spectroscopy The 13C NMR spectra of SAG, AM, PAM, and SAG-g-PAM11 copolymer are shown in Figure 3. The 13C NMR spectrum of SAG (Figure 3a) shows three distinct resonance signals. The chemical shift (δ) at 102.3 ppm is referenced to the anomeric carbon atom (C-1) of the oxygen linkage. 142 142/146
Figure 2. FTIR spectra of (SAG), (PAM), and SAG-g-PAM11 copolymer. Polímeros, 28(2), 139-146, 2018
Ultrasound-assisted synthesis of polyacrylamide-grafted sodium alginate and its application in dye removal
Figure 3. (a) 13C NMR spectra of SAG; (b) AM; (c) PAM; (d) SAG-g-PAM11 copolymer.
shows two distinct stages of mass loss with Tmax at 300 ºC and at 411 ºC (Figure 4b). The continuous mass loss until 150 ºC is due to the moisture desorption from the sample. Above 230 ºC, the polyacrylamide first undergoes oxidation with 25% of mass loss due to ammonia by imidization (intra- and intermolecular) and water by dehydration. The second event with Tmax of 411 ºC is due to the chains cyclization process with mass loss approximately of 46%[47], remaining at 600 ºC a residue equivalent to 24%. The degradation profile of the SAG-g-PAM11 copolymer was lightly different of that presented by PAM, but also shows two mass loss event with Tmax at 280 ºC and at 398 ºC (Figure 4b). This little difference is consequence of the alginate chains modification by the PAM long chain branching. The mass loss from 200 ºC to 350 ºC is due to SAG decomposition and as well PAM oxidation. In the graft copolymer thermogram, there is an overlap of the mass loss of the SAG (200-270 ºC) with the first mass loss of the PAM (200-320 ºC), occurring an unique mass loss in the range of 200 ºC to 350 ºC. Before 200 ºC, the SAG-g-PAM11 copolymer presents better thermal resistance than SAG because it losses lower water content due to the functional groups chemical modification, and there is also a new type of chemical interaction between both polymers. According to the literature[48], hydrogel grafted with acrylamide, instead of acrylic acid, also presented better thermal stability, probably due to different types of covalent bonds in the grafting copolymer backbone. Polímeros, 28(2), 139-146, 2018
Figure 4. (a) Mass loss and (b) derivative TGA curves of the neat polymers (SAG) and (PAM), and SAG-g-PAM11 copolymer. 143/146 143
Feira, J. M. C., Klein, J. M., & Forte, M. M. C. 3.4 Decolorization efficiency of methylene blue dye solution Figure 5 shows the absorption capacity (q mg/g) of methylene blue dye, by the SAG-g-PAM11 copolymer, as a function of the pH, at an initial fixed concentration of dye 2.2 X 10-6 mmolL−1 and SAG-g-PAM11 of 500 mgL−1. The adsorption capacity of the SAG-g-PAM11 copolymer was higher as the pH of the methylene blue solution increased. The maximum absorption capacity was at pH 10 (69 mg/g), since the alkaline medium favors stronger electrostatic force of attraction between the anionic groups of graft copolymer and the polar groups of the dye MB. This result indicates that the adsorption process is controlled by a charge neutralization mechanism, where the cationic groups of MB are electrostatically attracted by the anionic carboxylate groups (COO- Na+) of the SAG and the free electrons of the NH2 groups of the acrylamide of the
Figure 5. Absorption capacity of MB, by the SAG-g-PAM11 copolymer, as a function of the pH.
Figure 6. Schematic illustration of the MB dye adsorption by SAG-g-PAM copolymer.
grafting. Figure 6 illustrates the adsorption process of the cationic dye molecules by the anionic groups of SAG-g-PAM copolymer. When the carboxylate groups completely neutralize the cationic charges of MB, the decolorization efficiency reached the maximum value. However, when the molar ratio is exceeded, destabilization can take place by electrostatic repulsion between the cationic dye molecules that are already bound by the copolymer[49]. At alkaline pH, the number of positively charged sites decreases and the number of negatively charged sites increases that favors the removal of MB. The results are in accordance with literature reports on the adsorption of cationic dyes by other polymeric adsorbents; similar trends were observed for the adsorption of Basic Violet 7 on sodium alginate modified with acrylamide[17], and for cationic dye removal by natural adsorbents[50]. Table 2 shows the %DE and adsorption capacity (q mg/g) of methylene blue dye by the SAG-g-PAM copolymers evaluated in pH 10. After 8h, the decolourization efficiency and adsorption capacity of methylene blue dye was of the same order for all graft copolymers, independent on the percentage grafting efficiency. Since the SAG-g-PAM6 copolymer with lower grafting efficiency (%GE =75%) showed the same performance that the others, this must be preferentially used in MB removing, because lower acrylamide content is used in the synthesis, and fewer polyacrylamide residue will be present in wastewater, reducing environmental contamination. An extra time or higher concentration of SAG-g-PAM copolymers seems to be useless in our experiments, since decolouring efficiency was almost hundred per cent, thus 69 q mg/g must be the adsorption capacity of the grafting copolymers produced. However, larger amounts of adsorbent would imply in greater surface area and a higher number of available sites for dye adsorption[17]. In both conditions, solution acid and basic, the adsorption of MB by the SAG was not observed, since no chemical interaction is facilitated or favoured. In neutral solution (pH 7), the neat alginate salt will be dissociated, and the ions COO- and Na+ will be solvated by water molecules. In a neutral solution, the cationic groups of MB and the sodium ions will compete by the carboxylic ions, and the system will reach an equilibrium. By lowering the solution pH with HCl, the salt NaCl is produced and the carboxylic acid is regenerated. By increasing the solution pH with NaOH, harder will be the SAG salt dissociation and free dissociated species in solution.
Table 2. Decolourization efficiency (%DE) and adsorption capacity (q) of the SAG-g-PAM copolymers in solution of MB pH 10. Sample SAG-g-PAM6 SAG-g-PAM11 SAG-g-PAM23 SAG
144 144/146
Grafting efficiency (%) 75 79 98 -
Decolourization efficiency (%DE) 99.3 ± 0.1 98.7 ± 0.0 98.7 ± 0.1 No
Adsorption capacity (q mg/g) 69.13 ± 0.1 69.10 ± 0.1 69.00 ± 0.1 No
Polímeros, 28(2), 139-146, 2018
Ultrasound-assisted synthesis of polyacrylamide-grafted sodium alginate and its application in dye removal
4. Conclusion In this work was shown that graft copolymer of sodium alginate and with different concentrations of acrylamide (SAG-g-PAM) could be synthesized in a short period of time by using ultrasound energy at low frequency. An improvement of grafting reaction was reached by using ultrasound or sonication process, since the time reaction was significantly reduced (< 10 min) in comparison to 24 h needed in the conventional grafting reaction process reported in the literature[18,40]. The grafting efficiency %GE was dependent on the acrylamide monomer concentrations and the SAG-g-PAM copolymers were obtained with %GE higher than 75%. All the graft copolymers obtained by ultrasound-assisted method showed high decolorization efficiency (99%) independent of grafting efficiency, and the maximum amount of methylene blue adsorbed was 69.13 mg/g. The efficiency of dye adsorption was dependent on the pH of the methylene blue solution. The best SAG-g-PAM copolymer to be used as adsorbent was the one with lower content of incorporated acrylamide (%GE = 75%), because lower amount of acrylamide is used in the synthesis, and there will be lesser wastewater contamination by polyacrylamide, attenuating environmental contamination.
5. Acknowledgements The authors thank the Brazilian governmental agencies CNPq and CAPES for financial support.
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ISSN 1678-5169 (Online)
http://dx.doi.org/10.1590/0104-1428.12216
Preparation and characterization of composites from plastic waste and sugar cane fiber Ricardo Yoshimitsu Miyahara1*, Fábio Luiz Melquiades2, Ezequiel Ligowski3, Andressa do Santos4, Silvia Luciana Fávaro5 and Osmar dos Reis Antunes Junior3 Departamento de Física, Universidade Estadual do Centro-Oeste – UNICENTRO, Guarapuava, PR, Brazil 2 Departamento de Física, Universidade Estadual de Londrina – UEL, Londrina, PR, Brazil 3 Departamento de Química, Universidade Estadual do Centro-Oeste – UNICENTRO, Guarapuava, PR, Brazil 4 Departamento de Química, Universidade Estadual de Maringá – UEM, Maringá, PR, Brazil 5 Departamento de Engenharia Mecânica, Universidade Estadual de Maringá – UEM, Maringá, PR, Brazil
1
*rmiyahara@unicentro.br
Abstract This study presents the preparation and characterization of composite materials using Plastic Waste from Hydrapulper (PWH) from paper industries extruded with sugar cane fiber residues from ethanol industries. The factorial design showed that composite material with 40% of sugar cane fiber, pressed with 5 ton was the optimized condition. The main findings attested that the composite is resistant up to 250 °C and its hardness is increased compared to the raw PWH. The material presented woodsy aspect although water absorption has increased. So, this study offers a good alternative for the use of plastic waste generated as a by-product of recycled paper industry as well as a destination to the sugar cane bagasse. Keywords: cane bagasse fiber, composite material, hidrapulper equipment, plastic waste, alternative material for wood.
1. Introduction The multiple application possibilities of plastic materials, stimulated by the diversity of its properties and characteristics have been conducted to its frequent use in several applications, including civil engineering. The high cost of metals and cement for the production of reinforced concrete and the lack of wood allied to the large amount of plastic material available for recycling increased the use of thermoplastic composites. The use of materials which are able to combine basic requirements, such as the conservation of natural resources and the environment preservation, can be essential for the future of the planet. Considering these problems, a challenge is the production of environmental friendly materials or perform the correct destination and reutilization of plastic materials[1,2]. The transformation process of plastic waste in composites materials is a viable option to stimulate practices that prioritizes recycling and development of new materials through the utilization of environmental liabilities. A polymeric composite is formed by two phases: a continuous one, that is a thermoplastic base matrix and the dispersive phase, which are reinforcing fillers of organic or inorganic fibers[3]. The incorporation of fiber have been widely investigated and used as reinforcement in polymer matrices such as sponge gourd fiber[4], sisal[5], curauá[5], sugar beet pulp[6] and sugarcane bagasse fiber[7]. Acacia bark residues was used as reinforcing filler in polypropylene composite can produce higher impact properties and higher degradation temperature[8]. Taflick et al.[8] atributed this behavior to the of tiny and short fibers due distributed more homogeneously of composite. Researchers done by Martins et al.[9] and
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Rzatki and Barra[10] proved the effectiveness of the natural amorphous silicate short fibers as reinforcing agents increasing the tensile strength of reinforcement matrix of the epoxy and poly (butylene terephthalate), respectively. The final product will be a polymeric and natural fiber composite as an alternative or replacement to wood and its derivatives, reducing deforestation and pollution. The residues from the Hydrapulper machinery represent high destination costs both financially and environmentally. The technological equipment, Hydrapulper, is used at the separation of polymer residues, plant fibers and additives coming from containers of the recycled paper industry. This machinery is similar to a giant blender with a pond in cylinder form . The procedure is performed adding water to the paper that will be recycled and by a physical agitation process, the paper fibers are detached forming a cellulose pulp. These fibers are washed and purified, being used for the production of paper in the form of cardboard boxes, napkins, and other products. In a sieve located below the spinning rotor, the impurities in the process, such as various plastics, non disaggregated fibers and small quantities of metals that are retained originating the Plastic Waste from Hydrapulper (PWH), which is a very difficult waste to be recycled and is commonly sent to specific landfills causing concern with the disposal and vast environmental problems. On the other hand, the abundance of natural fibers such as sisal, coconut, jute, ramie bast, eucalyptus pulp, banana, hemp, flax, pineapple leaf, bamboo, palm, cotton, waste of
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O O O O O O O O O O O O O O O O
Miyahara, R. Y., Melquiades, F. L., Ligowski, E., Santos, A., Fávaro, S. L., & Antunes Junior, O. R. mate-tea and sugarcane fiber among others and the possibility of using these fibers as reinforcement in composites stimulate the studies of different products combining these raw materials[11-15]. Each of these natural fibers has a wide range of mechanical and physical properties governing its wider applications. They are renewable, biodegradable, non-toxic, non-abrasive, they have low density with specific strength and low cost. Moreover, they are available worldwide been characterized as “environmentally-friendly” materials[16-18]. Sugarcane bagasse is the by-product obtained after sucrose extraction from the sugar cane plant or ethanol fuel plant. It is especially interesting because it is considered the main residue from Brazilian agro industry, which generates around 270 kg of bagasse per ton of cane ground with an annual accumulation of 132 million tons. They are used in industrial furnace and for second generation ethanol, organic fertilizing, animal feed and cellulose and paper production. Sugarcane fibers are composed of cellulose, hemicelluloses and lignin as well as small amounts of mineral, wax, and other compounds[19,20]. The objective of this study was to prepare and characterize composite materials using PWH from paper industries extruded with sugar cane fiber (SCF) residues from ethanol industries. In Specific (a) a factorial design has been conducted to optimize the proportions and compaction pressure in the composite preparation and (b) characterization tests consisting of physical, chemical and mechanical tests, thermal stability and morphology of the composites materials.
2. Materials and Methods 2.1 Raw materials and composite The PWH was obtained from a paper recycling factory from Guarapuava, PR, Brazil. The material showed moisture content ~50% and was first centrifuged and then dried in an oven at 100 oC. Clamps, clips, wires, nails and other metallic residues were magnetically separated by a device adapted in the entrance of a Wiley knife mill. In the sequence the plastic waste was bonded for 5 minutes in a thermo-kinetic mixer. Finally the material was cut into small particles using the same mill with a 30 mesh sieve (Figure 1a-c). Sugar cane fibers were collected from a sugar and ethanol plant in São Pedro do Ivaí, PR, Brazil. The material was dried in an oven at 100 oC and milled in Wiley knife mill with 30 mesh sieve (Figure 1d).
2.2 Preparation of the composites The PWH and sugar cane particles were mixed and homogenized at 120 rpm for 5 minutes at room temperature using a thermo-kinetic mixer. The proportions of fiber and polymer contents were performed according to the factorial design that is presented in section 2.4. For the processing of PWH and the formulated composites, a laboratory single-screw extruder with 45 mm cannon by 300 mm in length and a screw with 25 mm of diameter by 350 mm in length was used, Tornoeste brand (model ET 001).
Figure 1. (a) PWH after first milling and removal of metals. (b) PWH after agglutination process (c) PWH after micronization. (d) Particles of Sugar cane fiber. 148 148/154
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Preparation and characterization of composites from plastic waste and sugar cane fiber A rotation process of 100 rpm and four heating temperatures were employed, respectively: 160 ± 5 °C (power zone), 175 ± 5 °C (compression zone), 190 ± 5 °C (dosage) and 200 ± 5 °C on the coupling head. It was possible to check the viscosity and fluidity of plastic waste and the polymeric composites, proposed by the factorial design, by the pressure of extruded paste, evaluated on the pressure gauge of the extruder. After the extrusion process, the composites were molded for the mechanical testes according to ISO and ASTM standards. The samples were pressed with 10 or 5 tons for 90 s, according to the factorial design (section 2.4). Additives or coupling agents were not used in the molding process.
2.3 Thermal analysis Thermal analyses have been performed to analyze the thermal behavior of the raw PWH, the raw sugarcane fiber and the composite material. Tests were done with a equipment SDT Q600 from TA Instruments. Thermogravimetry (TG) and differential scanning calorimetry (DSC) have been applied to determine the weight gain, the mass loss rate and the degradation process. Five milligrams of each samples have been analyzed at the temperature range from 20 °C to 1000 °C at a heating rate of 10 °C.min-1 in a nitrogen flow of 100 mL.min-1.
2.4 Factorial design A 22 factorial design was performed with the purpose to optimize the production conditions of the polymeric composite[21], considering fiber proportions and compaction pressure. Initially, one sample of each of the following charge content (sugar cane fiber) has been prepared: 20%, 30%, 40% and 50%. The factorial design was performed considering 30% and 40% because with 50% the extrusion process is very hard and the material has low hardness. With 20% of fiber content the material loses its wood appearance. The compaction factor was evaluated to analyze if differences in the pressure eliminates some empty spaces in the interfacial region of the materials which interfere in moisture and mechanical hardness of the samples. The studied values were 10 and 5 ton. Table 1 present the factorial design experiment. Five samples of each condition were prepared and evaluated for tensile strength, impact, hardness, moisture modifications and composite density.
2.5 Physical and mechanical tests Water absorption was carried out according to ASTM D-1037 with 24 h immersion time in a thermostatic batch at 20°C. The body tests were weighed before and Table 1. Factorial design description. Essay 1 2 3 4
Factor 1
Factor 2
(proportion) 70% PWH/30% bagasse fiber (-) 60% PWH/40% bagasse fiber (+) 70% PWH/30% bagasse fiber (-) 60% PWH/40% bagasse fiber (+)
(compaction) 5 ton (-) 5 ton (-) 10 ton (+) 10 ton (+)
PWH – plastic waste from hydrapulper.
Polímeros, 28(2), 147-154, 2018
after the immersions. Water absorption was determined using Equation 1: % moisture =
wet sample mass − dry sample mass × 100 (1) dry sample mass
The density of the composites and of the raw PWH was carried out with a Multipycnometer from Quantachrome Insruments using He gas. The measurements were performed in triplicate. The mechanical properties of the raw PWH and of the composites were evaluated through tensile strength at break (ASTM D-638) in an EMIC DL 10000 machine, flexural strength (ASTM D-790) in an LLOYD LR10K, Impact strength (Izod) with V notch (ASTM D-256) at room temperature in a RESIL impactor 5R, compression strength (ISO 604) in a EMIC DL 30000 and hardness shore (ABNT 7456) with a WOLTEST GSD 702D. All mechanical properties were determined using at least 5 samples for each test.
2.6 Morphological analysis Analysis of fracturated surfaces from raw PWH and from polymeric compounds obtained from Izod Impact test were performed by SEM using a TM300 Hitachi microscope. The samples were fixed in a carbon tape prior to analysis. The PWH and sugar cane particles were mixed and homogenized at 120 rpm for 5 minutes at room temperature using a thermo-kinetic mixer. The proportions of fiber and polymer contents were performed according to the factorial design that is presented in section 2.4. For the processing of PWH and the formulated composites, a laboratory single-screw extruder with 45 mm cannon by 300 mm in length and a screw with 25 mm of diameter by 350 mm in length was used, Tornoeste brand (model ET 001). A rotation process of 100 rpm and four heating temperatures were employed, respectively: 160 ± 5 °C (power zone), 175 ± 5 °C (compression zone), 190 ± 5 °C (dosage) and 200 ± 5 °C on the coupling head. It was possible to check the viscosity and fluidity of plastic waste and the polymeric composites, proposed by the factorial design, by the pressure of extruded paste, evaluated on the pressure gauge of the extruder. After the extrusion process, the composites were molded for the mechanical testes according to ISO and ASTM standards. The samples were pressed with 10 or 5 tons for 90 s, according to the factorial design (section 2.4). Additives or coupling agents were not used in the molding process.
3. Results and Discussions 3.1 Thermal analysis Figure 2 presents the thermograms for raw fiber, raw PWH and composites with 30% and 40% of fiber charge. It was observed that all these materials keep stable up to 250 °C eliminating retained moisture. It demonstrates that these raw materials are apt for extrusion, injection and thermoforming process, since 200 ± 5 °C was the maximum temperature applied in the composites preparation. The curves for raw PWH show mass loss and small degradation in the 250 °C to 450 °C interval. In the end of 149/154 149
Miyahara, R. Y., Melquiades, F. L., Ligowski, E., Santos, A., Fávaro, S. L., & Antunes Junior, O. R. the composition process, the residual mass is smaller than in the polymeric composites. Between 250 °C and 450 °C the composite with 40% fiber loses less mass and degrades more slowly than the sample with 30% fiber. It is in agreement with the literature which shows that natural fiber compounds have higher heat resistance and consequently high resistance to decomposition. So, cellulose and lignin contribute to the thermal resistance of the polymeric composite with higher fiber charge [22]. For temperatures over 450 °C the mass loss is equivalent. Endothermic crystalline transitions are evidenced by DSC curves, Figure 3. For temperatures below 120 °C second order transitions are present, demonstrating the presence in smaller quantities of amorphous polymers (PS, ABS, PVC). Over 200 °C the discontinuities occur due to the presence of semi-crystalline high fusion polymers (PET, PA 6/6.6, PC)[23,24]. The main events are polyolefin fusion (polyethylene and polypropylene) at 125.6 °C and 160.8 °C, besides the onset thermal degradation at 266.5 °C and offset at 465.9 °C, demonstrated in heat absorption peaks. As for sugar cane bagasse fibers, an endothermic peak was observed at 66.6 °C related to the moisture loss. The peaks at 300.3 °C (endothermic) and 355.3 °C (exothermic) represent hemicelluloses and cellulose degradation, respectively. Lignin degradation occurs at 395 °C as shown by the exothermic peak[25,26].
is denser than the polyolefin due to smaller proportion of thermoplastics, thermosets and elastomers which enlarge the final density.
Figure 2. Thermograms (TG) curves for sugar cane fiber (SCF), and polymeric composites with 30% (I 30) and 40% (I 40) fiber charge, and pure plastic waste from hydrapulper (PWH).
3.2 Physical and mechanical properties Incorporated as polymeric matrix in composites, the purpose is to use PWH to replace raw polymers like polyolefin that are the most used thermoplastics. Table 2 presents the results of physical and mechanical properties of the raw PWH and also compares it with the main polymers commonly used to prepare composites. The PWH absorbs higher water quantity when compared to polyethylene and polypropylene because in this plastic waste there is an incorporation of small quantities of cellulose residues in the plastic surface. The presence of these fibers and other impurities increases water retention. The PWH
Figure 3. Differential Scanning Calorimety (DSC) curves for sugar cane fiber (SCF), polymeric composites with 30% (I 30) and 40% (I 40) fiber charge, and pure plastic waste from hydrapulper (PWH).
Table 2. Mechanical and physical properties of PWH compared to virgin and recycled polymers. Polymeric Matrix PWH PEHD Recycled PEHD Virgin PP Recycled PP Virgin
Hardness Shore D (0 to 100) 41.4 ± 0.7 62.0 ± 0.79[27] -----------64.0[30] 62.9 ± 0.55[27] 65.0[30] -----------68.2 ± 0.57[27] ----------------------71.0 ± 0.71[27] -----------------------
Flexural strength
Tensile strength
18.6 ± 0.3 18.99 ± 0.45[27] ≅ 20.5[2] 24.87[30] ------------24.87[28] 19.76 ± 0.39[31] 41.74 ± 0.49[27] -----------18.2 [33] 43.48 ± 0.51[27] -----------55.74[36] ------------
9.6 ± 0.4 20.73 ± 0.44[27] ≅ 21.0[2] 25.66[30] 24.53 ± 0.12[27] 24.73[30] 18.9 ± 0.2[31] 28.96 ± 0.67[27] 22.78 ± 1.14[32] 16.1[33] 30.72 ± 0.45[27] 37.36 ± 1.87[33] 34.11[36] 23.4 ± 1.3[37]
Izod impact strength (J/m) 73.9 ± 2.0 -----------≅ 46.5 [2] 33.9[30] -----------38.34[30] ----------------------26.25 ± 4.5[32] -----------------------25.05 ± 1.6[33] 5.68[36] 8.5[37]
Moisture absorption (%)
Density (g cm-3)
0.95 ± 0.06 < 0.01[28,29]
1.00 ± 0.06 0.94-0.98[28,29]
< 0.03[34,35]
0.90-0.91[29,35]
PWH = Plastic waste from hydrapulper; PEHD = polyethylene high-density; PP = Polypropylen.
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Preparation and characterization of composites from plastic waste and sugar cane fiber Shore D hardness and tensile strength presents smaller values in the PWH compared to the other polymers. For flexural strength the values were similar to virgin or recycled PEAD. As mentioned before, the PWH is composed by a mix of thermoplastics, thermosets and elastomers that are melted to form the polymeric matrix. But the polymers that have high melting temperatures are incorporated to the composite as particulate charge generating micrometric empty spaces between the particulate and the polymeric matrix. This fact restricts the mechanical properties that require higher interface adhesion of the materials as noted in the tensile strength test. In the Izod impact test is possible to verify that the PWH has greater resistance compared to the polyolefin. So the composite could be employed in applications that require high impact resistance. Factorial design results and the comparison among the raw PWH and the composites are presented in Table 3. When shore hardness is analyzed, it is possible to verify that with 5 ton the results are equivalent in both proportions, but with 10 ton and 40% fiber the shore hardness decreased. For tensile strength all the results are equivalent considering the standard deviation with a slightly decrease for higher pressure quantities. The same equivalence was obtained for
traction test. In the impact test, the samples with higher fiber proportion have significant higher Izod impact strength, reaching around 58 J.m-1. The resistance to compression was slightly better for 5 ton compaction. The moisture physical test shows that increasing the fiber content, it increases the moisture. And also at 10 ton the moisture is higher than in 5 ton. Density of the material remains constant independently on the proportion and compaction value. Comparing the composite material with raw PHW it is possible to verify that the reinforcement with sugar cane fiber, although present visually good appearance, decreases the mechanical resistance and increases moisture and density in the samples. Inserting the fiber content in the PWH makes it harder than the composite. Analyzing the factorial design results the application of 10 ton seems to be inefficient. We suppose that with higher compaction pressure the natural fibers shifts from the interface region decreasing the interaction with the polymeric matrix. With the pressure excess, cracks were formed in the sample decreasing the mechanical properties of the final composite. Considering the four possibilities of the factorial design, the tests with higher fiber proportion and lower compaction pressure demonstrated the best result. Table 4 present a
Table 3. Factorial Design results for the polymeric composites compared to the mechanical and physical properties of raw PWH. Values and standard deviation for 5 different samples. Composite 1
Composite 2
Composite 3
Composite 4
70% PWH/30% SCF
60% PWH/40% SCF
70% PWH/30% SCF
60% PWH/40% SCF
Property
Raw PWH
Hardness (Shore D)
41.4 ± 0.7
5 ton 51.9 ± 2.8
5 ton 52.2 ± 1.6
10 ton 51.5 ± 1.3
10 ton 48.1 ± 1.3
18.6 ± 0.3
16.9 ± 0.6
18.1 ± 0.4
15.6 ± 0.7
14.3 ± 0.5
9.6 ± 0.4
7.5 ± 0.3
7.6 ± 0.7
7.5 ± 0.4
7.9 ± 0.8
73.9 ± 2.0
45.6 ± 1.5
57.5 ± 4.0
50.5 ± 4.7
58.0 ± 2.7
27.7 ± 0.9
11.9 ± 0.8
13.9 ± 0.5
8.9 ± 1.2
10.3 ± 1.8
0.90 ± 0.06 1.00 ± 0.06
1.8 ± 0.1 1.22 ± 0.01
2.1 ± 0.1 1.19 ± 0.01
1.9 ± 0.1 1.20 ± 0.01
2.4 ± 0.2 1.18 ± 0.01
(0-100 scale) Flexural strenght (MPa) Tensile resistance (MPa) Impact resistance (Izod) (J/m) Compression strength (MPa) Moisture retention (%) Density (g cm–3)
PWH = Plastic waste from hydrapulper, SCF = Sugar cane bagasse fiber.
Table 4. Comparison of the optimazed condition reached in this study with literature results. Tensile
Flexural
Conditions/references
strength
strength
Izod impact strength
(MPa) 7.6 ± 0.7 4.69 22.71 ± 0.10 28.62 ± 0.13 8.11 ± 1.44 6.42 ± 1.11 20.4 ± 0.3 21.2 22.64 ± 1.25
(MPa) 18.1 ± 0.4 3.96 ± 0.16 35.27 ± 1.02 49.74 ± 1.39 13.16 ± 1.64 17.76 ± 1.06 25.3 31.66 ± 3.25
(J/m)
(MPa)
60% PWH/40% sugar cane fiber (our study) 1-80% Epoxi resin/15% FBC/5% glass fiber[13] 2-80% raw PEAD/20% sisal fiber[38] 2-80% raw PP/20% sisal fiber[38] 3-50% recycled PEAD/50% wheat straw[39] 3-50% recycled PP/50% wheat straw[39] 4-90% recycled PEAD reciclado/10% banana fiber[40] 5-90% recycled PEAD/10% sisal fiber[2] 6-50% PSAI/50% pineapple leaf fiber[41]
57.5 ± 4.0 24.01 ± 0.28 65.5 ± 2.37 34.60 ± 2.56 47.37 ± 6.42 11.06 ± 1.50 45.2 ± 1.9 62.0 24.39 ± 1.34
13.9 ± 0.5 9.60 -
Compression strength
PWH = plastic waste from hydrapulper; PEHD = polyethylene high-density; PP = Polypropylen; HIPS = high impact polystyrene.
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Miyahara, R. Y., Melquiades, F. L., Ligowski, E., Santos, A., Fávaro, S. L., & Antunes Junior, O. R.
Figure 4. Scanning electron microscopy of the (a) fracture of raw plastic waste from hydrapulper and (b) higher magnification with empty spaces.
comparison with the best result of our study with other similar studies in the literature.
3.3 Scanning electron microscopy (SEM) An image of the fractured surface from the Izod impact sample of raw PWH is presented in Figure 4. It can be noted that there are components that were not melted and kept in the samples as charged particles. The Figure 4b highlights the empty spaces between melted and not melted polymers (white points).This indicates the necessity of some kind of reinforcement charge or some coupling agent to increase the interfacial adhesion[42,43]. Thakur et al. have shown that is possible a surface modification, for example, of natural polymers using silane coupling agent[44] or graft copolymerization[45], this change in morphology can causes changes in the properties of the polymers. Figure 5a and b present the micrographs for the composites with 30% and 40% of natural fibers. Compared to the micrography of the raw PWH is noted that the reinforcement charge incorporation filled the porous to obtain a more homogeneous material. Figure 5c, with amplification power of 1000 times in the sample with 40% fiber it still presents several “white dots” related to the polymers not melted, it attests that sugar cane fiber charge was not able to cause an efficient adhesion. However, the final product has interesting properties and applications. 152 152/154
Figure 5. Scanning electron microscopy of the fracture of composites (a) 30% sugar cane fiber, 100x; (b) 40% sugar cane fiber, 100x; (c) 40% sugar cane fiber, 1000x.
4. Conclusions Thermal analysis attested that the composite is resistant up to 250 °C, so it is not degraded in extrusion, injection and thermoset process. The PWH is composed in its majority by polyolefin (PEAD, PP, PEBD) as concluded by DSC analysis. Compared to the raw PWH, the addition of sugar cane as reinforcement fiber increased the hardness of the material. The other physical and mechanical properties were not improved. Higher fiber proportion in the composite presented positive effects, mainly in the compression and impact tests, although water absorption increased. It was observed that the material presented a woodsy aspect losing its plastic appearance. So, the composite material with 40% of sugar cane fiber, pressed with 5 ton just after extrusion was the optimized condition reached in this study. Higher Polímeros, 28(2), 147-154, 2018
Preparation and characterization of composites from plastic waste and sugar cane fiber pressure values caused a decrease in resistance and an increased in moisture. In the segment of recycled paper and cellulose this study contributes to give a profitable destination to the plastic waste generated as a by-product of recycled paper industry. The polymeric composite produced presents a material with strong environmental appeal as an alternative do wood and some derivatives to be used in buildings and furniture industries.
5. Acknowledgements The authors wish to thank CNPq (National Council of Technological and Scientific Development), CAPES (Coordination for the Improvement of Higher Education Personnel) and MCTI (Ministry of Science, Technology and Innovation), Conv. n°: 01.0010.00/2011.
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Polímeros, 28(2), 147-154, 2018
ISSN 1678-5169 (Online)
http://dx.doi.org/10.1590/0104-1428.02517
Ultrasound assisted miniemulsion polymerization to prepare poly(urea-urethane) nanoparticles André Eliezer Polloni1, Alexsandra Valério1, Débora de Oliveira1, Pedro Henrique Hermes de Araújo1 and Claudia Sayer1* Laboratório de Controle de Processos e Polimerização – LCP, Departamento de Engenharia Química e Engenharia de Alimentos – EQA, Universidade Federal de Santa Catarina – UFSC, Florianópolis, SC, Brazil
1
*claudia.sayer@ufsc.br
Abstract Recently, the physical and chemical effects of ultrasound in polymeric materials synthesis have attracted great attention. This work presents the synthesis of novel polymeric materials by polymerization of isophorone diisocyanate with different polyols. Polymers were synthesized by step miniemulsion polymerizations, using ultrasound bath and thermostatic bath. The effects of ultrasound, temperature and polyol type were evaluated by Fourier transform infrared spectroscopy, gel permeation chromatography, dynamic light scattering and titrimetry. Polymerization under ultrasound bath showed that different reaction temperatures in the range between 50 °C and 80 °C directly influence the molecular weight of the polymers, urea/urethane formation and increase of diisocyanate consumption rate. In addition, different polyols used in polymerizations in miniemulsion had a significant effect on the characteristics of the resulting poly(urea-urethane) nanoparticles. Finally, ultrasound assisted polymerizations showed a faster diisocyanate consumption rate, but did not lead to enhanced molecular weights. Keywords: miniemulsion polymerization, poly(urea-urethane), ultrasound.
1. Introduction Synthesis of new polymeric materials have attracted the interest of researchers, chemical and pharmaceutical industries. Continuous improvement of polymers and polymerization methods is a still growing field of interest. Production of materials with different characteristics, as surface characteristics, molecular weight, crosslinking degree and inorganic materials addition is a very active research field[1-7]. Miniemulsion is a polymerization technique that has been gaining much attention due to some advantages when compared to conventional emulsion polymerization[8,9]. Miniemulsions are described as aqueous dispersions of relatively stable oil droplets within a size range (from 50 to 500 nm), prepared in a system containing a dispersed phase (organic), continuous phase (aqueous), an emulsifier and a co-stabilizer[10,11]. To obtain this dispersion, a mechanism of high shear stress is required to break the monomer droplets into submicron droplets, reaching a steady state obtained by balancing the rates of coalescence and breakage of droplets that are kinetically, but not thermodynamically stable. In this way, since these droplets are protected against molecular diffusion (Ostwald Rippening) and coalescence, the size of polymeric particles formed after reaction is expected to remain practically the same as that of the droplets formed during miniemulsification[12,13]. In addition, the latex formed by miniemulsion polymerization may also exhibit viscosity and colloidal stability different from those of conventional emulsion polymerizations, resulting in more stable latexes with same surfactant concentration, possibility of obtaining latexes with a high solids content and hybrid organic/inorganic and organic/organic materials[13,14].
Polímeros, 28(2), 155-160, 2018
The development of poly(urea-urethane) (PUU) has been extensively studied in the last decades due to their excellent physical properties[2,15-17] ranging from very soft elastomers to very rigid plastics[4,16,18]. Poly(urea-urethane) nanoparticles obtained by miniemulsion polymerization are being used in different areas such as pharmaceutical, medical and cosmetic, especially as a means of drug delivery because of their excellent physical and biocompatibility properties[2,4,19-21]. In PUU synthesis, the main reaction involves the formation of urethane segments in the reaction between isocyanates (NCO) and OH groups from a polyol, Figure 1. Urea groups may be formed in a secondary reaction with water and with the release of CO2, as shown in Figure 2, wherein water reacts with a diisocyanate to form an amine and CO2. The formed amine can rapidly react with an isocyanate group to generate a compound with a urea bond, modifying the final properties of PUU[22]. The use of ultrasound bath in organic synthesis has been broadly expanded in recent years since it may enhance the reaction rate and selectivity of the product rather significantly[23-25]. Many studies have been carried out and advantages of ultrasound procedures as good yields, short reaction times and mild conditions have been reported[24,26]. In this way, the use of ultrasound in organic synthesis is now recognized as a viable environmentally benign alternative[27]. When ultrasound propagates in sound-bearing media, it can cause some effects on this media including mechanical, thermal and cavitation effect[23,26,28,29]. Few works report the use of ultrasound with the aim to affect the course of
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O O O O O O O O O O O O O O O O
Polloni, A. E., Valério, A., Oliveira, D., Araújo, P. H. H., & Sayer, C. polymerization reactions[23,24,30,31]. Price et al.[31] present an investigation of the effect of ultrasound on the synthesis of polyurethane in bulk polymerization using a diisocyanate:diol molar ratio of 1:1. The ultrasound system used was a sonic horn system operating at 23 kHz. Polymerization times and molecular weights have been improved by the use of high intensity ultrasound, and increasing the ultrasound intensity lowered the reaction times but had no significant effect of the molecular weight of the polymers formed. The authors attributed this effect of ultrasound heating to local heating around collapsing cavitation bubbles together with the enhanced mass transfer caused by the fluid motion. In this work the influence of ultrasound bath, temperature and different polyols on NCO consumption rate and on final characteristics (urea/urethane ratio, molecular weight and particle diameter) of poly(urea-urethane) nanoparticles produced by step miniemulsion polymerization were evaluated. Among the polyols used some were oil soluble and others water soluble leading to interfacial polymerization.
2.Materials and Methods 2.1 Materials For the synthesis of poly(urea-urethane) nanoparticles the following reactants were used: castor oil (100%, Mw 928 g mol-1) from Linfar, isophorone diisocyanate (IPDI, 98%, Mw 222 g mol-1), poly(ethylene glycol) diol with nominal molecular weights of 400 Da (PEG400) and 1000 Da (PEG1000), 1,6-hexanediol (99%), glycerol (99%), and surfactant sodium dodecyl sulfate (SDS, from Sigma–Aldrich). N-Dibutylamine (Vetec, 99%), toluene (Vetec, 99.5%), propanol (Synth, 99.5%), chloridric acid 1N (HCl, Cronoline, PA), and bromophenol blue 0.1% (Lafan) were used for free NCO quantification. All reagents were used as received.
2.2 Methods
This miniemulsification step was conducted using an ice bath to prevent polymerization. After that, the miniemulsion was placed in two different reactors: a jacketed reactor with constant magnetic stirring at 70 °C unless mentioned, for 3 hat low pressure (1 atm) or in glass vials inserted in an ultrasound bath (without agitation) (USC-1880A, 37 kHz, 132 watts, 3.8 liters, UNIQUE) at 70 °C unless mentioned, for 3 h at low pressure (1 atm). Special care was taken with samples inserted in ultrasound bath, ie.; samples were always placed in the same position inside the ultrasound bath and the amount of destilled water inside the bath was always the same, avoiding differences in cavitation formation during the reactions. A schematic representation of miniemulsification and polymerization processes is presented in Scheme 1. 2.1.2 Characterization Free NCO concentration was quantified based on the NCO standard dibutylamine back-titration method[32], titrations were conducted in duplicate. The molecular weight was evaluated by gel permeation chromatography (GPC in a High Performance Liquid Chromatograph (HPLC, model LC-20A, Shimadzu, equipped with a RID-10A detector in tetrahydrofuran (THF) at 35 °C). GPC analyses were carried out by injecting 20 μL of a 0.5 wt% polymer solution (solvent THF, Merck), previously filtered through a Teflon-filter with a mesh size of 450 nm. A column set was employed consisting of three 300 x 8 mm columns in series (GPC-801, GPC-804 and GPC-807). Polystyrene standards between
Figure 1. Polyurethane formation by the reaction between a polyol and a diisocyanate.
2.1.1 Synthesis of PUU nanoparticles by miniemulsion polymerization PUU nanoparticles were prepared by miniemulsion step polymerization according to a procedure based on a previous work[4]. Aqueous phase (90 wt % relative to the total formulation) was prepared with 10 wt% of surfactant (SDS) relative to the organic phase. The organic phase was prepared with IPDI and polyol, keeping the molar ratio NCO:OH constant for all reactions at 1.5:1. When castor oil/PEG 400 and castor oil/PEG 1000 were used as polyols, they were dissolved in organic phase along with IPDI and remained inside the monomer droplets; thereby the polymerization was inside monomer droplets. When 1,6-hexanediol and glycerol were used as polyols, they were dissolved in the aqueous phase while the IPDI was dispersed in the organic phase and polymerization occurred in the organic/aqueous phase interface. At first, the aqueous phase was added slowly to the organic phase under magnetic stirring and kept for 5 min at room temperature to form a coarse emulsion. In sequence, the miniemulsion was prepared by sonication of the previous emulsion for 120 seconds at 70% of amplitude (Ultrasonic Dismembrator model 500, Fisher Scientific – 400 W). 156 156/160
Figure 2. Urea formation by reaction between a isocyanate group and water.
Scheme 1. Schematic representation of PUU nanoparticles synthesis via miniemulsion polymerization. Polímeros, 28(2), 155-160, 2018
Ultrasound assisted miniemulsion polymerization to prepare poly(urea-urethane) nanoparticles 580 g mol-1 and 3,800,000 g mol-1 were used to calculate average molecular weights. Average diameters (intensity averages – Dpz) of PUU nanoparticles were measured at 25 °C using dynamic light scattering equipment (DLS, Zetasizer Nano S, from Malvern). Analyses were carried out in duplicate. Fourier transform infrared spectroscopy (FTIR, IR Prestige-21, Shimadzu), using the resolution of 4.0 cm-1, was used to identify features of IPDI, the absorption band with peak location at 2270 cm-1, related to N=C=O stretching vibration of isocyanate groups was used. At the end of the reaction, the band located between 1680-1650 cm-1 relative to N-H group of urea and the absorption band between 1740-1700 cm-1 due to stretching vibration of C=O group of urethane, were used to identify peaks of poly(urea-urethane)[4].
3.Results and Discussions 3.1 Influence of polyol type in the PUU synthesis using ultrasound and thermostatic bath For the study of the polyol type influence on the synthesis of PUU nanoparticles, reactions were carried out in ultrasound bath and thermostatic bath, using as monomers castor oil/PEG 400, castor oil/PEG 1000, 1,6-hexanediol, and glycerol, besides IPDI. Reactions were conducted at 70 °C for both reaction systems and when in ultrasound bath, power was kept constant at 132 W. When thermostatic bath was used, reactions were kept under magnetic stirring. FTIR (Figure 3 and Figure 4) was used to confirm urethane and urea formation during these miniemulsion polymerizations. The characteristic carbonyl stretching was observed at 1740-1700 cm-1, indicating the presence of urethane linkage. The absorption band of urea groups (–NH) was observed between 1680-1650 cm-1. When reactions were conducted under ultrasound bath, absence of absorbance at 2270cm-1 (N–C–O stretching vibration) indicates that all isocyanate groups were consumed during the reaction. In the reactions conducted in the thermostatic bath, the absorption band located at 2270 cm-1 was still found and that could be explained by the fact that these reactions were slower when compared with the same reactions in as ultrasound bath, thereby the isocyanate groups were not completely consumed at the end of reaction time. Figure 5 shows urethane/urea area ratios calculated from FTIR spectra of PUU synthesized by step miniemulsion polymerization using ultrasound bath at 132 W or thermostatic bath, both at 70°C.To obtain such ratios, the areas of the respective peaks (1740-1700 cm-1 for urethane bonds and 1680-1650cm-1 for urea groups) were integrated and then compared. When the water soluble polyols, 1,6-hexanediol and glycerol were used, it was not possible to separate the peaks related to urethane linkages in the FTIR spectra. Therefore, the ratio between urethane/urea peaks was calculated only for reactions using castor oil/PEG 400 and castor oil/PEG 1000 as polyols. It can be observed that reactions conducted under ultrasound bath resulted in a slightly lower ratio between the urethane/urea peaks. This indicates that the use of ultrasound bath led to a small increase in the reaction between isocyanate groups and OH from water, generating higher amount of polyurea. Polímeros, 28(2), 155-160, 2018
Figure 3. FTIR spectra of PUU synthesized by step miniemulsion polymerization using ultrasound bath at 132 W at 70 °C with different polyols type.
Figure 4. FTIR spectra of PUU synthesized by step miniemulsion polymerization using thermostatic bath at 70 °C with different polyols type.
Figure 5. Urethane/urea ratio calculated from FTIR areas of PUU produced using (a) castor oil/PEG 400 (ultrasound bath), (b) castor oil/PEG 400 (thermostatic bath), (c) castor oil/PEG 1000 (ultrasound bath) and (d) castor oil/PEG 1000 (thermostatic bath). 157/160 157
Polloni, A. E., Valério, A., Oliveira, D., Araújo, P. H. H., & Sayer, C. As can be observed in Table 1, when replacing 1,6-hexanediol (molecular weight 118 g mol-1) by the mixture of castor oil/PEG 400 and castor oil/PEG 1000 as monomers (molecular weight 928/400 and 928/1000 g mol-1, respectively) PUU molecular weight was increased. Moreover, reactions conducted in US bath resulted in lower molecular weight values when using 1,6-hexanediol and castor oil/PEG 1000 as monomers, probably due to the increase of the reaction forming urea bonds, as observed in the FTIR results. This reaction uses two NCO groups for each OH group from water, thus reducing the molecular weight. Due to the fact that these reactions were conducted in miniemulsion (using water as continuous phase) these results are different from those of Price et al.[31] who polymerized different isocyanates with different diols in bulk using an ultrasound and observed that the rate of reaction could be accelerated and molecular weights were increased in comparison with reactions without the use of ultrasound.
increase of the NCO consumption rate, as already observed for reactions under conventional heating[4]. This behavior is attributed to the higher water solubility at 80 °C increasing the comtribution of side reactions generating urea bonds. The ratio between the diameter of PUU nanoparticles and the initial diameter of the monomer droplets at different reaction temperatures is shown in Figure 7, it is possible to see that PUU nanoparticles size remained constant during the reaction indicating that droplets/particles remained stable
3.2 Effect of temperature in the PUU synthesis using ultrasound and thermostatic bath Initially, the influence of temperature in a range from 50 °C up to 80 °C was studied and the results using ultrasound bath were compared with conventional reaction at 70 °C. The use of thermostatic bath at 70 °C was based on previous studies that related 70 °C as optimum for PUU synthesis[4]. The polyol used to study the influence of temperature was castor oil/PEG 400 in a molar ratio of 9:1. Results in terms of NCO consumption are shown in Figure 6. As shown in Figure 6, lower reaction rates were obtained at 50 °C, what can be explained by the fact that this temperature will not result in an efficient mobility of monomer molecules hindering the reaction between functional groups (NCO-OH), as reported in the literature[19]. Results obtained for reactions at 70 °C using ultrasound bath showed higher NCO consumption rate when compared with the same reactions using thermostatic bath. The increase in NCO consumption rate can be explained by the ultrasound effect, since it causes an increase in molecules mobility, and at 70 °C the effect of cavitation and bubble collapse are higher,enhancing the reaction of NCO groups with OH groups of the polyol and with water to form, respectively, urethane and urea linkages. Another explanation for the higher NCO consumption under ultrasound bath would be the heating caused by sonication: the collapse of micro bubbles formed by cavitation results in an elevation of local temperature and, thus, local reaction rate, but does not affect the system temperature as a whole. Increasing reaction temperature further to 80 °C did not lead to a further
Figure 6. NCO consumption for poly(urea-urethane) reactions obtained by miniemulsion using thermostatic bath at 70 °C and ultrasound bath at 132 W at different temperatures.
Figure 7. Diameter evolution of PUU nanoparticles during step miniemulsion polymerizations with thermostatic bath at 70 °C and ultrasound bath at 132 W at 50, 70 and 80 °C.
Table 1. Weight average (Mw) and number average (Mn) molecular weights and dispersities (Ð) of PUU nanoparticles obtained in ultrasound bath at 70°C, using different polyols as monomers. Polyol
Mw
US 132 W Mn
1,6-hexanediol Glycerol Castor oil/PEG 400 Castor oil/PEG 1000
(g mol-1) 6200 * 19100 20650
(g mol-1) 3500 * 9450 9400
Đ 1.7 * 2.0 2.2
Mw
Thermostatic bath Mn
(gmol-1) 12600 9100 20400 29700
(g mol-1) 5500 4500 9600 10700
Đ 2.3 2.0 2.1 2.8
*not soluble in THF.
158 158/160
Polímeros, 28(2), 155-160, 2018
Ultrasound assisted miniemulsion polymerization to prepare poly(urea-urethane) nanoparticles Table 2. Weight average (Mw) and number average (Mn) molecular weights and dispersities (Ð) of PUU nanoparticles obtained by step miniemulsion polymerizations using ultrasound bath at different temperatures using isophorone diisocyanate and castor oil/PEG 400 as monomers. Temperature (°C) 50 70 80
Mw (g/mol) 10800 19100 9690
Mn (g/mol) 6900 9450 6550
Đ 1.56 2.02 1.48
during polymerization and that the ultrasound bath did not affect their size (around 200 nm). Table 2 shows the molecular weights of PUU nanoparticles obtained using ultrasound bath at 132 W and at different temperatures. When the temperature was increased from 50 °C to 70 °C, one can observe an increase in molecular weight from 10800 to 19100 g mol-1. On the other hand, a further increase in temperature led to a decrease of molecular weight to around 9690 g mol-1. This decrease can be attributed to side reactions of isocyanate group with water to form urea linkages with the release of carbon dioxide, resulting in lower molecular weights, since two isocyanate groups are consumed in the formation of each urea group. This occurs because the concentration of water in the organic phase increases with temperature[4].
4. Conclusions In this work, the effect of ultrasound bath in step miniemulsion polymerization of IPDI with different polyols was evaluated. Results show that molecular weights of poly(urea-urethane) was strongly influenced by the polyol type. Higher molecular weights were obtained using the mixture of castor oil/PEG 400 and castor oil/PEG 1000 as polyols. In addition, the reaction temperature also affected the molecular weight and reaction rates of materials produced. Results obtained for reactions at 70°C using ultrasound bath showed higher reaction rates when compared with the same reactions using thermostatic bath. Finally, the samples produced using ultrasound bath showed lower values of molecular weight when compared with those produced when thermostatic bath was used, because ultrasound waves facilitate the hydrolysis of isocyanate, thus forming urea. Ultrasound bath, besides of accelerating the reaction, led to formation of larger fractions of urea and this can be a useful tool for producing polymers with different urea/urethane ratios, creating materials with distinctive characteristics.
5. Acknowledgements The authors thank the financial support from CAPES (Coordenação de Aperfeiçoamento de Pessoal de Nível Superior) and CNPq (Conselho Nacional de Desenvolvimento Científico e Tecnológico).
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Polloni, A. E., Valério, A., Oliveira, D., Araújo, P. H. H., & Sayer, C. chitosan. Polymer, 51(9), 1926-1933. http://dx.doi.org/10.1016/j. polymer.2010.03.008. 16. Zhou, L., Yu, L., Ding, M., Li, J., Tan, H., Wang, Z., & Fu, Q. (2011). Synthesis and characterization of ph-sensitive biodegradable polyurethane for potential drug delivery applications. Macromolecules, 44(4), 857-864. http://dx.doi. org/10.1021/ma102346a. 17. Miao, S., Wang, P., Su, Z., Liu, Y., & Zhang, S. (2012). Soybean oil‐based shape‐memory polyurethanes: Synthesis and characterization. European Journal of Lipid Science and Technology, 114(12), 1345-1351. http://dx.doi.org/10.1002/ ejlt.201200219. 18. Wang, Z., Yu, L., Ding, M., Tan, H., Li, J., & Fu, Q. (2010). Preparation and rapid degradation of nontoxic biodegradable polyurethanes based on poly(lactic acid)‐poly(ethylene glycol)‐ poly (lactic acid) and L‐lysine diisocyanate. Polymer Chemistry, 2(3), 601-607. http://dx.doi.org/10.1039/C0PY00235F. 19. Valério, A., da Rocha, S. R. P., Araújo, P. H. H., & Sayer, C. (2013). Degradable polyurethane nanoparticles containing vegetable oils. European Journal of Lipid Science and Technology, 116(1), 24-30. http://dx.doi.org/10.1002/ejlt.201300214. 20. Zanetti-Ramos, B. G., Lemos-Senna, E., Cramail, H., Cloutet, E., Borsali, R., & Soldi, V. (2008). The role of surfactant in the miniemulsion polymerization of biodegradable polyurethane nanoparticles. Materials Science and Engineering C, 28(4), 526-531. http://dx.doi.org/10.1016/j.msec.2007.04.041. 21. Gaudin, F., & Sintes-Zydowicz, N. (2011). Poly(urethane–urea) nanocapsules prepared by interfacial step polymerization in miniemulsion: the droplet size: a key-factor for the molecular and thermal characteristics of the polymeric membrane of the nanocapsules? Colloids and Surfaces A: Physicochemical and Engineering Aspects, 384, 698-712. http://dx.doi.org/10.1016/j. colsurfa.2011.05.050. 22. Barrére, M., & Landfester, K. (2003). High Molecular Weight Polyurethane and Polymer Hybrid Particles in Aqueous Miniemulsion. Macromolecules, 36(14), 5119-5125. http:// dx.doi.org/10.1021/ma025981+. 23. Price, G. J. (1996). Ultrasonically enhanced polymer synthesis. Ultrasonics Sonochemistry, 3(3), 229-238. http://dx.doi. org/10.1016/S1350-4177(96)00031-4. 24. Price, G. J., Lenz, E. J., & Ansell, C. W. G. (2002). The effect of high-intensity ultrasound on the ring-opening polymerisation
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Polímeros, 28(2), 155-160, 2018
ISSN 1678-5169 (Online)
http://dx.doi.org/10.1590/0104-1428.12416
Effect of heat cycling on melting and crystallization of PHB/TiO2 compounds Nichollas Guimarães Jaques1, Ingridy Dayane dos Santos Silva1, Manoel da Cruz Barbosa Neto1, Andreas Ries2, Eduardo Luis Canedo3 and Renate Maria Ramos Wellen1* Departamento de Engenharia de Materiais, Universidade Federal da Paraíba – UFPB, João Pessoa, PB, Brazil 2 Departamento de Engenharia Elétrica, Universidade Federal da Paraíba – UFPB, João Pessoa, PB, Brazil 3 Departamento de Engenharia de Materiais, Universidade Federal de Campina Grande – UFCG, Campina Grande, PB, Brazil 1
*wellen.renate@gmail.com
Abstract Compounds of poly(3-hydroxybutyrate) (PHB) and titanium dioxide (TiO2) with filler content between 1% and 10% were prepared in a laboratory internal mixer. The effect of heating and cooling rates on the crystallization and melting of PHB/TiO2 compounds was investigated by differential scanning calorimetry (DSC). Melt and cold crystallization rates rise with increasing cooling/heating rates. A higher cooling rate translates to a lower melt crystallization temperature, while a higher heating rate results in a higher cold crystallization temperature. TiO2 promotes melt crystallization of PHB, behaving as a nucleant agent. The total crystallinity developed after melt and cold crystallization decreases for low levels of TiO2, i.e. 2% per weight, and is almost independent of the heating/cooling rate. The melting temperatures and rates are minimally affected by both the heating rate and filler content. The results suggest that the desired PHB microstructure can be controlled by filler content and adjusted heating/cooling rate. Keywords: crystallization, DSC, melting, PHB, titanium dioxide.
1. Introduction Polymers are fundamental for most common materials of our modern society because they present several desired features like lightness, easy processability, softness and low cost. As polymers are easily suitable for multiple uses, companies have been using them increasingly in various areas, such as packaging, automotive, medical-hospital, construction, electro-electronics industries among many others. Nevertheless, the production of the most useful polymers is based on fossil fuels. Considered to be non-biodegradable, they need decades to degrade and are an important source of pollution; in particular situations ecological disasters are verifiable. Concerned with the ecological equilibrium in nature, the society has made a lot of efforts to preserve the environment and to improve it where possible. Engaged with this issue, polymer researchers are seeking answers on how to decrease the use of petroleum based products, thus reducing pollution and toxic volatiles emission. The use of natural and biodegradable polymers appears as good option. Among these polymers, the polyhydroxyalkanoates (PHAs) seem to be a good choice. PHAs are linear aliphatic polyesters produced in nature by bacterial fermentation of sugar or lipids. They are used by the microorganisms to store carbon and energy. Many different monomers can be combined within this family to give materials with different properties. In the present work poly(3-hydroxybutyrate) (PHB) was selected for studying[1-6]. PHB possesses good oxygen permeability and ultra-violet resistance, it is nontoxic and has been used in the packing industry for years; it is biocompatible and hence suitable for
Polímeros, 28(2), 161-168, 2018
medical applications; due to its good mechanical properties and fast biodegradation rates, it has been employed in restaurant disposables, i.e., cups and cutleries. Furthermore, a wide range of additives and polymers may be incorporated to PHB, producing blends and compounds, to improve its performance and to expand its possible applications[7-12]. Few articles are concerned with PHB crystallization and melting behaviour in the presence of fillers. For instance carbon black (CB) and babassu natural fibers induce PHB partial crystallization from the melt during cooling and partial cold crystallization on reheating[13-15]. The amount of polymer crystallization in each stage depends strongly on the cooling rate and the filler content, being the melting subtly affected by the fillers and experiment rates. In a recent study concerned with the influence of Zinc oxide (ZnO) on the PHB crystallization[16] it was found, that ZnO can neither be classified as a crystallization accelerator, nor as a crystallization inhibitor; the influence of zinc oxide on PHB crystallization is irregular and strongly dependent on its concentration, the melting of PHB/ZnO compounds was little modified by ZnO content and heating/cooling rates tested. To the best of our knowledge, there is not any study in the literature dealing with the melting as well as with the crystallization of PHB/TiO2. From the literature is known that TiO2 could greatly improve the total solar reflectance, thus an appropriate content of TiO2 filler counteracts the thermal and photodegradation of plastics exposed to the sun. Antibacterial activity of TiO2 has been presented in the
161/168 161
O O O O O O O O O O O O O O O O
Jaques, N. G., Silva, I. D. S., Barbosa Neto, M. C., Ries, A., Canedo, E. L., & Wellen, R. M. R. literature; the packing, textile and health-care industries have successfully used TiO2 in their products with the objective to reduce the risk of microorganism transmission. Adding to the biodegradable nature of PHB this research is dealing with a very promising compound[17-26]. The non-isothermal melt and cold crystallization, and melting processes of PHB/TiO2 nanocompounds were investigated by differential scanning calorimetry (DSC) applying heating/cooling/reheating cycles (six different heating/cooling rates); thermal parameters related to these processes are presented with the main objective to provide a reference of crystallization and melting parameters for these compounds.
2. Materials and Methods 2.1 Materials Poly(3-hydroxybutyrate) (PHB), a random copolymer with approximately 4% of 3-hydroxyvalerate units, was supplied by PHB Industrial SA (Brazil). Its characteristic thermal transition temperatures are presented in Table 1. Melt crystallization (Tmc), cold crystallization (Tcc) and melting peak (Tmp) temperatures were obtained during cooling and reheating, according to ASTM D-3418 standard, applying a cooling/heating rate of 10°C/min. Titanium dioxide (TiO2), was purchased from Evonik Degussa Co. (manufacturer’s specification P25), with a surface area of 50 m2/g and a 75:25 anatase:rutile ratio. According to the manufacturer, the mean crystal sizes of the anatase and rutile phases are approximately 25 and 94 nm, respectively.
2.2 Methods PHB and TiO2 were used as received. The components were mixed for 10 minutes using a Haake Rheomix 600 laboratory internal mixer fitted with high intensity (roller type) rotors with the chamber wall kept at 190˚C and a nominal rotor speed of 60 rpm. The TiO2 content ranged from 1% to 10% by weight. Scanning electron microscopy images were acquired with a LEO 1430 unit, from Zeiss. The polymer samples were fractured in liquid nitrogen to avoid plastic deformation and coated with a carbon layer in order to avoid the accumulation of charges. Thermal analysis was performed in a TA Instruments DSC Q20 V24.9 differential scanning calorimeter, under a nitrogen flow of 50 mL/min to minimize oxidative degradation, to which the PHB is susceptible[27-29]. Samples of approximately 5 mg were tested in aluminiun pans. A new specimen was used for each run. A blank curve was obtained for each heating/cooling/reheating stage to ensure that no contamination of the instrument had taken place. A thermal cycle of four stages was used: (1) heating from 25˚C to 190˚C (first heating stage); (2) isothermal stage: the samples were held at 190°C for 3 min to eliminate any residual crystallinity and erase the previous thermal history; Table 1. Thermal transition temperatures (°C) of PHB obtained by the authors. Tg
Tmc
Tcc
Tmp
2
71.9
54.3
172.6
162 162/168
(3) the melt was cooled to 20˚C (cooling stage) and then (4) reheated to 190˚C (second heating or reheating stage). Tests were conducted at constant heating and cooling rates of 5, 7.5, 10, 15, 20 and 30°C/min. Figure 1 shows a typical DSC output with the indicated cycles. Four thermal events were identified in most DSC tests. They were denominated as follows: F1: melting during the first heating stage; C1: melt crystallization during cooling; C2: cold crystallization during reheating, and finally F2: a second melting event. For each thermal event, the starting and end points of departure from the underlying baseline were visually established in a plot of energy flow (J) versus time (t). The fractional crystallization (or melting) x for the event was computed as a function of time by integration[30]: = x(t )
1 t ∫ J (t ′) − J 0 (t ′) dt ′ (1) E0 t1
Where: J0 is the virtual baseline during the event (a straight line in the present case), and E0 is the total latent heat of the phase change: = E0 ∫tt2 J (t ) − J 0 (t ) dt (2) 1
and t1 and t2 are the initial and final times. The rate of the phase change (crystallization or melting) c is: c(= t)
dx = dt
J (t ) − J 0 (t ) E0
(3)
from which the peak (maximum) and average crystallization rates may be computed. The fractional crystallization/melting x and the rate of crystallization/fusion c may be expressed as functions of temperature (T) by knowing the linear relationship between time and temperature during the event: T= T1 + φ(t − t1 ) (4)
Where: T1 is the sample temperature at the starting point t1, and ϕ is the (constant) rate of heating or cooling during the event. The specific latent heat of crystallization or melting
Figure 1. Typical DSC output for PHB/1%TiO2 heating/cooling/ reheating at 10°C/min (exothermic peaks up), showing the phase change events: first melting (F1), melt crystallization (C1), cold crystallization (C2), and second melting (F2). Polímeros, 28(2), 161-168, 2018
Effect of heat cycling on melting and crystallization of PHB/TiO2 compounds (or enthalpy, because the phase change occurs at constant pressure) is computed from E0, the polymer fraction wP and the sample mass mS:
The mass crystallinity change ∆XC during the event is estimated, taking into account the heat of fusion of PHB 100% crystalline:
E ∆H = 0 (5) wP mS
∆H ∆X c = 0 (6) ∆H m
A value of ∆H0M =146 J/g at the equilibrium melting temperature T0M = 185°C was previously reported in the literature[31].
3. Results and Discussions Figure 2 presents a SEM image of PHB/10%TiO2; despite the fact that TiO2 nanoparticles were not treated with a surface modifier, SEM image confirms their homogeneous distribution in the PHB matrix.
Figure 2. Scanning electron micrograph of PHB/10%TiO2, as obtained after mixing.
Figure 3 displays the DSC scans for neat PHB and all PHB-TiO2 compounds for a fixed heating/cooling rate (10°C/min). Four events are recognized: the first melting (A), melt crystallization (B), cold crystallization (C) and the second melting (D).
Figure 3. DSC scans of neat PHB and PHB-TiO2 compounds obtained during the first heating (A), melt crystallization (B), cold crystallization (C) and second melting stages (D) (constant 10°C/min scanning rate). Polímeros, 28(2), 161-168, 2018
163/168 163
Jaques, N. G., Silva, I. D. S., Barbosa Neto, M. C., Ries, A., Canedo, E. L., & Wellen, R. M. R. As can be seen from Figure 3, the TiO2 filler has no significant influence on the first melting event (A). Melt crystallization and cold crystallization are the most affected events by the filler. As TiO2 promotes melt crystallization in a similar way as a nucleant agent, the cold crystallization events vanish, as the crystallization process has been more completed during the melt crystallization event. The higher the filler content, the smaller is the observed cold crystallization peak. Results gathered from previous works performed with neat PHB, PHB/Babassu, PHB/Carbon black and PHB/ZnO compounds about the influence of heating and cooling rates on PHB phase transition are very similar to these presented in this paper, in general, it is observed that PHB partially crystallizes from the melt during the cooling and partially cold crystallizes on reheating, and that the relative amount of polymer crystallizing in each stage strongly depends on the cooling rate and filler content. Regarding the influence of the fillers, addition of babassu filler in higher concentrations (10 to 50 wt%) increased the crystallization peak temperature and crystallization rate; carbon black acted as a nucleating agent during PHB crystallization. The effect of ZnO on the PHB crystallization is unclear and strongly dependent on its concentration, behaving as an accelerator or an inhibitor upon changing of filler content[13-16]. The second melting events (D) are different from the first melting events (A), a fact that can be seen from the shape of the peaks. Note that the first heating refers to the melting of PHB as obtained some time after processing (fully crystalline phase, no cold crystallization could be observed prior to first melting), whereas the second heating refers to the melting of PHB obtained during the more controlled cooling stage of the experiment. In this stage, the samples had undergone melt crystallization, followed by cold crystallization. During reheating stage, i.e., for the second fusion event (D) all samples showed complex (double) endothermic peaks upon addition of several filler contents: a first small melting peak followed by a larger peak at a higher temperature; this first small peak can also manifest as a shoulder of the larger peak. This behavior is more pronounced when the TiO2 content is high, and the heating rate is lowest (results not shown). According to the literature[32-42], such double peaks appearing in the fusion region may originate from (1) melting, recrystallization and re-melting during heating, (2) the presence of more than one crystal modification (polymorphism), (3) different morphologies (lamellar thickness, distribution, perfection or stability), (4) physical aging or/and relaxation of the rigid amorphous fraction, or (5) different molecular weight species. Although PHB degradation has been reported in some papers, for instance, PHB/Bentonite Organoclay compounds[43], in this work the degradation of PHB, as investigated by thermogravimetry (TG) (results not shown), was verified occurring at temperatures higher than 270oC, that is, well above the melting range observed in Figure 3. Figure 4 presents the development of the relative crystallinity as function of temperature for neat PHB and PHB-TiO2 compounds; the data were obtained from DSC thermal scans. All plots have the sigmoidal shape characteristic of phase transformation in polymers. 164 164/168
One may observe the displacement of the sigmoid plots in Figure 4 to higher temperatures with increasing TiO2 concentration.
For a given cooling/heating rate and TiO2 content, the melt crystallization peak temperature is higher than the cold crystallization peak temperature, as shown in Figures 5 and 6. Numerical values are given in Table 2. The melt crystallization peak temperature (Figure 5) decreases with an increase in the cooling rate and increases with an increase in the filler content. The melt crystallization peak rates (Figure 7) are proportional to the cooling rate and strongly dependent on the TiO2 content. The cold crystallization temperature (Figure 6) and peak rate (Figure 8) increase with an increase in the heating rate when TiO2 filler concentration
Figure 4. Relative crystallinity versus temperature for neat PHB and PHB-TiO2 compounds considering melt crystallization at cooling rate 10°C/min.
Figure 5. Crystallization peak temperature for melt crystallization of neat PHB and PHB-TiO2 compounds as a function of filler content and cooling rate. Polímeros, 28(2), 161-168, 2018
Effect of heat cycling on melting and crystallization of PHB/TiO2 compounds Table 2. Melt and cold crystallization parameters of compounds - Neat PHB. Neat PHB Melt Crystallization
Cold Crystallization
ϕ
Tmp
cmax
ΔHm
ΔXc
Tmp
cmax
ΔHm
ΔXc
(°C/min)
(°C)
(min−1)
(J/g)
(%)
(°C)
(min−1)
(J/g)
(%)
5
77.2
0.26
44.01
30.1
--
--
--
--
7.5
69.6
0.25
27.81
19.1
50.8
1.06
11.45
7.8
10
71.9
0.31
17.74
12.2
54.3
1.28
21.61
14.8
15
64.0
0.54
5.21
3.6
64.0
1.52
38.66
26.5
20
--
--
--
--
71.8
1.68
53.56
36.7
30
--
--
--
--
81.2
1.98
28.20
19.3
PHB + 1%TiO2 Melt Crystallization
Cold Crystallization
ϕ
Tmp
cmax
ΔHm
ΔXc
Tmp
cmax
ΔHm
ΔXc
(°C/min)
(°C)
(min−1)
(J/g)
(%)
(°C)
(min−1)
(J/g)
(%)
5
84.3
0.27
31.30
21.4
--
--
--
--
7.5
68.8
0.30
26.78
18.3
46.9
1.13
4.21
2.9
10
65.6
0.42
18.81
12.9
49.1
1.49
9.17
15
63.1
0.57
7.72
5.3
56.3
1.82
34.0
23.3
20
--
--
--
--
56.3
1.94
31.04
21.3
30
--
--
--
--
78.0
2.12
51.95
35.6
6.3
PHB + 2%TiO2 Melt Crystallization
Cold Crystallization
ϕ
Tmp
cmax
ΔHm
ΔXc
Tmp
cmax
ΔHm
ΔXc
(°C/min)
(°C)
(min−1)
(J/g)
(%)
(°C)
(min−1)
(J/g)
(%)
5
79.2
0.36
35.99
24.7
--
--
--
--
7.5
76.9
0.43
30.54
20.9
--
--
--
--
10
70.3
0.47
28.49
19.5
--
--
--
15
65.3
0.56
20.15
13.8
53.5
1.86
11.90
20
--
--
--
--
59.6
2.24
27.08
18.5
30
--
--
--
--
78.2
1.82
36.92
25.3
-8.15
PHB + 4%TiO2 Melt Crystallization
Cold Crystallization
ϕ
Tmp
cmax
ΔHm
ΔXc
Tmp
cmax
ΔHm
ΔXc
(°C/min)
(°C)
(min−1)
(J/g)
(%)
(°C)
(min−1)
(J/g)
(%)
5
86.3
0.43
29.69
20.3
--
--
--
--
7.5
83.2
0.56
31.12
21.3
--
--
--
--
10
77.8
0.68
28.61
19.6
--
--
--
--
15
69.7
0.71
24.12
16.5
52.9
2.69
0.67
0.5
20
66.4
0.86
26.42
18.1
55.0
2.52
6.24
4.3
30
--
--
--
--
72.5
1.95
15.16
10.4
PHB + 10%TiO2 Melt Crystallization
Cold Crystallization
ϕ
Tmp
cmax
ΔHm
ΔXc
Tmp
cmax
ΔHm
ΔXc
(°C/min)
(°C)
(min−1)
(J/g)
(%)
(°C)
(min−1)
(J/g)
(%)
5
91.2
0.58
36.52
25.0
--
--
--
--
7.5
87.9
0.78
37.75
25.9
--
--
--
--
10
83.7
0.92
35.56
24.4
--
--
--
--
15
78.3
1.05
30.82
21.1
--
--
--
--
20
72.9
1.21
31.54
21.6
--
--
--
--
30
--
--
--
--
--
--
--
--
Polímeros, 28(2), 161-168, 2018
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Jaques, N. G., Silva, I. D. S., Barbosa Neto, M. C., Ries, A., Canedo, E. L., & Wellen, R. M. R.
Figure 6. Crystallization peak temperature for cold crystallization of neat PHB and PHB-TiO2 compounds as a function of filler content and heating rate.
Figure 8. Crystallization maximum rate for cold crystallization of neat PHB and PHB/TiO2 compounds as a function of the filler content.
Figure 7. Crystallization maximum rate for melt crystallization of neat PHB and PHB/TiO2 compounds as a function of the filler content.
Figure 9. Total crystallinity developed during melt and cold crystallization as function of filler content for several cooling/ heating rates.
is low (up to 1%); their full dependence on filler content is unclear, but in general, the temperature decreases with an increase in the TiO2 content.
drop in crystallinity for the highest heating/cooling rate might be explained with a suppressed crystallization under the effect of fast cooling and heating process.
Figure 7 and 8 show the melt and cold crystallization peak rates which increase significantly with increasing filler concentration. This allows to classify titanium dioxide as a crystallization accelerator. Figure 9 presents the results for the degree of crystallinity, which was computed using Equation 6. Initially, there is a decrease in the crystallinity for PHB compounds with TiO2 content up to 2%. For higher filler concentrations no trend is observable; results are independent of the heating/cooling rate as well as of the filler content, even considering the high uncertainty of DSC estimates of crystallinity[44]. The observed 166 166/168
4. Conclusions The cooling and heating rates have a strong influence on the melt and cold crystallization of PHB and its compounds with titanium dioxide. Increasing the cooling rate results in lower melt crystallization temperatures and promotes higher crystallization rates. Both, cold crystallization rates and cold crystallization peak temperatures increase with the heating rate. The addition of titanium dioxide accelerated the melt crystallization, suggesting TiO2 as a nucleant agent for PHB. The total crystallinity as obtained after melt and successive Polímeros, 28(2), 161-168, 2018
Effect of heat cycling on melting and crystallization of PHB/TiO2 compounds cold crystallization, decreases for low levels of TiO2, i.e., less than 2% per weight, and it is quite independent of the heating/cooling rate after that. The melting temperatures and rates are minimally affected by both the heating rate and filler content. Summing up, controlling the cooling and heating rates as well as adding TiO2 makes PHB crystallization changeable, and consequently its processing parameters, thus a cost-effective processing may be tailored.
5. Acknowledgements Authors would like to thank PHB Industrial SA (Brazil) for kindly supplying PHB resin. NGJ, IDSS and MCBN thank CNPq for their scholarships.
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Polímeros, 28(2), 161-168, 2018
ISSN 1678-5169 (Online)
http://dx.doi.org/10.1590/0104-1428.03117
The effect of molecular weight and hydrolysis degree of poly(vinyl alcohol)(PVA) on the thermal and mechanical properties of poly(lactic acid)/PVA blends Iván Restrepo1,2, Carlos Medina3, Viviana Meruane4, Ali Akbari-Fakhrabadi4, Paulo Flores3* and Saddys Rodríguez-Llamazares5 Departmento de Ingeniería de Materiales, Universidad de Concepción, Concepción, Chile 2 Unidad de Desarrollo Tecnológico, Universidad de Concepción, Coronel, Chile 3 Departamento de Ingeniería Mecánica, Universidad de Concepción, Concepción, Chile 4 Laboratorio de Materiales Avanzados, Departamento de Ingeniería Mecánica, Universidad de Chile, Santiago, Chile 5 Centro de Investigación de Polímeros Avanzados, Concepción, Chile 1
*pfloresv@udec.cl
Abstract The effect of molecular weights and hydrolysis degrees (HD) of polyvinyl alcohol (PVA) on thermal and mechanical properties and crystallinity of polylactic acid (PLA)/PVA blends was investigated. Blends were prepared by the melt blending method using PLA/PVA ratios: 80/20, 90/10 and 97/3 wt. %. A single glass transition temperatures was observed for all PLA/PVA blends, suggesting the formation of binary compatible blends at concentration range studied. Thermogravimetric analysis results showed a better thermal stability for PLA/PVA blends containing PVA of higher Mw and HD. According to mechanical properties, low quantities of PVA (3 wt. %) do not affect the tensile strength of blends (irrespective of Mw and HD). However, as the PVA content increases, tensile strength tends to lower values, especially for blends with 20 wt.% of PVA, with 98% of HD. Keywords: material testing, melt blending, polylactic acid, polyvinyl alcohol, polymer blend.
1. Introduction Polylactic acid (PLA) is a linear aliphatic polyester obtained from renewable sources such as starch and sugar. It is one of the most widely used bioplastics due to its optical, mechanical and barrier properties, and good processability by conventional transformation techniques of thermoplastics[1]. The worldwide production capacity of PLA in 2021 is estimated around 297.000 tons, 40% more than that produced in 2016[2]. PLA is classified by the American food and drug administration as generally recognized as safe and hence it is very common in the food packaging industry. Its barrier, thermal and mechanical properties are in fact similar to those of synthetic polymers such as polystyrene[3-6]. PLA can be biodegraded under compost conditions; it takes up to a year to degrade in real and simulated soil burial conditions. This slow degradation compared to other polymers such as for example poly (hydroxy alkanoates), is due to the fact that PLA must be hydrolyzed before microorganisms can use it as a source of nutrients[7,8]. Adding small quantities of compatible hydrophilic polymers such as PVA, which is biodegradable, hydrophilic and flexible[9] is a way to enhance the biodegradability of PLA. For example, it has been reported that polyvinyl alcohol (PVA) has the ability to accelerate the degradation of PLA, by increasing the hydrophilicity of the blend and breaking the crystallinity of PLA[10]. Furthermore, the hydroxyl groups in PVA readily form hydrogen bonds
Polímeros, 28(2), 169-177, 2018
with the ester groups of PLA, which favors the compatibility of their blends. Yeh et al., (2008)[11] found that a compatible blend PLA: PVA (80: 20 wt. %: wt. %) prepared by melt blending method using PLA Mw = 37 kg/mol and PVA with HD = 97-98.5% and Mw = 75 kg/mol. Similar results were reported by Shuai et al., 2001[12] for blends PLA/PVA 10 to 90 wt. % of PVA with polymerization degree of 2000 and HD=99%. They report the formation of interpolymer hydrogen bonds in the amorphous region of PLA/PVA blends (PVA content higher than 50 wt. %), which contributes to compatibility of these blends. However, PVA and PLA crystallized as isolated phases, for that tensile strength and elongation at break first declines for PVA content from 0 to 50 wt. % and then increases with increasing when the PVA content. Several authors have reported the effect of PVA content with different HD and Mw on mechanical properties of PLA/PVA blends. For example, Lipsa et al. (2008)[13] prepared films by these blends using PVA with HD = 98% and Mw = 18 kg/mol. They found that the blends are partially compatible for PVA content from 70 to 90 wt. %, and slight reduction of mechanical properties, such as lower tensile strength were observed in blends. Tsuji and Muramatsu (2001)[14] found that the Young’s modulus, tensile strength and percent elongation at break
169/177 169
O O O O O O O O O O O O O O O O
Restrepo, I., Medina, C., Meruane, V., Akbari-Fakhrabadi, A., Flores, P., & Rodríguez-Llamazares, S. of the blends increase with increasing PVA content from 50 to 90 wt. % (HD=99.5%). Hoai N. et al. (2014)[15] prepared nanofibrillary structures from PLA/PVA blends, varying content of PVA from 0 to 100 wt. %. They found that increasing the amount of PVA results the thermal degradation of blends more stable. However, a systematic study that show the influence of content, molecular weight and hydrolysis degree of PVA on compatibility, and therefore on the thermal and mechanical properties and crystallinity of PLA has not found. In this work, PLA/PVA blends prepared by the melt blending method was studied.
2. Materials and Methods 2.1 Materials The commercial PLA grade 3251D, Mw of 55.4 kg/mol and isomer D lactic acid content of 1.2%[16] was purchased from Nature Works®, United States. Four different types of PVA with different hydrolysis degree and molecular weight were provided by Sigma Aldrich. The main specifications of PVA are summarized in Table 1. The pure PLA processed under similar conditions was studied as reference material.
temperature (Tcc), cold crystallization enthalpy (ΔHcc), melting temperature (Tm), melting enthalpy (ΔHm), and degree of crystallinity (XC) were determined from the second heating scans. XC of PLA in the blends was calculated by the Equation (1): Xc =(∆H m / ( ∆H °m W ))100% (1)
where ΔHM is the melting enthalpy of the blends, W is the weight fraction of PLA and ΔH°M is the melting enthalpy of 100% crystalline PLA (93 J/g)[17,18]. 2.3.2 TGA Thermal stability of pure PLA, PVA’s and their blends was evaluated using a thermogravimetric analyzer Netzsch TGA 209 F3 (Tarsus, Selb, Germany). TGA scans were carried out at 10 °C/min under nitrogen atmosphere (20 ml/min), from 35 to 600 °C. The following parameters were reported: (i) temperature at maximum decomposition rate for each step of decomposition (Tmax), (ii) weight loss associated with Tmax (WL Tmax), (iii) the onset temperature (Tonset) corresponding to each decomposition step and (iv) the temperature corresponding to 5% weight loss (T5%).
2.2 Processing
2.3.3 XRD
The PLA/PVA blends, ratios of 80/20, 90/10 and 97/3 wt. % for each type of PVA (Table 1), were prepared via melt blending in a Brabender mixer (Plastograph ® EC plus, Mixer 50EHT32, Germany) at 60 rpm for 8 min. The temperature was set at 190 °C. The PLA was previously dried at 40 °C overnight in a vacuum oven. After blending, the samples were pressure-molded at 30 bar and 10 min in order to obtain plates for XRD and impulse excitation technique (IET) tests. The temperature was set between 170 and 210 °C, depending on the type of PVA used in the preparation of the sample.
The XRD measurements were recorded on a Bruker diffractometer Endeavor model D4/MAX-B (United States) operated at 40 kV, 20 mA, at room temperature, using a CuKα source and λ= 1.5405 Å. The diffraction spectra were taken in the range 3°<2θ<45° at 0.02° steps and a scanning rate of 1°/min. The pure PLA and PLA/PVA blends were previously annealed for 20 min at 120 °C under vacuum conditions.
2.3 Blend characterization 2.3.1 DSC The thermal properties of PLA/PVA blends were evaluated using a differential scanning calorimetry (DSC) analyzer (Netzsch DSC 204 F1 Phoenix), Germany. DSC curves were scanned at 10 °C/min under nitrogen atmosphere (20 mL/min). The DSC scans were performed from room temperature to 250 °C. Then, the samples were cooled to 25 °C and a second heating scans were made up to 250 °C. Glass transition temperature (Tg), cold crystallization Table 1. Nomenclature of blends. PVA Code
PVA
Range of HD Mw (g/mol) (%) 1PVA 13000-23000 87-89 2PVA 13000-23000
98
3PVA 31000-50000 87-89 4PVA 31000-50000
170 170/177
98
Sample code for PLA/PVA blends 80/20 90/10 97/3 wt.% wt.% wt.% 1PVA 1PVA 1PVA 80/20 90/10 97/3 2PVA 2PVA 2PVA 80/20 90/10 97/3 3PVA 3PVA 3PVA 80/20 90/10 97/3 4PVA 4PVA 4PVA 80/20 90/10 97/3
2.3.4 Tensile test The mechanical properties of PLA/PVA blends were evaluated using tensile test performed on a Karg Industrietechnik machine (Germany) according to ASTM D638 standards. Samples type V were fabricated using a mini injection equipment type Haake Minijet II, and injection conditions were 350 bar of pressure and injection temperature of 180 °C. Prior testing, specimens were conditioned under 25°C and 50% relative humidity (R.H) for 7 days. The crosshead speed was set at 10 mm/min. The Young’s modulus, strength and elongation percentage at break were obtained from the stress–strain curves. At least 5 individual measurements were carried out for each formulation and mean values and standard deviations were reported. Collected data were evaluated with a one-way analysis of variance at the 95% confidence level. For comparison purpose, elastic modulus of PLA/PVA blends were evaluated by Impulse excitation technique (IET). The rectangular samples were suspended by soft and tiny tape sticks to simulate a “free-free” boundary condition and were excited by an impact hammer. The samples have an aspect ratio of a/b=1.5, where a, b are the plate’s length and width dimensions, respectively. The vibration of the samples is captured by a microphone which was connected to a data acquisition system. A signal processing software computed the frequency content of the measured signals from which the experimental resonant frequencies were identified. Once the experimental resonant frequencies Polímeros, 28(2), 169-177, 2018
The effect of molecular weight and hydrolysis degree of poly(vinyl alcohol)(PVA) on the thermal and mechanical properties of poly(lactic acid)/PVA blends were identified, the elastic modulus was computed using the relationships presented in Equation (2)[19]: ω=
21.603h a
2
(
E
)
12ρ 1 − ν 2 (2)
where ʋ = 0.4 is Poisson´s ratio, ρ is the density, E is the elastic modulus and h=0.85 and a=20 are plate’s dimensions (thickness and length respectively, in millimetres). The variable ω corresponds to the second resonant frequency of the plate, which was easier to be excited and measured. 2.3.5 SEM The morphology of cryogenic-fractured cross-sections of PLA/PVA blends was analyzed by scanning electron microscopy (JEOL JSM 6380 LV, Tokyo-Japan), operated at 20 and 5 kV. The samples were fractured under liquid
nitrogen and sputtered with a gold coating of ca. 50 nm. The magnifications were 100x and 500x.
3. Results and Discussion 3.1 DSC Analysis DSC thermograms of PLA, PVA’s and PLA/PVA blends are shown in Figure 1. The thermograms for pure PLA and PLA/PVA blends display successive peaks corresponding to the glass transition, cold crystallization and melting. The Tg, Tcc and Tm are resumed in Table 2. The only one peak of Tg is observed for PLA/PVA blends, which is ranged from 42 to 55 °C, intermediate between the Tg values of pure PLA and PVA. This behavior is characteristic of compatible blends. Shuai et al., 2001[12] found by 13C solid-state nuclear magnetic resonance spectroscopy, the presence of hydrogen bonds between hydroxyl groups of PVA and carbonyl groups of PLA in the amorphous region of blends with
Figure 1. Second heating run DSC thermograms for: (a) Pure PLA and PVA’s; (b) PLA/PVA blends.
Table 2. Thermal characteristics of pure PLA, PVA’s and PLA/PVA blends. Sample code Pure PLA 1PVA 1PVA 80/20 1PVA 90/10 1PVA 97/3 2PVA 2PVA 80/20 2PVA 90/10 2PVA 97/3 3PVA 3PVA 80/20 3PVA 90/10 3PVA 97/3 4PVA 4PVA 80/20 4PVA 90/10 4PVA 97/3
Tg (°C) 61 37 42 43 50 40 44 42 53 44 55 43 50 42 51 42 54
Polímeros, 28(2), 169-177, 2018
Tcc (°C) 98 ___ 96 91 95 ___ 99 91 97 ___ 89 92 94 ___ 91 88 94
Tm1 (°C) 150 167 116 114 142 210 118 118 143 167 111 121 143 212 112 117 149
Tm2 (°C) 169 ___ 124 133 153 ___ 130 136 157 ___ 132 138 157 ___ 130 136 159
ΔHcc (J.g-1) 38 ___ 13 24 39 ___ 13 28 41 ___ 16 27 45 ___ 14 32 40
ΔHm (J.g-1) 54 20 14 22 41 66 14 25 49 21 15 23 52 61 16 27 47
Xc (%) 58 ___ 19 26 45 ___ 19 30 54 ___ 20 27 57 22 32 53
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Restrepo, I., Medina, C., Meruane, V., Akbari-Fakhrabadi, A., Flores, P., & Rodríguez-Llamazares, S. PVA content of 30 wt. %. A single Tg of PLA/PVA blends with PVA content lower than 20 wt. % is attributed to this kind of interpolymer interaction. Tg peak is shifted to lower temperatures with increasing PVA content in the blends. The blends and pure PLA exhibited two endothermic peaks of melting characteristic of the formation of two PLA crystal structures during its cold crystallization, known as α and β forms with melting temperatures Tm2 and Tm1 respectively. The melting temperature of the α-form is higher because of the better quality and higher size of its crystals[20,21]. The melting temperatures of the blends are also shifted toward lower temperatures, with increasing the PVA content irrespective of Mw and HD. The Xc values for PLA in blends decrease, especially when 20 wt.% of PVA was added. Thus, the crystallization of PLA in presence of PVA is affected due to the formation of imperfect crystals in PLA by considering that the presence of the PVA can cause a dilution of the PLA nuclei, combined with the possible interference between the chains of PLA and PVA due to partial compatibility between them. Opposite results were obtained by Yeh et al., (2008) [11] who found that highest PLA crystallinity, appears in PLA/PVA blends containing 20 wt. % of PVA (MW: 75KDa, HD: 98%), attributed to the interactions between PLA and PVA that promote the crystallization. However, Tsuji and Muramatsu (2001)[14] report a zero Xc value of PLA for all PLA/PVA blends (ratio 90:10 to 50:50 wt. %) indicating that PLA is amorphous in the blend. On the other hand, Shuai et al., (2001)[12] reported a decrease in Xc values of PLA when PVA is added in contents more than 70 wt.%, result of some depression of the PLA crystallinity upon blending with a large amount of PVA. Variations in values of Tcc, Tm, and Xc in PLA/PVA blends are usually attributed to interactions between the components. In general, the decrease in PLA crystallinity is related to the partially compatible nature of the blends[11,22].
3.2 TGA Analysis Thermal stability of pure PLA, PVA’s and PLA/PVA blends was evaluated by TGA analysis. Table 3 summarizes the thermal parameters. The TGA curve for pure PLA showed one-step degradation with a weight loss of 99%, associated with the loss of ester groups by unzipping depolymerization[23,24]. Tmax for pure PLA was 364 °C. In contrast, pure PVA’s displayed three degradation steps. The first one, between 31 and 188 °C and a weight loss below 5 wt. %, is associated with the loss of absorbed moisture. The second step, between 190 and 388 °C and a weight loss ranged from 65 to 77%, is related to the loss of low molecular weight substances, such as residual acetate groups, non-conjugated polyenes, acetic acid and H2O. The third step, between 351 and 426 °C, is associated to the breakdown of polymer backbone[25]. The main product of thermal degradation of PVA is water, which is formed by the elimination of hydroxyl side-groups[26]. For the same Mw, PVA with higher HD (98%) showed the lowest values of Tmax-1 in the second step of degradation with values around 260 °C and weight loss associated with this Tmax-1 was 69–71% (see Table 3). Therefore, PVA with HD of 87-89%, present better thermal stability. Acetate groups present in PVA with lower HD, confer higher thermal decomposition temperatures to PVA, favoring its thermal stability[27]. Two degradation steps were observed for PLA/PVA blends, (see Figure 2) suggesting that thermal decomposition of PVA and PLA happens as a combined process[25]. The first step of decomposition was between 188 and 386 °C with a weight loss of 75–98 wt.%. The second one was between 381 and 507 °C with a weight loss 0.2-2.2 wt. %. With increasing PVA content in the blends, the main degradation step is shifted significantly to lower temperatures. Water absorbed by PVA, could be favoring hydrolytic degradation of PLA.
Table 3. TGA and DTG parameters for pure PLA, PVA and PLA/PVA blends. Sample code Pure PLA 1PVA 1PVA 80/20 1PVA 90/10 1PVA 97/3 2PVA 2PVA 80/20 2PVA 90/10 2PVA 97/3 3PVA 3PVA 80/20 3PVA 90/10 3PVA 97/3 4PVA 4PVA 80/20 4PVA 90/10 4PVA 97/3
172 172/177
TGA T5%
T on set-1 (°C)
Tmax-1 (°C)
WL-1 (%)
DTG T On set-2 (°C)
Tmax-2 (°C)
WL-2 (%)
334 241 268 280 301 186 272 286 301 261 275 285 301 215 270 288 311
284 216 191 232 256 216 188 231 251 221 218 225 241 206 220 240 261
364 306 301 312 353 261 305 330 348 314 305 322 345 260 380 331 356
99 77 90 92 97 69 93 95 97 75 88 93 97 71 87 94 98
____ 391 381 392 411 351 383 406 411 391 388 395 406 351 390 393 421
____ 426 425 427 420 421 428 419 421 426 432 422 427 416 435 427 431
____ 12 7 5 1 15 4 2 1 16 7 3 1 14 4 3 1
Polímeros, 28(2), 169-177, 2018
The effect of molecular weight and hydrolysis degree of poly(vinyl alcohol)(PVA) on the thermal and mechanical properties of poly(lactic acid)/PVA blends The Tmax-1 shifts to higher values for those blends with PVA’s of higher Mw and higher HD. With increasing Mw of PVA, the entanglement along the chains of PLA and PVA is favored. These results are in agreement with some literature reports[28-30] for higher Mw of PVA, which favors intramolecular entanglement between PVA and PLA, primarily by virtue of esterification of PVA hydroxylic groups and PLA carboxylic groups, improving the thermal stability of PVA in the blend.
3.3 XRD analysis The XRD diffraction patterns for pure PLA, PVA and selected PLA/PVA blends are shown in Figure 3. The results are shown only for two types of blends, considering that the diffraction patterns for all PLA/PVA blends and pure PLA were very similar. The most intense peaks for PVA’s appeared around 2θ=19.3°, 20° and 22.7°, related to reflections of planes 10̅1, 101, and 200, respectively, associated with a monoclinic unit cell[14]. Assender and Windle, 1998[31] reported that
PVA chains are lying along the b-axis of the unit cell. Two characteristic peaks of pseudo orthorhombic α–phase crystallites structure (Space group P32) were observed for pure PLA, around 16.5 and 18.8° (2θ), corresponding to the reflection of 110/200 and 203 planes with two chains in a helical conformation[11,32,33]. The absence of characteristic PVA peaks in blends is attributed to the absence of PVA crystallization in the presence of PLA suggesting that PVA molecules was trapped in an amorphous state in the PLA phase and/or PLA molecules reduce the nucleus density of PVA crystallites[14]. In this context, Shuai et al. (2001)[12] reported the presence of two isolated crystalline phases (called co-crystalline phase) coexisting in PLA/PVA blends (ratio 1:1), that have the same crystal structure without inter-polymer interactions between them. Furthermore, with the addition of PVA in PLA, all of the blends show no shift in characteristic diffraction peaks of PLA, implying that, there is no significant effect of PVA on the crystal structure of PLA[11].
Figure 2. TGA diagrams for (a) Pure PLA, PVA’s; (b) blends PLA/PVA: 80/20 wt. % blends; (c) blends PLA/PVA: 97/3 wt. %.
Figure 3. Diffractions patterns for: (a) PLA/1PVA blends and (b) PLA/4PVA blends. Polímeros, 28(2), 169-177, 2018
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Restrepo, I., Medina, C., Meruane, V., Akbari-Fakhrabadi, A., Flores, P., & Rodríguez-Llamazares, S. 3.4 Mechanical analysis Tensile strength of pure PLA and PLA/PVA blends are shown in Figure 4. Compared with pure PLA, the tensile strength of blends containing 3 wt. % of PVA, is very similar (irrespective of Mw and HD). However, as the PVA content increases, tensile strength tends to lower values, especially for blends with 20 wt.% of PVA, with 98% of HD. Influence of Mw is observed for blends with 10 wt.% of PVA. Thereby, blends containing PVA with higher Mw showed higher values in the tensile strength, irrespective of the HD. Young’s modulus results from tensile test analysis for pure PLA and PLA/PVA blends are presented in Figure 5. As compared with pure PLA, it is observed a decrease in Young´s modulus with increasing the content of PVA, particularly for blends with 20 wt.% of PVA (HD: 98%), indicating that presence of PVA affect the rigidity of PLA. Influence of HD of PVA is clearer in Young´s modulus values obtained by IET analysis presented in Table 4. For higher HD, increasing content of PVA from 10 to 20 wt.%,
decrease Young´s modulus values, and for the same Mw, samples containing 20 wt.% PVA with higher HD, present lower young´s modulus. PVA present flexible characteristic, in this way is expected a decrease of rigidity in PLA/PVA blends, suggesting that acetate groups present in PVA with lower HD influence the mobility of polymeric chains in the blends. In this context, according to DSC results, the lowering of crystallinity of PLA in presence of PVA can lead a decrease in mechanical properties of PLA/PVA blends. Results reported for Shuai et al., 2001 and H. Tsuji, 2001[12,14], pointing out that the lowering of mechanical properties with increasing PVA content was attributed to the partial compatibility of the blends, due to the weak interfacial adhesion between the two phases in PLA/PVA blends. However, PVA with higher Mw conferring more resistance to PLA/PVA blends acts like a cross-linked component between molecules of PVA and PLA.
3.5 SEM Analysis Figure 6 shows cross-section SEM images of selected PLA/PVA blends. In general, is observed a porous morphology for all blends, related to the partially compatible nature of PLA/PVA blends. However, the content of PVA had a remarkable influence on blends morphology, as evidenced in the characteristics of the observed pores. In this context, blends with higher content of PVA have higher density of pores, indicating a greater phase separation, unlike blends containing less Table 4. Young´s modulus from IET for pure PLA and PLA/PVA blends. Sample code
Figure 4. Tensile strength for pure PLA and PLA/PVA blends.
Pure PLA PLA/1PVA PLA/2PVA PLA/3PVA PLA/4PVA
Young’s Modulus (GPa) 80/20 wt.% 90/10 wt.% 100 wt.% 4.7 4.3 4.1 2.1 4.0 4.1 4.0 3.0 4.2
Figure 5. Young’s modulus from tensile test for pure PLA and PLA/PVA blends. 174 174/177
Polímeros, 28(2), 169-177, 2018
The effect of molecular weight and hydrolysis degree of poly(vinyl alcohol)(PVA) on the thermal and mechanical properties of poly(lactic acid)/PVA blends
Figura 6. SEM images of selected PLA/PVA blends: (a) 3PVA 80/20; (b) 4PVA 80/20; (c) 3PVA 97/3; (d) 4PVA 97/3.
PVA (particularly with high Mw and independent of the HD), which have smaller pores, suggesting that the compatibility of PVA and PLA was improved on addition of low quantities of PVA with higher Mw. In this context, Zhang et al., (2012)[34] reported that increasing the PVA content (polymerization degree:1700, HD:88%) from 10 to 80 wt. % in PLA/PVA blends produced a less smooth surface. This was attributed to an excessive PVA content that led to strong and extensive intermolecular hydrogen bonding, which in turn resulted in PLA aggregation. Other reports showed development of pores with an average size of 5 µm for 50/50 PLA/PVA blends, suggesting the formation of two separate phases; a PLA-rich phase forming domains leading to porosity and a continuous PVA-rich phase[14].
4. Conclusions The partially compatible nature of PLA/PVA blends were evidenced from thermal and mechanical results. In this context, were observed one peak of Tg for PLA/PVA blends, intermediate between the Tg values of pure PLA and PVA, characteristic behavior of compatible blends. Additionally, for a higher content of PVA, decrease crystallinity of PLA, thus, PLA/PVA blends with 20 wt.% of PVA tend to reduce mechanical properties of PLA, particularly for those with 98% HD. However, the addition of low quantities of PVA (3 wt. %) Polímeros, 28(2), 169-177, 2018
do not reduce the tensile strength in PLA/PVA blends, because of better compatibility between PLA and PVA. PVA’s with higher Mw favors the entanglement of PVA and PLA improving the thermal stability of PLA/PVA blends. With increasing PVA content in the blends, the main degradation step is shifted significantly to lower temperatures, however PVA with HD of 87-89%, present better thermal stability due to acetate groups conferring higher thermal decomposition temperatures to PVA, favoring the thermal stability of the PLA/PVA blends.
5. Acknowledgements This work has been financed by projects of Comisión Nacional de Investigación Científica y Tecnológica CONICYT (Beca de Doctorado Nacional – Proyecto PAI 781411004), CONICYT-REGIONAL R08C1002 and Programa de Financiamiento Basal para Centros Científicos y Tecnológicos de Excelencia PFB-27. The authors thank to Carmen Pradenas for sample testing.
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Received: Apr. 28, 2017 Revised: June 14, 2017 Accepted: July 05, 2017
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ISSN 1678-5169 (Online)
http://dx.doi.org/10.1590/0104-1428.013116
O O O O O O O O O O O O O O O O
Polysaccharides of red alga Gracilaria intermedia: structure, antioxidant activity and rheological behavior Joana Paula Lima de Castro1, Luís Eduardo Castanheira Costa1, Maísa Pessoa Pinheiro1, Thiago dos Santos Francisco2, Pedro Hermano Menezes de Vasconcelos3, Lizandra Mistrello Funari4, Renata Moschini Daudt4, Gustavo Ramalho Cardoso dos Santos5, Nilo Sérgio Medeiros Cardozo4* and Ana Lúcia Ponte Freitas1 Departmento de Bioquímica e Biologia Molecular, Universidade Federal do Ceará – UFC, Fortaleza, CE, Brazil 2 Departmento de Química Orgânica e Inorgânica, Universidade Federal do Ceará – UFC, Fortaleza, CE, Brazil 3 Departmento de Química, Instituto Federal de Educação, Ciência e Tecnologia do Ceará, Fortaleza, CE, Brazil 4 Departmento de Engenharia Química, Universidade Federal do Rio Grande do Sul – UFRGS, Porto Alegre, RS, Brazil 5 Laboratório de Tecido Conjuntivo, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brazil 1
*nilo@enq.ufrgs.br
Abstract A sulfated polysaccharide fraction from the red alga Gracilaria intermedia (PLS) was obtained by papain digestion (60 °C, 30 min). The extract was subjected to colorimetry and turbidimetry analysis, Fourier transform infrared (FTIR) spectroscopy, 1H, 13C and 2D 1H COSY nuclear magnetic resonance (NMR) and gas chromatography/mass spectrometry analysis. Antioxidant activity tests were performed (chelation of ferrous ion, total antioxidant capacity, and scavenging of DPPH radicals); significant activity of the extract indicated that these polysaccharides may be used as non-synthetic antioxidants. The rheological behavior of aqueous polysaccharide solutions was studied at 25 ± 1 °C using steady-shear and dynamic oscillatory measurements. All the solutions analyzed showed pseudoplastic behavior and potential to act as a thickening agent, as proved through a preliminary comparison with a commercial product used for this application. Keywords: antioxidant activities, Gracilaria intermedia, polysaccharides, rheological behavior, structure.
1. Introduction Polysaccharides are polymers that show a wide range of properties depending on their monomeric composition and have applications in different types of industries. Applications include the use as a stabilizer, thickener, flocculant, and water-retaining agent in the textile, food, pharmaceutical, and biotechnology industries[1-4]. Carrageenans and agarans are sulfated polysaccharides obtained from red algae. The industrial applications of these polymers are closely dependent on their rheological properties. Such properties make these two polymers important gelling agents and thickeners, especially in the food industry[5]. Sulfated polysaccharides also exhibit various biological activities, including antioxidant activities, among others[6-9]. Therefore, they also have potential application as agents to reduce cellular damage and to prolong the shelf time of foods; these are areas where compounds with antioxidant activity are currently in great demand[10-12]. Non‑synthetic compounds, particularly polysaccharides, with antioxidant potential are currently under investigation by several research groups[13-18] as an alternative to synthetic antioxidant compounds traditionally used in the food and pharmaceutical industries such as butylated hydroxytoluene
178 178/186
(BHT), butylated hydroxyanisole (BHA), and tert-butyl hydroquinone (TBHQ) due to their suspected damage to liver tissue and carcinogenic potential[9,13,19]. Red algae from the genus Gracilaria are economically important in the phycocolloids industry[20] and are currently the source of about 65% of the 7.5 tons of agar produced annually throughout the world[21]. Sulfated polysaccharides obtained from species of the genus Gracilaria contain 3-linked-β-D-galactopyranose (G unit) and 4-linked3,6‑anhydro-α-L-galactopyranose (LA unit)[22-25] with substitution of hydroxyl groups by ester sulfate, methyl groups, and pyruvic acid[23-26] (Figure 1). Gracilaria is very common in the Northeast Brazilian coast and its extraction has been pointed as an alternative for the economic inclusion of part of the population of this region[27]. Yield of extraction and final properties of agarans depend strongly on the seaweed species, the methodology of extraction of polysaccharides, seasonal variations and the place of origin of the specimens. So the chemical, structural, and rheological characterization of polysaccharydes extracted from seaweeds is necessary[26,28]. In the specific case of the
Polímeros, 28(2), 178-186, 2018
Polysaccharides of red alga Gracilaria intermedia: structure, antioxidant activity and rheological behavior solution was then filtered through a nylon membrane, and the homogenate was retained. The polysaccharides in solution were precipitated with 16 mL of 10% cetylpyridinium chloride (CPC) solution. After 24 h at room temperature (25 °C), the mixture was centrifuged at 2,560 × g for 20 min at 20 °C. The polysaccharides were washed with 500 mL of 0.05% CPC solution, dissolved with 100 mL of a 2 mol.L-1 NaCl‑ethanol (100:15, v/v) mixture, and the excess of salts was removed by precipitation and wash with 200 mL of absolute ethanol. After 24 h at 4 °C, the precipitate was collected by centrifugation (2,560 × g for 20 min at 20 °C), washed extensively with ethanol-80%, then absolute ethanol. After this, the polysaccharides (PLS) were washed with acetone, which was followed by hot air drying (60 ºC) until all the acetone was removed.
2.3 Chemical analysis Figure 1. Chemical structure of agar type molecules with the different types of sugar units and substituents.
species Gracilaria intermedia, the related scientific research Gracilaria intermedia is still scarce. To the knowledge of the authors, only studies on its taxonomy[28-31] and ecological aspects[32] have been reported. In this context, the goal of this paper was to characterize sulfated polysaccharides extracted from the red seaweed Gracilaria intermedia. This characterization was performed by spectroscopy, turbidimetric and colorimetric methods. The antioxidant activity of the polymer and its rheological behavior in solution were also evaluated.
2. Materials and Methods
Sulfate content of PLS was determined by the barium‑gelatin method[34] after hydrolysis of the sample in 1 mol.L-1 HCl (5 h, 105 °C) using sodium sulfate (Na2SO4) as standard. To measure the amount of total neutral carbohydrates the phenol-sulfuric method was performed[35] with a standard curve prepared with D-galactose. The protein content was measured by the Bradford method[36] using bovine serum albumin (BSA) as standard.
2.4 Monosaccharide composition Samples of the polysaccharides extracted from G. intermedia (5 mg) were hydrolyzed with 5 mol.L-1 trifluoroacetic acid for 4 h at 100 °C, reduced with borohydride, and the alditols were acetylated with acetic anhydride:pyridine (1:1, v/v). The alditols acetates were dissolved in chloroform and analyzed in a gas–liquid chromatograph/mass spectrometrer (GCMS-QP2010 Shimadzu, Japan) with a DB-5ms column (Agilent)[37].
2.1 Materials
2.5 Infrared spectroscopy
Papain, cysteine and cetylpyridinium chloride (CPC) were obtained from Vetec (Brazil), ammonium molybdate was purchased from Dinâmica (Brazil) and D-galactose from Acros Organics (USA). Ferrozine, 1, 1-diphenyl2‑picrylhydrazyl (DPPH) and bovine serum albumin (BSA) were purchased from Sigma–Aldrich (USA). Ascorbic acid was obtained from Synth (Brazil) and Carboxymethyl cellulose was obtained from Mix (Brazil). All other solvents and chemicals were of analytical grade.
The Fourier transform infrared spectra (FTIR) were obtained with a Shimadzu IR spectrophotomer (FTLA 2000, ABB‑BOMEM, Canada) with measurements in the wavenumber range of 400-4000 cm-1 using 20 scans. The samples were analyzed as KBr pellets.
2.2 Extraction of sulfate polysaccharides Specimens of G. intermedia were collected in April 2013 on the Atlantic coast of Brazil (Taiba beach, São Gonçalo - Ceará). The collected seaweed were cleaned of epiphytes, washed with distilled water, and stored at -20 °C. A voucher specimen (n° 2386) was deposited in the phycological Herbarium of the Laboratory of Marine Sciences, Universidade Federal do Ceará, Brazil. The enzymatic extraction of polysaccharides was performed according to the methodology of Farias et al.[33], with some modifications. The dried tissue (5 g) was suspended in 250 mL of 0.1 mol.L-1 sodium acetate buffer (pH 5.0), containing 1 g of papain, 5×10-3 mol.L-1 EDTA, and 5×10-3 mol.L-1 cysteine, and incubated at 60 °C for 30 min. The incubation Polímeros, 28(2), 178-186, 2018
2.6 Nuclear magnetic resonance (NMR) spectroscopy The spectroscopic technique of nuclear magnetic resonance (NMR) is a method commonly used for structural characterization of seaweed polysaccharides[38-41]. 1 H, 13C and 2D 1H COSY NMR spectra in D2O were recorded at 353 K on a Fourier transform spectrometer (Bruker Avance DRX 500, USA) with an inverse multinuclear gradient probe-head equipped with z-shielded gradient coils.
2.7 Determination of antioxidant activity 2.7.1 Total antioxidant capacity The total antioxidant capacity test, based on the reduction of Mo6+ to Mo5+, was performed by the methodology proposed by Prieto, Pineda & Aguilar[42]. Aliquots of the polysaccharide solution (0.1 mL) of different concentrations (0.1% to 1.5%) were mixed with 1 mL of the reagent solution (0.6 mol.L-1 sulfuric acid, 28×10-3 mol.L-1 sodium phosphate, 179/186 179
Castro, J. P. L., Costa, L. E. C., Pinheiro, M. P., Francisco, T. S., Vasconcelos, P. H. M., Funari, L. M., Daudt, R. M., Santos, G. R. C., Cardozo, N. S. M., & Freitas, A. L. P. and 4×10-3 mol.L-1 ammonium molybdate). This step was followed by incubation at 95 °C for 90 min. Subsequently, absorbance was read at 695 nm. A standard curve under the same conditions was prepared with solutions of ascorbic acid. Thus, the results are presented as equivalence of ascorbic acid (mg/g EAAsc). 2.7.2 Iron (Fe 2+) chelating activity The iron ion is associated with lipid peroxidation due to the Fenton reaction. This process is related to a series of diseases[43-45]. The antioxidants form a complex with this ion, thus preventing cell damage[46]. The ferrous ion chelation activity of PLS was investigated according to the methodology used by Zhang et al.[18]. Aliquots of 1 mL of the polysaccharide solution at different concentrations (0.1% to 1.5%) were mixed with 0.05 mL of FeCl2 (2×10-3 mol.L-1), 0.2 mL of ferrozine (5×10-3 mol.L-1), and 2.75 mL of water. The solution was agitated and incubated at room temperature for 10 min. Absorbance of the solution at 562 nm was measured. The chelating activity was calculated using the following Equation 1: A1 − A2 Chelating activity ( % ) = 1 − ×100 (1) A0
where A0 is the absorbance of the blank, A1 is the absorbance of the test solution, and A2 is the absorbance of a solution identical to A1 with the substitution of FeCl2 for the same aliquot of water. Ascorbic acid was used as the positive control. 2.7.3 Scavenging of DPPH radicals The 1,1-diphenyl-2-picrylhydrazyl free radical (DPPH) scavenging activity of PLS was measured using the method used by Wu et al.[16]. Ascorbic acid was used as a positive control. The inhibition (%) was calculated using the following Equation 2: Scavenging of DPPH radicals = (%)
( A0 − A1 ) A0
×100 (2)
where A0 and A1 are the absorbance of the blank and sample, respectively.
2.8 Rheological behavior Rheological measurements were carried out in a rotational rheometer (Ares, TA Instruments, New Castle, DE, USA), using cone-plate geometry (50 mm diameter, cone angle of 0.0399 rad, gap of 0.0553 mm). All measurements were run at 25 ± 1 °C. Dynamic tests were performed to evaluate the behavior of storage (G’) and loss (G”) moduli as a function of frequency. The frequency sweep tests were performed in the linear viscoelastic region (LVR), determined through strain sweep tests. Three samples containing the extracted sulfated PLS (1.0%, 1.25%, and 1.5%) were analyzed. Flow curves were obtained by recording shear stress values when shearing the samples at an increasing shear rate from 0.1 to 100 s-1 with increment of 4 s-1 and then reducing it through the same path. This test was performed for samples containing PLS in different concentrations (1.0%, 1.25%, and 1.5%). Additionally, for testing the potential of PLS as thickening agent, solutions containing 180 180/186
carboxymethylcellulose (CMC), a well-known thickening agent, and blends of CMC and PLS were also analyzed. The concentrations of the referred CMC solutions were of 0.75% of CMC and 1.5%. In the case of the blends, two formulations were prepared, one containing 1.0% of PLS and 0.5% of CMC and another containing 0.75% of PLS and 0.75% of CMC. The data of the flow curves were fitted using the Ostwald‑de Waele model (power-law), Equation 3: η= K γ n −1 (3)
where: is the shear viscosity (Pa⋅s), K is the consistency index (Pa.s); is the shear rate (s-1), n is the power-law index (dimensionless) and K is the consistency index (Pa.s). The estimation of the parameters K and n was performed using the least-square method, in the software Microsoft Excel.
3. Results and Discussions 3.1 Yield and chemical analysis The yield of polysaccharides by mass of seaweed Gracilaria intermedia was 17.0 ± 1.18%. Studies which extraction took place by enzymatic digestion of seaweed Gracilaria cornea (Brazil) achieved yields of 18.0%[28] and 11.0 to 21.4%[47]. Souza et al.[12] in Brazil obtained a polysaccharide yield of 27.2% for the seaweed Gracilaria birdeae while Freile-Pelegrin and Robledo[48], on a seasonal study in Mexico, obtained yield from 25.0 to 39.3% for Gracilaria cervicornis. In extraction performed at room temperature, in Brazil, Maciel et al.[23] were able to yield 6.5% of polysaccharides from seaweed Gracilaria birdeae. Making the extraction of polysaccharides with autoclave in India, yields of 14.8% for Gracilaria debilis and 15.2% for Gracilaria Salicornia[49] were achieved. Polysaccharides from Gracilaria intermedia (PLS) exhibited 6.60 ± 0.13% of sulfate, which is within the range of sulfate polysaccharides content from the other Gracilaria species (2.30-8.90%)[23]. More recent studies indicate that this interval is actually wider, since percentages of sulphate of 0.76 ± 0.08% (Gracilaria debilis, India)[49], 1.00 ± 0.05% (Gracilaria caudata, Brazil)[26], and 15.66% (Gracilaria cornea, Brazil)[28] have already been reported. The carbohydrate content was 54.64 ± 1.19%, consistent with the percentage value of D - galactose found in other Gracilaria species. Amorin et al.[50] performed the test sulfuric phenol in different fractions of Gracilaria ornata, Brazil, and found levels of sugars ranging between 33.14 and 62.20%. In polysaccharide fractions of seaweed Gracilaria birdeae, Brazil, sugar content between 30.8 and 68.2%[51] were found. Proteins were not detected in the polysaccharide fraction obtained from seaweed Gracilaria intermedia.
3.2 Monosaccharide composition The monosaccharide composition of the polysaccharide extracted from G.intermedia was determined based on gas chromatography/mass spectrometry analysis of the alditol acetates formed after acid hydrolysis. The major monosaccharide detected was galactose. This finding is in agreement with results presented by Pomin & Mourão[52] Polímeros, 28(2), 178-186, 2018
Polysaccharides of red alga Gracilaria intermedia: structure, antioxidant activity and rheological behavior who describe the presence of sulfated galactans in red algae. No other sugar was detected up to a limit of < 2% as % of dry weight, ensuring the purity of the material.
3.3 Infrared spectroscopy Figure 2 shows the FTIR spectrum of the extracted polysaccharide fraction, expanded in the region between 1400 and 700 cm-1 to better identify the sulfate groups bands. The bands found at 1375 and 1258 cm-1 may be attributed to sulfate ester groups[51,52], the band at 1075 cm-1 corresponds to galactan[53,54], and the band at 892 cm-1 corresponds to the agar-specific band[55]. The band at 931 cm-1 corresponds to the C-O-O group present in the 3,6 -anhydrogalactose, and the bands between 820 and 860 cm-1 indicate the presence of sulfate groups[53-57].
3.4 NMR spectroscopy Figures 3 and 4 show, respectively, the 1D (1H and 13C) and 2D NMR spectra of the polysaccharide fraction extracted from G.intermedia. The 1H NMR spectrum is somewhat complex due to overlap and enlargement of the signs (Figure 3a). It shows the signals from the α-anomeric proton at δ 5.62, 5.14 and 4.69 (assigned, respectively, to 3,6-α-Lanhydrogalactose linked to β-D-galactose, α-L‑galactose6-sulfate linked to β-D-galactose and β-D‑galactose linked to 3,6-α-L-anhydrogalactose), and from the H-1 of β-D-galactose linked to 3,6-α-L-anhydrogalactose with a signal at δ 4.69[23,26,38-41,58]. However, the H-1 β-D-galactose linked to α-L-galactose-6-sulfate was not detected, probably due to the signal overlap. Others studies state that when this unit is linked with α-3,6-anhydrogalactose and L-α-Lgalactose-6-sulfate, there are minor chemical variations in the region of 4.65-3.90 ppm. Therefore, its identification is difficult[23,26,58,59]. The anomeric region of 13C NMR (90-110 ppm) shows four main signals (Figure 3b), which were assigned based on the literature data[23,26,38-41,58,60,61], just like the other carbons observed in the region of 59 – 85 ppm (Table 1). The C-1 of β-D-galactose linked to α-L-galactose-6-sulfato at δ 103.7; C-1 of β-D-galactose linked to 3,6-α-L-anhydrogalactose at δ 102.5; C-1 of α-L-galactose-6-sulfato linked to β-D‑galactose at δ 101.25 and C-1 of 3,6-α-L-anhydrogalactose linked to β-D-galactose at δ 98.4. 2D COSY was used to determine the proton resonance sequence (Figure 4). Regarding the dimer formed by 3,6-α-L-anhydrogalactose linked to β-D-galactose five couplings relating the 3,6-α-L-anhydrogalactose u and two couplings assigned β-D-galactose are observed. The protons that are coupled are H-1 (3,6-α-L-anhydrogalactose)/H-2 (3,6-α-L-anhydrogalactose) at δ 5.62/4.60; H-2 (3,6-α-L‑anhydrogalactose)/H-3 (3,6-α-L-anhydrogalactose) at δ 4.60/5.04; H-3 (3,6-α-L‑anhydrogalactose)/H-4 (3,6-α-L-anhydrogalactose) at δ 5.04/4.11; H-4 (3,6-α-L‑anhydrogalactose)/H-5 (3,6-α-L‑anhydrogalactose) at δ 4.11/4.26; H-1 (β-D-galactose)/H-2 (β-D-galactose) at δ 4.69/4.44 and H-2 (β-D-galactose)/H-3 (β-D-galactose) at δ 4.44/3.91. For the α-L-galactose-6-sulfate dimer linked to β-D-galactose a coupling is observed for the H-1 (α-L‑galactose-6-sulfato)/H-2 (α-L-galactose-6-sulfato) at δ 5.14/5.05. Although other engagements can be observed, Polímeros, 28(2), 178-186, 2018
Figure 2. The IR spectra of the PLSs were determined using a Fourier transform infrared spectrometer (FTIR).
Figure 3. 1H NMR (a) and 13C NMR (b) spectra of the PLSs extracted from Gracilaria intermedia in D2O solution.
it is not possible to identify the sequence of hydrogens because the one-dimensional 1H spectrum overlap them.
3.5 Determination of antioxidant activity The total antioxidant capacity was determined by forming a phosphomolybdenum complex when Mo6+ is reduced to Mo5+; PLS showed activity with 28.98 ± 1.86 mg/ g EAAsc. In studies performed by Costa et al.[62] the amount of ascorbic acid in the seaweeds Codium isthmocladum and Spatoglossum schroederi were (9.2 mg/g EAAsc) and (14.4 mg/g EAAsc) respectively. These values are inferior amount to the ones exhibited by the polysaccharides extracted from G. intermedia. The ability of the PLS to chelate iron (II) ions was dose‑dependent. Even though PLS led to higher chelation percentage than the ones obtained from ascorbic acid 181/186 181
Castro, J. P. L., Costa, L. E. C., Pinheiro, M. P., Francisco, T. S., Vasconcelos, P. H. M., Funari, L. M., Daudt, R. M., Santos, G. R. C., Cardozo, N. S. M., & Freitas, A. L. P. Table 1. 1H and 13C NMR chemical shifts for residues of G. intermedia polysaccharide. Residue
H chemical shift (ppm) H-3 H-4 5.04 4.11 3.91 nd* nd* nd* 13 C chemical shift (ppm) C-3 C-4 81.2 77.5 82.1 68.7 nd* 78.75 80.2 69.3 1
3,6-α-L-anhydrogalactose linked to β-D-galactose β-D-galactose linked to 3,6-α-L-anydrogalactose α-L-galactose-6-sulfato linked to β-D-galactose
H-1 5.62 4.69 5.14
H-2 4.60 4.44 5.05
3,6-α-L-anhydrogalactose linked to β-D-galactose β-D-galactose linked to 3,6-α-L-anhydrogalactose α-L-galactose linked to β-D-galactose β-D-galactose linked to α-L-galactose-6-sulfato
C-1 98.4 102.5 101.25 103.7
C-2 nd* nd* nd* 70.5
H-5 4.26 nd* nd*
H-6 nd* nd* nd*
C-5 nd* 73.8 71.9 75.6
C-6 69.5 61.5 67.5 61.8
nd, not detected; *Signal which can be overlapped with the signal of similar units.
Figure 4. 2D COSY spectrum of PLSs from Gracilaria intermedia in D2O.
(Figure 5a), its activity is low if compared to other polysaccharides extracted from seaweeds. The values for PLS were 11.64 ± 0.83% (PLS 0.1%) and 2819 ± 0.97% (PLS 1.5%). Costa et al.[62] found a value of 40.2% of iron (III) chelation in solutions of 1.5 mg.mL-1 of polysaccharides from Gracilaria caudata. Wang et al.[63] studied different polysaccharide fractions from Laminaria japonica and found a value of 29.48% for quelation in a 1.17 mg.mL-1 solution. The DPPH radical scavenging capacity of PLS was also dose-dependent, with a blockage varying from 9.89 ± 1.32% (at a concentration of 0.1%) to 41.83 ± 0.97% (at a concentration of 1.5%). Although significant, the DPPH radical scavenging capacity of PLS was lower than that ascorbic acid over the concentration range tested (Figure 5b) and also lower than the activity of other polysaccharides from seaweeds. Dore et al.[14], with polysaccharides from Sargassum vulgare, found values of 10.0 ± 0.7% a 22.0 ± 0.6% in solutions that varied from 0.15 to 3.0 mg.mL-1 182 182/186
3.6 Rheological characterization Based on the strain sweep tests performed to determine the region linear viscoelastic behavior of the samples in the dynamic oscillatory tests, the frequency sweep (FS) tests were carried out at 8% of a deformation. The results of the FS tests are presented in Figure 6, where it can be seen that the values of loss modulus (G”) were higher than those of the storage modulus (G’) at all concentrations and frequencies tested. These results indicate that the extracted polysaccharide fraction does not lead to gel-like solutions. This behavior can be attributed to the elevated amount of sulfate that was observed in the polysaccharide obtained by papain digestion. The structure of agarans is strongly influenced by the presence of charged groups that participate in intermolecular hydrogen bonds[48]; the strength of this polymer gel is inversely proportional to the amount of sulfate present in its structure[64]. The predominantly viscous response reflected in the dynamic and steady state tests suggest that a potential use Polímeros, 28(2), 178-186, 2018
Polysaccharides of red alga Gracilaria intermedia: structure, antioxidant activity and rheological behavior
Figure 5. Antioxidant activity of the PLSs extracted from Gracilaria intermedia. (a) Chelating activity of Fe2+ with ascorbic acid used as a positive control. (b) DPPH radical scavenging activity. Values are means ± SD (n = 3).
for the polysaccharides extracted from the red seaweed Gracilaria intermedia is as a thickening agent, especially for applications where antioxidant activity and absence of color and odor are mandatory requirements. To check this hypothesis, the flow curves at 25 ± 1 °C of solutions of the extracted polysaccharides, either pure or in mixture with carboxymethylcellulose (CMC), were compared to those of two solutions containing pure CMC. The concentrations of the pure CMC solutions were of 0.5% and 1.5%, corresponding, respectively, to the lower and upper limit of the range of concentrations typically used in commercial applications of CMC as thickening agent. In the mixtures of CMC with the extracted polysaccharides, the total concentration of thickening agents was 1.5%. The obtained flow curves are presented in Figure 7. Figure 7 shows that all the tested solutions presented pseudoplastic behavior, since an increase in shear rate led to a decrease in viscosity. Besides, for all samples, the upward and downward (not shown) curves of shear stress vs. shear rate were identical, with no observed hysteresis behavior. Therefore, the extracted polysaccahride fraction solutions, as well as the two solutions with pure CMC, showed no thixotropy in the range of concentrations tested. In the case of the extracted polysaccharide fraction solution, the pseudoplastic behavior is likely due to the rupture of the double helix structure present in agarans[65]. Additionally, the fact that this rupture only occurs above a certain value of tension (tcrit)[65] is in agreement with the two-region pattern observed in the flow curves, with constant viscosity at low shear rates (Newtonian plateau) and shear-thinning behavior at higher ones. The higher the polysaccharide concentration, the lower the upper limit of the Newtonian plateau because an increase in viscosity causes tcrit be reached at lower values of shear rate. Additionally, the data corresponding to the shear thinning region of all flow curves presented in Figure 7 were fitted to power-law model. The obtained values of consistency index (K) and power-law index (n) are presented in Table 2. It is observed that increase of the polysaccharide concentration in the PLS solutions led to increase of the consistency index and reduction of the power-law index. Taking into account Polímeros, 28(2), 178-186, 2018
Figure 6. Effect of the concentration of PLS extracted from Gracilaria intermedia on the storage and loss moduli measured during frequency sweeps.
F i g u re 7 . F l o w c u r v e s o f t h e P L S s e x t r a c t e d f r o m Gracilaria intermedia and of CMC at 25 ± 1 °C.
that the consistency index is related to the resistance of a fluid to the flow and power-law index indicates how quickly the viscosity is reduced with an increase in the shear rate, the increase of K and reduction of n with the increase of the concentration of PLS in the solution can be explained 183/186 183
Castro, J. P. L., Costa, L. E. C., Pinheiro, M. P., Francisco, T. S., Vasconcelos, P. H. M., Funari, L. M., Daudt, R. M., Santos, G. R. C., Cardozo, N. S. M., & Freitas, A. L. P. Table 2. Power index (n) and consistency index (K) values for PLS, CMC solutions, and blends of both samples. Sample PLS G. intermedia
CMC PLS + CMC
Concentration (%) 1.0 1.25 1.5 0.75 1.5 1.0 PLS + 0.5 CMC 0.75 PLS + 0.75 CMC
n 0.67 0.62 0.47 0.60 0.29 0.43 0.39
K (Pa.s) 2.31 5.76 11.64 1.44 47.23 15.83 10.11
in terms of the increase in the level of interaction among polysaccharide molecules resulting from the increase in the concentration. As observed in Figure 7, although pure CMC led to the highest viscosity, both pure PLS and the PLS/CMC blends were able to produce a significant increase of viscosity. Actually, all over the range of strain rates studied, pure PLS and the PLS/CMC blends provided viscosity values closer to the that obtained with the CMC solution of 1.5% than to those obtained with 0.5%. These results confirm the potential of the extracted sulfated PLS to be used as thickening agent.
4. Conclusions In this study, a water-soluble polysaccharide was successfully extracted from the red seaweed Gracilaria intermedia. The data obtained by spectroscopic methods suggest that the extracted polymeric material is rich in agaran. In vitro studies showed that PLS had antioxidant activity, which may be used for cellular protection against free radicals and to increase the viability of products. The rheological characteristics of PLS indicate that it can be used as a viscosity modifier in industrial processes. The antioxidant characteristics and rheological behavior combined with the organoleptic properties observed (absence of color and odor) make polysaccharides from the seaweed Gracilaria intermedia a promising agent to be used in manufacturing processes of food and pharmaceuticals.
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