Polímeros: Ciência e Tecnologia (Polimeros) 3rd. issue, vol. 30, 2020 by Polímeros: Ciência e Tecnologia (Polimeros)

Polímeros VOLUME XXX - Issue III - July/Sep., 2020

São Paulo 994 St. São Carlos, SP, Brazil, 13560-340 Phone: +55 16 3374-3949 Email: abpol@abpol.org.br 2020

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

P o l í m e r o s - I ss u e I I I - V o l u m e X X X - 2 0 2 0 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 a n d 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 ”

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

Editorial Committee

Antonio Aprigio S. Curvelo (USP/IQSC) - President

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

Members Adhemar C. Ruvolo Filho (UFSCar/DQ) Ailton S. Gomes (UFRJ/IMA) Alain Dufresne (Grenoble INP/Pagora) 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) 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) Luiz Antonio Pessan (UFSCar/DEMa) Luiz Henrique C. Mattoso (EMBRAPA) Marcelo Silveira Rabello (UFCG/UAEMa) 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) 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

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

www.editoracubo.com.br

“Polímeros” is a publication of the Associação Brasileira de Polímeros São Paulo 994 St. São Carlos, SP, Brazil, 13560-340 Phone: +55 16 3374-3949 emails: abpol@abpol.org.br / revista@abpol.org.br http://www.abpol.org.br Date of publication: September 2020

Financial support:

Available online at: www.scielo.br

Polímeros / Associação Brasileira de Polímeros. vol. 1, nº 1 (1991) -.- São Carlos: ABPol, 1991Quarterly v. 30, nº 3 (July/Sept. 2020) ISSN 0104-1428 ISSN 1678-5169 (electronic version)

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

1. Polímeros. l. Associação Brasileira de Polímeros. Polímeros, 30(3), 2020

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

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

Editorial Section News....................................................................................................................................................................................................E3 Agenda.................................................................................................................................................................................................E4 Funding Institutions.............................................................................................................................................................................E5

O r i g in a l A r t ic l e Investigation of Lactobacillus paracasei encapsulation in electrospun fibers of Eudragit® L100 Juliana Mikaelly Dias Soares, Ruan Emmanuell Franco Abreu, Mateus Matiuzzi da Costa, Natoniel Franklin de Melo and Helinando Pequeno de Oliveira........................................................................................................................................................................ 1-8

Separation of PET from other plastics by flotation combined with alkaline pretreatment Fernando Pita and Ana Castilho...................................................................................................................................................................... 1-9

Water vapor permeation and morphology of polysulfone membranes prepared by phase inversion Luis Guilherme Macedo Baldo, Marcelo Kaminski Lenzi and Daniel Eiras................................................................................................... 1-8

Rheological and thermal properties of EVA-organoclay systems using an environmentally friendly clay modifier Reinaldo Yoshio Morita, Juliana Regina Kloss, Ronilson Vasconcelos Barbosa, Bluma Guenther Soares, Luis Carlos Oliveira da Silva and Ana Lúcia Nazareth da Silva............................................................................................................................................................................ 1-9

Improved durability of Bisphenol A polycarbonate by bilayer ceramic nano-coatings alumina-zinc oxide Abdellah Moustaghfir, Agnes Rivaton, Bénédicte Mailhot and Michel Jacquet............................................................................................... 1-7

Layered cryogels laden with Brazilian honey intended for wound care Gabriela de Souza dos Santos, Natália Rodrigues Rojas dos Santos, Ingrid Cristina Soares Pereira, Antonio José de Andrade Júnior, Edla Maria Bezerra Lima, Adriana Paula Minguita, Luiz Henrique Guerreiro Rosado, Ana Paula Duarte Moreira, Antonieta Middea, Edlene Ribeiro Prudencio, Rosa Helena Luchese and Renata Nunes Oliveira.............................................................................................. 1-10

Reactive compatibilization effect of graphene oxide reinforced butyl rubber nanocomposites Sathishranganathan Chinnasamy, Rajasekar Rathanasamy, Harikrishna Kumar Mohan Kumar, Prakash Maran Jeganathan, Sathish Kumar Palaniappan and Samir Kumar Pal.......................................................................................................................................... 1-8

Nanoscale morphology, structure and fractal study of kefir microbial films grown in natura Robert S. Matos, Ellen C. M. Gonçalves, Erveton P. Pinto, Gerson A. C. Lopes,Nilson S. Ferreira and Cristiane X. Resende .................... 1-7

Physicochemical characterization, drug release and mechanical analysis of ibuprofen-loaded uhmwpe for orthopedic applications Loise Silveira da Silva, Izabelle de Mello Gindri, Gean Vitor Salmoria andCarlos Rodrigo de Mello Roesler.............................................. 1-8

Comparison of MA-g-PP effectiveness through mechanical performance of functionalised graphene reinforced polypropylene Saravanan Natarajan, Rajasekar Rathanasamy, Sathish Kumar Palaniappan, Suresh Velayudham, Hari Bodipatti Subburamamurthy and Kaushik Pal....................................................................................................................................................................................................... 1-7

Isolation and characterization of micro cellulose obtained from waste mango Miguel Angel Lorenzo-Santiago and Rodolfo Rendón-Villalobos.................................................................................................................... 1-8

Evaluation of fracture toughness of epoxy polymer composite incorporating micro/nano silica, rubber and CNTs Ronaldo Câmara Cozza and Vikas Verma....................................................................................................................................................... 1-14 Cover: Scanning electron microscopy images are shown for (1a), L. paracasei (L) (1b), EDGT (1c) EDGT-L.paracasei and (1d) average diameter EDGT, EDGT-L.paracasei and L. paracasei – calculated from sets of 25 different fibers per SEM image. Fluorescence images for samples: (a) L. paracasei treated with AO, (b, c and d) EDGT-L.paracasei treated with AO. Arts by Editora Cubo.

Polímeros, 30(3), 2020

The automotive industry is changing gears when it comes to thermoplastics As the engines of cars become more and more fuel efficient, ever more compact and increasingly more sophisticated in order to satisfy lower and lower emissions requirements, manufacturers face many industrial challenges. They have to explore how to build engines in ways that combine lightweighting, downsizing and the capacity to function in more aggressive environments (specifically with higher temperatures). On top of that, the rapid commercial take-off of hybrid and electric vehicles comes with its own challenges, as their engines are more complex. One result of all this is that cooling systems in cars are challenged too, as there are more elements to cool down, higher temperatures to deal with and less space to fit it all. “Thermal management systems have become more and more complex,” confirms Brian Baleno, Head of Marketing for Transportation, at Solvay’s Specialty Polymers business unit. “Yet when manufacturers design an engine, cooling systems are one of the last things they look at. So they need a solution enabling them to design with very complex geometries, simply because there isn’t much space left.” All this means the automotive industry is faced with the following conundrum: they can’t use metal because of its weight, they can’t use standard plastics because they couldn’t withstand the increasingly harsh conditions, yet they need a material that can be designed into intricate shapes. At least one thing is certain: traditional thermal management systems made of aluminum tubes bent into shape need to evolve - and that creates new opportunities in the high temperature thermoplastics market. Among Solvay’s wide portfolio of solutions that can help the transportation industry meet its lightweighting challenges, there is at least one thermoplastic that answers all of the abovementioned requirements. Ryton® polyphenylene sulfide (PPS) is a high performance thermoplastic that combines mechanical, chemical and thermal resistance with design flexibility. It’s also flame retardant, a desirable feature in high-temperature, highvoltage environments. With features such as these, Ryton® PPS has unsurprisingly been used as a replacement for metal in cars and in the electronics industry for several decades now, primarily for the 30% to 50% weight reduction it enables. What’s new is the increasing extent of its applications, driven by manufacturers’ need for innovative material solutions. “Our R&I teams adapted Ryton® PPS to make it compatible with extrusion manufacturing instead of traditional injection molding,” explains Brian. “Thanks to this, more and more OEMs are adopting it to make highly specific cooling line parts, for which injection molding isn’t an option, which created a new application for an existing product.” Moreover, the combination of properties offered by Ryton® PPS means it cannot only advantageously replace metal, but also ‘lower-performance’ plastics. On top of its resistance to heat, chemicals and mechanical stress, it also provides dimensional stability, meaning that parts maintain their original dimensions even when exposed to moisture or heat, which eliminates the risk of leakage, as gaps are less likely to appear between parts. With the complex thermal management systems of hybrid and electric vehicles, that’s a precious safety feature to have. New opportunities through electrification Beyond cooling systems, Ryton® PPS is also increasingly adopted for powertrain components as well as for parts in electric mobility systems, including belt starter generators, traction motors, power electronics, and fuel cells. Within these, applications such as sensors, connectors, bus bars, and rectifiers use Ryton® PPS. Here too, it replaces other thermoplastics that just don’t offer the same characteristics. “The electrical properties of Ryton® PPS are yet another argument for its adoption,” continues Brian. “E-mobility is definitely a new area of growth, as it creates more opportunities to replace metal, for example in electronically-driven pumps. There are more water, oil and vacuum pumps in a hybrid car because there are more elements to cool, and you need to make them with weight optimization in mind on top of everything else, so you need a material that will enable you to do that.” Source: Solvay - www.solvay.com

Polímeros, 30(3), 2020

Porsche and Circularise collaborate with Borealis, Covestro and Domo Chemicals to enable the traceability of plastics in the automotive sector Circularise, the blockchain supply chain transparency provider, as part of the Startup Autobahn innovation program, recently launched a project with Porsche and its pioneering material suppliers – Borealis, Covestro and Domo Chemicals – to enable the traceability of plastics on blockchain and to ensure that the use of sustainable materials in Porsche cars can be proven. By digitizing materials Circularise was able to create a digital thread through the whole supply chain, enabling material traceability, tracking the CO2 footprint and other sustainability metrics like water savings. Getting information from supply chains has always been a challenge. Not only because of the inherent complexity of the supply chains and the multitude of suppliers, but also due to concerns around trust, privacy and confidentiality. That is why blockchain is offering such a fitting solution to transparency challenges in supply chains. “We believe transparency should not come at the cost of reduced privacy and confidentiality. That is why we developed our patent pending technology for creating verified statements on public blockchains without revealing any underlying sensitive data. While this raw data is very valuable in a B2B setting, consumers demand a more distilled and interactive version. We are proud to present exactly that in collaboration with Porsche and some of their pioneering suppliers,” says Mesbah Sabur. Porsche has a large number of suppliers providing parts to its cars but it doesn’t stop the company looking for more information about the materials that go into its cars. According to Antoon Versteeg, Project Lead Innovation Research at Porsche, “We need to know more details on the parts and materials being used in our products, that means information on production processes deep down the supply chain, statements of recycled content and more. With the help of Circularise, as well as with the help of their partners we were able to trace for a number of specific cases plastics from raw material production to the final car.” A number of suppliers who can deliver sustainably produced materials for the automotive industry were involved in this project to realise the final outcome. Each batch of material was digitized on the blockchain receiving a digital copy called digital twin. The digital twin carries all relevant information regarding the batch, such as its environmental footprint and origin. This digital thread created transparency between project partners leading to an improved supply chain collaboration. However, the companies cannot simply create a digital twin. First, the batch of materials needs to be audited by an independent third party to verify that the material and related claims are true. “Verification is essential. Even with a supply chain involving blockchains we want independent auditors for our system. And this is how we gain the trust and confidence of all our value chain members.” says Christopher McArdle, Borealis Vice President Polyolefin Strategy and New Business Development. Once the materials are digitized, the parties along the supply chain can now update the digital twin mimicking the physical supply chain and reflecting the manufacturing processes along the lifecycle of the product. Burkhard Zimmermann, Head of Resin, Digital Transformation & Sustainability at Covestro’s Polycarbonates segment: “For us, it is really important to share information and be more transparent while maintaining confidentiality. For instance, the material composition is of competitive advantage so we would never share that openly. Here, Circularise helps us to maintain this confidentiality and only disclose the information needed from raw material producer to recycler. And with that, we can close the loop.” Not only this approach helps car manufacturers to make better decisions for the next generations of vehicles and support end-of-life recycling approaches, it also helps final consumers to learn more about their vehicle and its origins, enabling them to make more sustainable choices. Ultimately reducing the environmental impact across the whole value chain. Source: Borealis - www.borealisgroup.com

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May World Polymer Congress (IUPAC-MACRO2020+) Date: May 16-20, 2021 Location: Jeju Island, South Korea Website: www.macro2020.org Polymers 2021: New Trends in Polymer Science: Health of the Planet, Health of the People Date: May 17-19, 2021 Location: Turin, Italy Website: polymers2021.sciforum.net

June Gordon Research Conference — Polymers Date: June 5-6, 2021 Location: South Hadley, United States Website: www.grc.org/polymers-conference/2021 Biopolymer – Processing & Moulding Date: June 15-16, 2021 Location: Halle (Saale), Germany Website: https://polykum.de/en/biopolymer-mkt-2021 RosUpack - 25th International Exhibition for the Packaging Industry Date: June 15-18, 2021 Location: Moscow, Russia Website: www.rosupack.com Gordon Research Conference — Polyamines Date: June 27 - July 2, 2021 Location: Waterville Valley, United States Website: www.grc.org/polyamines-conference/2021

July 25th IUPAC International Conference on Physical Organic Chemistry Date: July 10–15, 2021 Location: Hiroshima, Japan Website: icpoc25.jp 84th Prague Meeting on Macromolecules - Frontiers of Polymer Colloids Date: July 18–22, 2021 Location: Prague, Czech Republic Website: mns-20.com 2nd International Conference on Materials and Nanomaterials (MNs-20) Date: July 26–28, 2021 Location: Rome, Italy Website: www.imc.cas.cz/sympo/84pmm

August International Conference on Electronic Materials (2021 IUMRS-ICEM) XIX Brazilian Materials Research Society Meeting (XIX B-MRS) Date: August 29 – September 2, 2021 Location: Foz do Iguaçu, Brazil Website: www.sbpmat.org.br/19encontro

7th Edition of International Conference on Polymer Science and Technology Date: September 13-14, 2021 Location: Berlin, Germany Website: https://polymerscience.euroscicon.com/ 9th International Conference on Fracture of Polymers, Composites and Adhesives Date: September 26-30, 2021 Location: Les Diablerets, Switzerland Website: www.elsevier.com/events/conferences/esistc4conference 36th International Conference of the Polymer Processing Society Date: September 26-30, 2021 Location: Montreal, Canada Website: www.polymtl.ca/pps-36/en 13th PVC Formulation Date: September 27-29, 2021 Location: Cologne, Germany Website: www.ami.international/events/event?Code=C1104

October International Conference on Materials Science and Engineering Date: October 11-14, 2021 Location: Brisbane, Australia Website: www.materialsconferenceaustralia.com Sustainable Polymers Date: October 17-20, 2021 Location: Safety Harbor, United States Website: www.polyacs.net/21sustainablepolymers 16th Brazilian Polymer Conference – (16thCBPol) Date: October 24-28, 2021 Location: Ouro Preto, Brazil Website: www.cbpol.com.br

November Performance Polyamides USA Date: November 2, 2021 Location: Cleveland, United States Website: www.ami.international/events/event?Code=C1149 Plástico Brasil Date: November 8-12, 2021 Location: São Paulo, Brazil Website: www.plasticobrasil.com.br Multilayer Flexible Packaging Date: November 23-25, 2021 Location: Barcelona, Spain Website: www.ami.international/events/event?Code=C1147

December Polymers in Flooring Date: December 9-10, 2021 Location: Berlin, Germany Website: https://www.ami.international/events/ event?Code=C1150

September Global Summit and Expo on Materials Science and Nanoscience Date: September 6-8, 2021 Location: Lisbon, Portugal Website: www.thescientistt.com/materials-science-nanoscience CIRM – Workshop — Directed Polymers and Folding Date: September 6-10, 2021 Location: Marseille, France Website: https://conferences.cirm-math.fr/2021-calendar.html 100 Years Macromolecular Chemistry Date: September 12-14, 2021 Location: Freiburg, Germany Website: veranstaltungen.gdch.de/tms/frontend/index. cfm?l=9162&sp_id=2 E4

Polímeros, 30(3), 2020

ABPol Associates Sponsoring Partners

Polímeros, 30(3), 2020

ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.03020

Investigation of Lactobacillus paracasei encapsulation in electrospun fibers of Eudragit® L100 Juliana Mikaelly Dias Soares1,2 , Ruan Emmanuell Franco Abreu2 , Mateus Matiuzzi da Costa1,2 , Natoniel Franklin de Melo3  and Helinando Pequeno de Oliveira1,2*  Rede Nordeste de Biotecnologia – RENORBIO, Universidade Federal Rural de Pernambuco – UFRPE, Recife, PE, Brasil 2 Universidade Federal do Vale do São Francisco – UNIVASF, Petrolina, PE, Brasil 3 Embrapa Semiárido, Petrolina, PE, Brasil

*helinando.oliveira@univasf.edu.br

Abstract Some species of Lactobacillus have demonstrated beneficial health effects being applied in the production of food supplements. Thus, the incorporation of viable cells as encapsulated probiotics represents an essential condition to be considered in new strategies for the controlled release of microorganisms. Herein, the massive encapsulation of Lactobacillus paracasei is provided by the use of alternative electrospinning technique. Is spite of the high voltage required for the production of fibers, a high density of viable cells is observed into the polymeric electrospun web, allowing the controlled release at targeted pH (characteristic of Eudragit® L100 polymer support). The reported procedure circumvents typical drawbacks of degradation of microorganisms under adverse conditions (storage, package and low pH) and preserves its biologic action after complete release from polymer fibers. Keywords: probiotics, encapsulation, electrospinning, food, Lactobacillus. How to cite: Soares, J. M. D., Abreu, R. E. F., Costa, M. M., Melo, N. F., & Oliveira, H. P. (2020). Investigation of Lactobacillus paracasei encapsulation in electrospun fibers of Eudragit® L100. Polímeros: Ciência e Tecnologia, 30(3), e2020025. https://doi.org/10.1590/0104-1428.03020

1. Introduction Some probiotic bacteria have positive physiological effects and have been considered as important components for the production of foods supplements[1,2]. In particular, Lactobacillus spp. has been considered as a promising probiotic that confers health benefits to the host. Recent studies considered the use of Lactobacillus spp. in the prevention and treatment of inflammatory bowel disease[3,4], food hypersensitivity[5], cardiometabolic disorders[6] and anti-tumor activity[7]. However, it has also been reported that bioactive living cells, such as the Lactobacillus spp. present low bioavailability/ biofunctionality as a consequence of transport through the gastrointestinal tract, in the processing and/or prolonged storage[8,9]. The acidic medium can induce changes on bacterial membrane components, modifying and disturbing the peptidoglycan components, lipids, proteins and DNA in Gram-positive bacteria[10]. As a consequence, the survivability and colonization in the digestive tract are considered critical to ensure optimal functionality of Lactobacillus species[11]. The poor survival rate of bioactive cells can be attributed to environmental conditions such as acidic medium, the toxicity of oxygen and UV light[12-16]. Healthy-promoting effects of probiotics are extremely dependent on cell viability degree and the concentration of living cells as high as 109 CFU/ day for

Polímeros, 30(3), e2020025, 2020

administration[17]. To circumvent the drawbacks related to low shelf-life of food products and adverse conditions at the acidic environment (stomach/ bile salts) the creation of an anaerobic environment for probiotics growth received increased attention in the literature with promising strategies to maintain the viability of cells until to reach the colon lumen. These encapsulation strategies are based on the production of fruit bubles[18], nanoencapsulation by electrospinning[17,19] and by the production of microcapsules[20,21]. The electrospinning technique has been drawing attention in the encapsulation of Lactobacillus spp.[22-24]. Despite the adverse conditions from experimental setup (high voltage and the nature of organic solvents) – that could be harmful to remaining viable cells in culture, studies are reporting that Lactobacillus-loaded electrospun fibers may preserve metabolic activity, with increased stability and protection[25-28]. The basic experimental setup for electrospinning production requires the dispersion of additives (molecules of interest) in a polymeric solution to be incorporated in a compartment (a syringe) kept at a fixed pressure. The needle in the syringe is connected to a high voltage source and depending on a series of factors (such as the distance of dip of the syringe and the grounded target, density of the solution, the intensity of the electric field, local humidity and infusion rate) the production of the fibers takes place[29,30]. Under an adequate

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Soares, J. M. D., Abreu, R. E. F., Costa, M. M., Melo, N. F., & Oliveira, H. P. combination of parameters, the atomization of a polymeric solution results in fibers with the diameter ranging from micrometer scale to nanometer scale[31,32]. The nature of the polymeric support is a critical parameter applied in the determination of the density and physical-chemical properties of the resulting fibers. The family Eudragit® is a commercial group of pHdependent block polymer materials that presents important properties for encapsulation and release of active molecules at targeted pH, avoiding side effects for adsorption of drugs at low pH conditions in the organism. In particular, the Eudragit® L100 is an anionic methacrylic acid and methyl methacrylate copolymer which presents dependent solubility, with rapid dissolution in the upper intestine (pH ≥ 6)[33], that protects the encapsulated species from degradation in the stomach. The fast dissolution of the polymeric matrix at high pH can be explored as an alternative strategy to protect probiotics against adverse conditions in the stomach. Despite the conventional use of enteric polymers of this family in the form of microparticles for encapsulation, it is observed an important possibility of use of the matrix as support for electrospinning – as a consequence, the massive production of matrix encapsulating species can be reached. Thus, the aim of this work was to developing electrospun fibers of enteric polymers for the encapsulation of L. paracasei. Several studies have already demonstrated efficient encapsulation of Lactobacillus-loaded electrospun fibers[22,26,28,34,35]. However, we have explored a simple procedure for encapsulation of probiotics in electrospun fibers from Eudragit® L100 polymer solution in alcohol. The successful encapsulation of microorganisms into this matrix represents a step forward in the direction of the massive production of encapsulated probiotics with welldefined targeted pH for release.

2. Materials and Methods 2.1 Materials The Lactobacillus paracasei probiotic strain was isolated from silage composed of elephant grass (Pennisetum purpureum cv. Cameroon) plus grape residue and had its identification confirmed by previous 16S rDNA sequencing. The microorganisms were kept in culture medium plates containing Rogosa and Sharpe Agar in anaerobic conditions for 48 hours at 37 °C[36]. The strain was phenotypically and genotypically characterized[37] by sequencing the 16S rRNA gene[38]. Eudragit® L100 was donated from Evonik, alcohol and sodium alginate were purchased from Sigma-Aldrich and MRS broth and agar from Neogen. Potassium monobasic phosphate and sodium hydroxide were purchased from Vetec Quimica Fina Ltda. Ultrapure water was obtained by the Milli-Q® equipment.

2.2 Microorganisms growth conditions The cultures of L. paracasei were prepared by transfer into MRS broth cultures and then incubated anaerobically at 37 °C for 48 h, as reported in the literature with some modifications[39]. Following the incubation step, the mediacontaining cells were centrifuged at 5,000 rpm for 10 min 2/8

at 10 °C, after that the supernatant was removed and the cells were further washed twice in sterile simulated intestinal fluid solution, with centrifugation after each step. The washed cells were suspended in sterile simulated intestinal fluid solution and stored for later use.

2.3 Preparation of electrospun fibers Polymeric solutions were prepared from dispersion of 0.4 g of Eudragit® L100 and 2% of sodium alginate (w/v) in 2 mL of alcohol. After that, 500 µL of suspended L. paracasei in sterile simulated intestinal solution was added in the previous solution. Solutions without and with L. paracasei were loaded into 5 mL syringes fitted with a capillary (metal needle), which was mounted horizontally on a syringe pump. The electrode at high-voltage power supply was clamped to the capillary and an aluminum plate was used as a collector was grounded. The voltage of 15 kV was established with the nozzle-to-collector distance of 20 cm and the flow rate of 1.0 mL/h. The resulting samples were: electrospun fibers of Eudragit® L100 (EDGT) and electrospun fibers with L. paracasei (EDGT-L.paracasei).

2.4 Characterization of electrospun fibers morphology Scanning Electron Microscopy (SEM) was performed in an SEM Vega 3XM Tescan. The electrospun fibers of EDGT and EDGT-L.paracasei were examined from SEM and the mean diameter was measured using the ImageJ from 25 electrospun fibers randomly selected.

2.5 Fourier transform infrared spectroscopy Fourier Transform Infra-Red (FTIR) analysis was performed using an IR Prestige-21 FTIR Shimadzu by KBr method in the range of 4000 cm-1 - 500 cm-1. The FTIR spectra were used to identify the influence of L. paracasei in the overall structure of electrospun fibers.

2.6 Viability of L. paracasei in Eudragit® L100 electrospun fibers The viabilities of the L. paracasei cells in electrospun fibers were determined from the dissolution of the fibers into sterile simulated intestinal fluid and then plating in MRS broth and agar. The sterile simulated intestinal fluid solution was prepared with potassium monobasic phosphate, sodium hydroxide and ultrapure water with pH adjustment to 6.8. All assays were performed in duplicate.

2.7 Acridine Orange/DAPI staining and data analysis The cells of L. paracasei encapsulated in electrospun fibers were stained with Acridine Orange - AO (1 mg/mL) for 20 min, which was mounted in a glass slide for posterior analysis under a fluorescence microscope. The material (cells and/or electrospun fibers) were analyzed using a Leica DM2000 epifluorescence microscope with a set of four filter cubes (A, L5, N3 and E4) and the images were captured with a Leica FX-350 camera using Leica QFish software. Polímeros, 30(3), e2020025, 2020

Investigation of Lactobacillus paracasei encapsulation in electrospun fibers of Eudragit® L100

3. Results and Discussions The morphology of L. paracasei, electrospun fibers of Eudragit® L100 (EDGT) and electrospun fibers with L. paracasei (EDGT-L.paracasei) were compared from SEM images, as shown in Figure 1 – the diameter was calculated from 25 different fibers per image. The EDGT fibers prepared in the absence of L. paracasei were uniform (with no imperfection) and presented a mean diameter of (2.134 ± 0.4127 µm) (Figure 1b). The morphology of pristine (non-encapsulated L. paracasei) is characterized by cells with an average length of (1.536 ± 0.370) µm and an average width of (0.526 ± 0.068) µm (Figure 1a). The incorporation of L. paracasei into electrospun fibers

is followed by the formation of imperfections localized along with the structure. As can be seen in Figure 1c, it is possible to identify aggregates of cells along with the polymeric structure. The EDGT-L.paracasei presented an average diameter of 1.508 ± 0.477 µm (Figure 1a). The EDGT-L. paracaseiloaded presented a reduction in diameter size (Figure 1d) and that the aggregates correspond to several L. paracasei cells, which individually evaluated show the same average diameter of L. paracasei (0.51 to 0.77 µm) (Figure 1d). The decrease in the average diameter of EDGT-L. paracasei samples may be due to the viscosity modification that results in reduced finer fiber size. Fung et al.[23] reported that PVA-based electrospun fibers with agrowaste containing

Figure 1. Scanning electron microscopy images are shown for (a), L. paracasei (L) (b), EDGT (c) EDGT-L.paracasei and (d) average diameter EDGT, EDGT-L.paracasei and L. paracasei – calculated from sets of 25 different fibers per SEM image. Polímeros, 30(3), e2020025, 2020

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Soares, J. M. D., Abreu, R. E. F., Costa, M. M., Melo, N. F., & Oliveira, H. P. L. acidophilus have decreased in mean diameter size by comparison with electrospun neat fibers due to the higher viscosity of pristine solution. An important aspect to be reported from these images is that there are no L. paracasei cells on the surface of electrospun fibers, in agreement with reported in the literature[27,28,40]. According to Heunis et al.[26], the high voltage applied during the electrospinning process inhibits bioactive cells not included in the electrospun fibers (non-protected species) that are prone to destruction in consequence of applied high voltage. In Figure 1c it is possible to observe that the L. paracasei cells were concentrated into the electrospun fibers as also randomly oriented along the electrospun fibers, such as reported in the literature[22,28]. In terms of the FTIR responses sample, it is possible to observe the presence of the peaks at 3507 cm-1 as a response of free carboxylic acid form and 3000 cm-1 and 2957 cm-1 assigned to vibrations CHx, 1723 cm-1 for esterified carboxylic groups and 1162 cm-1 given by carboxylic acid ester bonds stretching vibrations[41-43], see Figure 2. The spectrum of L. paracasei shows FTIR at positions that are in agreement with reported in the literature with the fingerprint of Lactobacillus spp. for peaks between 1300 and 900 cm-1 that indicate specific vibrational features of nucleic acids and bacterial proteins[40,44,45]. For samples EDGT-L. paracasei, the FTIR spectrum is characterized by peaks of the Eudragit® L100 polymer that stands out above the L. paracasei peaks. Ceylan et al. [24] reported similar results in which the peak assignments of each electrospun fiber component (pristine electrospun fibers and Lactobacillus loaded electrospun fibers) were highly similar to the individual components. Other studies have also reported the hardness visualization of Lactobacillus spp. peaks in electrospun fibers due to the complex overlap of polymer peaks and/or additives, showing that the evidence of encapsulation of Lactobacillus spp. was supported by the SEM images and fluorescence microscopies[28,40]. The staining procedure for the detection of cells (viable and killed organisms) is based on the interaction of fluorochromes and cells for the following identification of fluorescence levels in microscopy images[46]. AO and 4”

Figure 2. FTIR of EDGT, EDGT-L.paracasei and pristine L. paracasei. 4/8

6-diamidino-2-phenylindole (DAPI) are two of the most used fluorochromes in microbiology[47]. Particularly, the AO at low concentration binds with RNA and allows that microorganisms at high growth rates present high fluorescence at the red-orange region. On the contrary, for death cells, abundant DNA binds with AO and shifts the emission to the green region[48]. Based on these properties, it is possible to differentiate active cells´ signature (RNA fluorescence) from dead cells signature (prevailing DNA fluorescence)[49,50]. Assays of fluorescence microscopy were performed using AO and are summarized in Figure 3. The presence of free L. paracasei (viable cells) (from the fluorescence of AO at red region) can be visualized in Figure 3a. As shown, strong fluorescence reveals the viable character of cells before incorporation into the polymer solution. After encapsulation of microorganisms into electrospun fibers, it is possible to identify the presence of viable cells in fibers from strong red fluorescent dots (shown in Figures 3b, 3c and 3d) characterizing by the interaction of prevailing RNA of the viable cells and AO[50]. The fluorescence images are in agreement with previous SEM images that shown organisms dispersed as aggregates at specific sites of the fibers. The strong fluorescence reached form interaction with fluorochrome confirms the presence of viable cells encapsulated in electrospun fibers. It can be attributed to a covering layer of polymer that protects cells against the effects of high voltage and allows that a reasonable number of cells remain active after encapsulation. The confirmation of the presence of viable L. paracasei in the Eudragit® L100 electrospun fibers is already considered a notable result because to the best of our knowledge, there is no literature reporting the encapsulating of probiotics in electrospun fibers derived from Eudragit® L100 polymer solution in alcohol. Therefore, it is essential to test whether encapsulated L. paracasei are viable after fiber production and collection. For this, the L. paracasei-loaded Eudragit® L100 electrospun fibers were dissolved in the intestinal fluid solution for later plating. In the Figure 4 it is possible to see the growth of L. paracasei cells after electrospinning. Thereby, immediately after electrospinning of electrospun fibers still contained a high number of active L. paracasei as the survival rate of the viable cells, which is a remarkable result compared to the adverse effects of the electrospinning process and also of the polymeric solution used and tested for the first time for encapsulation of viable cells. The initial amount of L. paracasei added was 8.705 (log (CFU/mL)) and after electrospinning was 5.837 (log (CFU/mL)). The remaining viable cells confirm that technique and polymer template can be explored to encapsulate Lactobacillus species. Heunis et al.[26] assert that encapsulation in electrospun fibers has a high viability rate. Thus, electrospinning has been demonstrating a very useful and advantageous technique for the encapsulation of viable cells[22,27,34]. One of the factors may also have compensated the loss in viability that is the low diameter value of the fibers that favors the high surface to volume ratio of resulting material[26]. Besides that, the process of electrospinning may exclude environmental oxygen, Polímeros, 30(3), e2020025, 2020

Investigation of Lactobacillus paracasei encapsulation in electrospun fibers of Eudragit® L100

Figure 3. Fluorescence images for samples: (a) L. paracasei treated with AO, (b, c and d) EDGT-L.paracasei treated with AO.

Figure 4. Growth of L. paracasei cells (a) after the release of encapsulated species loaded on electrospun fibers (b).

contributing to the stability of Lactobacillus species[23]. Thus, although the decrease in water activity decreases during the electrospinning, a reduced oxygen level can improve storage stability[51]. In general, the effects observed in the literature for encapsulation of probiotics refer to the increase in the number Polímeros, 30(3), e2020025, 2020

of viable cells at prolonged contact with the simulated gastric fluid (SGF) and simulated intestinal fluid (SIF)[21]. Mojaveri et al.[19] reported that nanoencapsulation by electrospinning affects the number of viable cells (in 1-log reduction) as a result of extreme conditions (high electric field for synthesis). Despite this effect, the strong protection 5/8

Soares, J. M. D., Abreu, R. E. F., Costa, M. M., Melo, N. F., & Oliveira, H. P. provided by the fibers improves the viability rate of cells under SIF and SGF conditions. While a 3-log reduction is observed for free cells in an extreme environment, the production of the electrospun fibers reduces the number of viable cells from 8.37-8.44 log CFU/mL to 7.25-7.31 CFU/mL, revealing the potential of electrospun fibers. As shown from comparison with data reported in the literature, superior performance in terms of survivability degree of cells during the electrospun procedure and the lower reduction in the viable cells at SGF/ SIF conditions is reached for binary systems, such as introduced as Yilmaz et al.[17] that associated conventional polymer matrix and sodium alginate. The improvement in the survivability for experimental systems based on EDGT-based electrospun mats depends on adequate interaction of probiotics and polymer support with a third component that can be an alginate polymer, that acts as an extra layer to protect cells against high electric field and an barrier for controlled release under specific pH. The incorporation of additives for electrospun mats represents an important trend for this work, in a posterior step that tends to improve not only the retention of viable cells under electrospinning but also at adverse conditions (low pH) for long time assays.

4. Conclusions The electrospinning technique demonstrated to be a successful strategy applied in the encapsulation of L. paracasei by Eudragit® L100. Although some factors interfere in cell viability, the results revealed that Eudragit® L100 electrospun fibers offer a hydrophobic environment that provides adequate protection of L. paracasei cells against oxygen – preserving its viability. Viable cells were identified by fluorescence microscopy before and after release from controlled conditions, confirming that strategy of encapsulation of probiotics in enteric polymer-based electrospun fibers has been successfully established under acidic pH for following the release of viable cells – that remain protected against adverse conditions – and optimizing the characteristics of probiotics for prolonged action. In summary, these findings open new possibilities for use of a simple experimental system (alcoholic solution of EDGT) for encapsulation of probiotics in a promising template that can be enriched by the incorporation of additives such as sodium alginate in binary electrospun mats.

5. Acknowledgements This work was supported by Coordenação de Aperfeiçoamento de Pessoal de Nível Superior - CAPES, Financiadora de Estudos e Projetos – FINEP, Conselho Nacional de Desenvolvimento Científico e Tecnológico – CNPq and Fundação de Amparo a Ciência e Tecnologia do Estado de Pernambuco – FACEPE.

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Polímeros, 30(3), e2020025, 2020

ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.03320

Separation of PET from other plastics by flotation combined with alkaline pretreatment Fernando Pita1*  and Ana Castilho1  Departamento de Ciências da Terra, Faculdade de Ciências e Tecnologia, Centro de Geociências – CGEO, Universidade de Coimbra, Coimbra, Portugal

*fpita@ci.uc.pt

Abstract Plastics are naturally hydrophobic, so the selective flotation of plastic mixtures is impossible without changing their surface properties. The objective of this research is flotation separation of PET from PS, PMMA and PVC combined with sodium hydroxide pretreatment. NaOH pretreatment had a strong effect on PET’s floatability, medium effect on PMMA and PVC and limited effect on PS. The reduction of PET floatability is ascribed to a sharp decline of contact angle. The optimal pretreatment conditions were: 10% NaOH concentration, 70 °C (PET/PMMA and PET/PVC) to 80 °C (PET/PS) and plastic treatment times in the alkaline solution of 20 min (PET/PMMA and PET/PVC) to 30 min (PET/PS). Flotation separations were achieved efficiently and the best results were obtained for the PET/PS mixture, with a floated PS grade of 98% and a sunk PET grade of 100%. PET and PS was separated effectively, implying that sodium hydroxide treatment possesses superior applicability. Keywords: alkaline treatment, flotation, particle size, plastic. How to cite: Pita, F., & Castilho, A. (2020). Separation of PET from other plastics by flotation combined with alkaline pretreatment. Polímeros: Ciência e Tecnologia, 30(3), e2020026. https://doi.org/10.1590/0104-1428.03320

1. Introduction Since the discovery of plastic in the ´50s of the last century, its global production has been continuously rising, gradually replacing materials, like glass and metal. In the last decade, the world production of plastics has been grown around 3.5% per year, increasing from 230 million tonnes in 2005 to 359 million tonnes in 2017[1]. PET plays an important role in the packaging industry. The global PET production in 2017 was some 30.3 million ton, and in 2016, 485 billion PET bottles were produced worldwide. Despite the constant increase in plastic consumption none of the commonly used plastics is biodegradable. The vast majority of plastic waste ends up in landfills or the natural environment, or are incinerated, causing serious environmental problems. In 2015, around 55% of global plastic waste was discarded, 25% was incinerated, and 20% was recycled[2]. However, in Europe during 2017, 32.5% of plastic waste was recycled, 42.6% was recovered through energy recovery processes and 24.9% was landfilled[1]. Thus, it is urgent to substantially reduce the use of plastics and lessen plastic waste by recycling and reusing. However, in order to recycle plastic waste it is necessary to separate the plastic mixtures into individual plastics, because different plastics cannot be recycled together due to chemical incompatibilities, different melting points and thermal stabilities[3,4]. For example, in PET contaminated with PVC, PVC will degrade at the

Polímeros, 30(3), e2020026, 2020

PET processing temperature and produce char, leading to product discoloration[5]. Froth flotation, the most common separation process used by the mineral industry, is a possible alternative for separating plastic mixtures. Froth flotation allows the separation of hydrophobic and hydrophilic materials. However, since most plastics are naturally hydrophobic, it is necessary to selectively modify the plastic surface before plastic flotation. In the past few years, many modification methods were applied to plastic flotation, including wetting reagents[6-13] and surface chemical modification by treatment with chemical reagents, particularly with alkaline solution of NaOH[3,14-30]. Surface modification for plastic flotation aims to raise the hydrophilicity by introducing hydrophilic groups on plastic surface. In previous studies, it has been found that the alkaline solution of NaOH was able to destroy the hydrophobicity of PET, making it hydrophilic, whereas the hydrophobicity of others plastics was only slightly affected[14,16,17]. In this study, an alkaline treatment of PET, PVC, PS and PMMA particles with sodium hydroxide (NaOH) solutions followed by froth flotation was performed. Separation of PET from other plastics was achieved by froth flotation combined with NaOH pretreatment in bi-component mixtures of plastics. The mechanism of NaOH treatment was examined by contact angle measurements. The analysed parameters were: NaOH concentration, temperature, and treatment time of alkaline solution, and particle size.

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

Pita, F., & Castilho, A.

2. Experimental

2.3 Flotation experiments

2.1 Materials

The froth flotation experiments were performed in a Denver cell, with a capacity of 3 dm3, at a low rotational speed of 600 rpm. Each flotation test used 40 g of plastics, previously treated with NaOH, that was conditioned with the frother (MIBC) for about 2 minutes before the flotation tests. MIBC was added at a constant concentration of 30x10-3 g/L in all experiments. After conditioning, the air valve was opened and the floated product was collected over 6 minutes. Both the floated and the sunk (non-floated) were dried and weighed. Tap water was used in the flotation tests.

Four different kinds of post-consumer plastics were used: Polyethylene Terephthalate (PET) (colorless, transparent, lamellar); Polyvinyl Chloride (PVC) (light green to white, translucid, lamellar), Polystyrene (PS) (dark colored, translucid, irregular) and Polymethyl methacrylate (PMMA) (colorless to white, transparent, irregular) (Figure 1). The samples were previously ground and the sieve size fractions used in this study were +2-2.8 mm and +2.8-4 mm. The density of these plastics, measured by an Ultra Pycnometer (AccuPyc 1330), are as follows: PET: 1.364 g/cm3; PVC: 1.326 g/cm3; PS: 1.047 g/cm3; and PMMA: 1.204 g/cm3.

The pH in the flotation cell was not adjusted, but it was measured periodically along with the experiment. The pH remained approximately constant, in the range of 7.0-7.3.

Sodium hydroxide was used in the alkaline treatment as wetting agent, and Methyl isobutyl carbinol (MIBC) (109916 Sigma Aldrich) was used as frothing reagent.

Firstly, flotation tests were carried out with one-component plastic samples previously treated with the alkaline solution. According to the floatability of plastics, it was possible to separate the four plastics into two groups: the first group constituted only by PET, which has low floatability, and the second group that includes the other three plastics (PS, PMMA and PVC), which have similar floatability. Then, flotation separation of binary plastic mixtures was performed using three bi-component mixtures: PET/PS, PET/PMMA and PET/PVC. Plastic mixtures were previously treated with the alkaline solution, and each plastic contributed with 50% (20 g) for the total mixture weight.

2.2 Alkaline pretreatment Plastic samples were treated with NaOH solutions, using a Denver stirrer (400 rpm) with a plot plate. The treatment was done in a 1 L glass beaker using 40 g of plastic at a solids concentration of 20 wt%. The beaker was placed on a hot plate in order to adjust and control the temperature. The alkaline treatment was controlled by the operating parameters: NaOH concentration, temperature, and treatment time (Table 1). Plastic samples were treated in alkaline solutions at NaOH concentrations of 2%, 6% and 10%, at a temperature range between 20 °C and 80 °C and a treatment time between 2.5 and 30 min (Table 1). After alkaline pretreatment, the plastics were taken out from alkaline solutions, rinsed in a stream of tap water to remove the treatment solution and used to conduct flotation tests.

The effectiveness of the flotation tests was evaluated by the grade and recovery of each type of plastic in the floated and in the sunk products, and by the separation efficiency, defined by Schulz[31] as η =RP1-RP2 (where η is the separation efficiency, RP1 is the recovery of plastic 1 in the floated and RP2 is the recovery of plastic 2 in the floated). In the flotation tests of the plastic mixtures, the plastics type presented in

Table 1. Experimental range and levels of the independent variables for the alkaline treatment. Parameters NaOH concentration (%) Temperature (°C) Treatment time (min)

Symbol A

Range values and coded level () 10

(-1) 20

(0) 40

(+1) 70

(-1) 2.5

(-0.333) 5

(+0.666) 10

(+1) 20

(-1)

(-0.818)

(-0.455)

(+0.273)

(+1)

-1: factor at low level; 0: factor at medium level; +1: factor at high level.

Figure 1. Original pictures of the four plastics. 2/9

Polímeros, 30(3), e2020026, 2020

Separation of PET from other plastics by flotation combined with alkaline pretreatment the floated and the sunk were separated from each other by manual sorting, weighed, and flotation recovery and grade were calculated based on mass balance. This was possible due to the differences in colours and shapes of the plastics particles. Experiments were done three times under similar operating conditions. A second order polynomial equation was chosen to investigate the effect of different operating parameters of the alkaline treatment on the floatability of the plastics (Equation 1): Y = b0 +b1A+b 2 B+b3C+b12 AB+b13AC+ b 23BC+b11A 2 +b 22 B2 +b33C2

(1)

where, Y is the predicted response, b0 is model constant; b1, b2 and b3 are linear coefficients; b12, b13 and b23 are the interaction coefficients; and b11, b22 and b33 are the quadratic coefficients. This model represents the effect of NaOH concentration (A), temperature (B), treatment time (C) and their interactions on the plastics floatability. The list of the independent variables (A, B and C) with their coded and levels are presented in Table 1. The significance of model equation, individual parameters, and factor interactions were evaluated by analysis of variance (ANOVA) at the confidence intervals of 95% (α = 0.05).

2.4 Contact angle measurements There is a positive correlation between the hydrophobicity and the floatability, i.e., the flotation recovery increases

with the increase of the contact angle. Contact angles were measured in the Data Physics Instruments OCA20 equipment, using the sessile drop method. A drop of distilled water was placed onto the surface of plastic particles through a microsyringe and the contact angle was measured. This process was repeated five times for each plastic and the average value was considered to be the contact angle of the plastic.

3. Results and Discussion The four plastics (PET, PS, PMMA and PVC) are naturally hydrophobic, and the flotation recovery of untreated plastics is near 100%. Therefore, in order to separate plastic mixture by flotation, it is necessary to selectively modify the plastics floatability.

3.1 Effect of alkaline pretreatment on PET floatability Figure 2 shows the effects of NaOH concentration, temperature, and treatment time of the alkaline solution on the flotation recovery of PET of the two size fractions. NaOH treatment affects significantly the PET floatability. The flotation recovery of PET decreased with increasing NaOH concentration, temperature, and treatment time. Also, others studies[3,14,16-19,24,25] verified that recovery of PET in the floated decreased with increasing NaOH concentration, temperature, and treatment time. They found that alkaline treatment rendered the PET surface more hydrophilic, which may be a result of the hydrolysis of ester bonds in PET chains.

Figure 2. Influence of NaOH concentration, temperature, and treatment time of the alkaline solution on floatability of PET for fractions +2-2.8 mm and +2.8-4 mm. Polímeros, 30(3), e2020026, 2020

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Pita, F., & Castilho, A. The overall effect of particle size was minimal, since the results were similar for both size fractions (Figure 2). The reduction of PET flotation is more pronounced in strong alkaline solutions, under a long treatment time, and at an elevated temperature, since it substantially improves the kinetics of the hydrolysis reaction of the PET surface. So, the lowest recovery of PET in the floated (2.5% for fraction +2.8-4 mm and 2.8% for fraction +2-2.8 mm) was obtained with the highest NaOH concentration (10%), the highest treatment time (30 min) and the highest temperature (80 °C). The temperature of 20 °C had not effective impact in the PET hidrophlization, since its recovery is close to 100%. This situation can be explained because NaOH could not hydrolyze on the PET surfaces at this temperature. Other studies[14,16,23] verified similar behavior, saying that PET particles start to be affected by NaOH only when the temperature reaches about 30 °C. Also, for low NaOH concentration and short pretreatment time, PET floatability is still significant (Figure 2). Similar behaviour was found by Burat et al.[16] and Wang et al.[23]. At 40 °C and a 2% NaOH concentration, a change in treatment time had a low effect in PET recovery (Figure 2b). But for higher NaOH concentrations (6% or 10%), the flotation recovery of PET decreased significantly with increasing treatment time. At 70 °C and 80 °C (Figure 2c and 2d) the flotation recovery of PET decreased with increasing treatment time. At a temperature of 80 °C and for fraction +2.8-4 mm, the flotation recovery dropped from 98.6% to 34.5%, when the treatment time increased from 2.5 min to 30 min for a NaOH concentration of 2%. However, for higher concentrations of NaOH (10%), the hydrophilization of PET was achieved for lower treatment times. To find the effect of the NaOH concentration, temperature, and treatment time in PET recovery, statistical analysis of the experimental data was done and models were developed for optimization of the parameters. A quadratic relationship was found to describe the dependence of the plastic floatability on the three operating variables of the alkaline treatment. The equations presented are in terms of coded levels: the low parameter levels were coded as -1 and the high as +1 (Table 1). Therefore, the relative impact of the NaOH concentration, temperature, and treatment time in PET recovery can be identified by comparing the coefficients of the equation. The analysis of variance (ANOVA) for the PET recovery model of the two size fractions is shown in Table 2. Based on all statistical analysis, the model presented was considered adequate to predict PET floatability after alkaline treatment. The coefficients of determination (R2) obtained for the PET recovery of size fractions +2-2.8 mm and +2.8-4 mm were 0.8885 and 0.8881 respectively. The significance level

of each independent variable, as well as their quadratic terms and interaction between the variables, was evaluated based on corresponding F-values and p-values. For the two size fractions the model F-value was about 44 at 99.99% confidence level and the model Prob > F value is less than 0.05, shows that the model is significant. The quadratic effect of the three variables and the interaction between NaOH concentration and time treatment and between temperature and treatment time had no statistical significance. The actual model equation for fractions +2-2.8 mm and +2.8-4 mm are given in Equations 2 and 3, respectively: PET recovery (%) = 62.46 − 22.12A − 33.38B − 14.83C − 15.88AB + 0.31 AC − 3.54BC + 7.52A 2 − 3.84B2 + 1.83C2

PET recovery (%) = 61.83 − 22.10A − 33.70B − 15.10C − 15.27AB +0.50 AC − 5.31BC + 7.43A 2 − 3.78B2 + 4.38C2

(2) (3)

where A is the NaOH concentration (%), B the temperature (°C), and C the treatment time (min). The coefficients of the three parameters were negative values indicating a negative correlation between PET floatability and parameter levels. For both size fractions, PET recovery presented an equal order of relative impact of the operating parameters. The equation coefficients clearly showed that PET recovery was mainly affected by temperature, followed by NaOH concentration, treatment time, and interaction between the NaOH concentration and temperature.

3.2 Effect of alkaline pretreatment on PS floatability The floatability of PS was not influenced by alkaline pretreatment, since the results of the PS recovery in the float were always of 100%, suggesting that the alkaline pretreatment had not promoted the hydrophilization of PS. Also Wang et al.[23] verified that alkaline treatment has a low effect on PS floatability.

3.3 Effect of alkaline pretreatment on PMMA floatability Flotation recovery of PMMA decreased slightly with increasing NaOH concentration, temperature, and treatment time (Figure 3). The temperature had a considerable effect on the alkaline pretreatment of PMMA. At 20 °C, for the three NaOH concentrations and for all treatment times, there was no hydrophilization of PMMA, since PMMA recovery was about 100%. The effects of NaOH concentration, temperature, and treatment time of the alkaline solution on the flotation recovery of PMMA for the two fractions (+2-2.8 mm and +2.8-4 mm) were similar. However, PMMA recovery of fraction +2.8-4 mm was lower than that observed for fraction +2-2.8 mm. Also, other authors[6,15,32-34] found that large plastic particles were more difficult to float than smaller ones.

Table 2. Analysis of Variance (ANOVA) of the response surface quadratic model for PET recovery of the two size fractions. +2-2.8 mm

p-value Prob>F

model <0.0001

A <0.0001

B <0.0001

C <0.0001

AB <0.0001

AC 0.9635

BC 0.1169

A2 0.0588

B2 0.4583

C2 0.1557

Sum of Square=79831; Mean of square= 8870; degree of freedom=9; Fmodel=44.3; R2=0.8885; Ajusted R2=0.8680 +2.8-4 mm

p-value Prob>F

<0.0001

0.0002

<0.0001

0.8789

0.1242

0.0667

0.4283

0.2593

Sum of Square=79999; Mean of square=8889; degree of freedom=9; Fmodel=44.1; R2=0.8881; Ajusted R2=0.8678

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Polímeros, 30(3), e2020026, 2020

Separation of PET from other plastics by flotation combined with alkaline pretreatment

Figure 3. Influence of NaOH concentration, temperature, and treatment time of the alkaline solution on floatability of PMMA, for fractions +2-2.8 mm and +2.8-4 mm. Table 3. Analysis of Variance (ANOVA) of the response surface quadratic model for PMMA recovery of the two size fractions. +2-2.8 mm

p-value Prob>F

+2.8-4 mm

p-value Prob>F

model <0.0001

A <0.0001

B <0.0001

C <0.0001

AB <0.0001

AC 0.2060

BC 0.0001

A2 0.3750

B2 <0.0001

C2 0.086

Sum of Square=4316; Mean of square= 480; degree of freedom=9; Fmodel=102.4; R2=0.9485; Ajusted R2=0.9393 <0.0001

<0.0001

0.0002

0.0088

0.1793

0.3509

0.5321

<0.0001

0.088

Sum of Square=6645; Mean of square= 738; degree of freedom=9; Fmodel=128.0; R2=0.9584; Ajusted R2=0.9509

The lowest recovery of PMMA in the floated was obtained with the highest NaOH concentration (10%), the highest temperature (80 °C) and the highest treatment time (30 min), with 67.6% recovery for fraction +2-2.8 mm and 56.5% recovery for fraction +2.8-4 mm.

effect of NaOH concentration and treatment time had no statistical significance.

The analysis of variance (ANOVA) for PMMA recovery model of the two size fractions is shown in Table 3. For fractions +2-2.8 and +2.8-4 mm, the R2 values were 0.9485 and 0.9584, respectively, and the F-value at 99.99% confidence level and the Prob > F value were less than 0.05, showing the model goodness of fit.

PMMA recovery (%) = 89.08 − 3.65A − 11.63B − 4.64C −

For fractions +2-2.8 mm and +2.8-4 mm, the Equations 4 and 5, respectively, dictate that linear terms of the NaOH concentration, temperature, and treatment time, linear term of interaction between NaOH concentration and temperature, and quadratic term of the temperature had a negative effect on PMMA floatability of the two size fractions. Interaction between temperature and treatment time had a negative effect on PMMA floatability for the size fraction +2.8-4 mm. The linear term of interaction between NaOH concentration and treatment time, and the quadratic Polímeros, 30(3), e2020026, 2020

PMMA recovery (%) = 95.12 − 3.35A − 9.01B − 3.17C − 2.79AB − 0.32 AC − 1.98BC + 0.37A 2 − 5.23B2 + 0.58C2

1.30AB − 0.70 AC − 1.49BC + 0.41A 2 − 4.25B2 + 2.14C2

(4) (5)

For both size fractions, PMMA recovery presented an equal order of relative impact of the three independent variables. PMMA recovery was mainly affected by temperature, followed by NaOH concentration (A), treatment time (C) and interaction between the NaOH concentration and temperature (AB).

3.4 Effect of alkaline pretreatment on PVC floatability Floatability of PVC was influenced by NaOH concentration, temperature, and treatment time (Figure 4). However, this effect was much smaller than that observed for PET, and slightly smaller than that observed for PMMA. PVC recovery for fraction +2-2.8 mm was greater than that for 5/9

Pita, F., & Castilho, A.

Figure 4. Influence of NaOH concentration, temperature, and treatment time of the alkaline solution on floatability of PVC, for fractions +2-2.8 mm and +2.8-4 mm.

fraction 2.8-4 mm. Flotation recovery of PVC decreased slightly with increasing NaOH concentration, temperature, and treatment time. Thus, the lowest recovery of PVC in the floated was obtained with the highest NaOH concentration (10%), the highest temperature (80 °C) and the highest treatment time (30 min), with 87.8% recovery for fraction +2-2.8mm and 76.4% recovery for fraction +2.8-4 mm. These values are slightly lower than those observed by other authors[14,16,18,19,24], whose recoveries were close to 95%. At a temperature of 20 °C, for all NaOH concentrations and all treatment times, there was no hydrophilization of PVC, since PVC recovery was about 100%. Also, at 40 °C, the hydrophilization of the PVC was small, since the PVC recovery was close to 100%. The analysis of variance (ANOVA) for PVC recovery model of the two size fractions is shown in Table 4. For size fractions of +2-2.8 mm and +2.8-4 mm, the R2 values of 0.928 and 0.872 respectively, imply that the model fits well. The linear term of interaction between NaOH concentration and treatment time and the quadratic term of NaOH concentration and treatment time had no statistical significance. For fractions +2-2.8 mm and +2.8-4 mm, the Equations 6 and 7, respectively, dictate that NaOH concentration, temperature, and treatment time of alkaline pretreatment had a negative effect on PVC floatability of the two size fractions. Also, the interaction between NaOH concentration and temperature, interaction between temperature and treatment time, and quadratic term of temperature had a negative effect on the PVC floatability. 6/9

PVC recovery (%) = 99.01 − 0.75A − 2.64B − 1.29C − 0.97AB − 0.30AC − 1.33BC − 0.79A 2 − 1.81B2 + 0.59C2 PVC recovery (%) = 97.37 − 1.38A − 5.84B − 1.85C − 1.57AB − 0.30AC − 1.42BC − 0.08A 2 − 4.30B2 + 0.75C2

(6) (7)

PVC recovery was mainly affected by temperature (B). The three independent variables presented less impact on PVC recovery than on PET recovery.

3.5 Effect of alkaline pretreatment on contact angle Contact angle measurements were conducted to assess the effects of alkaline pretreatment on plastics surface. In the absence of alkaline pretreatment, all plastics untreated showed large contact angle, with PS presenting the highest value, followed by PVC, PMMA and PET. Contact angles measured in the the four plastics decreases with increasing NaOH concentration, temperature, and treatment time (Table 5). However, the contact angle of PET drops remarkably after alkaline treatment, while the contact angle of PS and PVC decrease only slightly. This is consistent with the results of flotation tests of single plastics. So, the dropping of flotation recovery of PET with alkaline treatment is ascribed to a significant decline of contact angle, which implies an increase on surface wettability. On the other hand, the contact angle of PS is larger than those of the other plastics in any treatment conditions, meaning that PS is the most hydrophobic plastic. Polímeros, 30(3), e2020026, 2020

Separation of PET from other plastics by flotation combined with alkaline pretreatment Table 4. Analysis of Variance (ANOVA) of the response surface quadratic model for PVC recovery of the two size fractions. +2-2.8 mm

p-value Prob>F

model <0.0001

A <0.0001

B <0.0001

C <0.0001

AB <0.0001

AC 0.1006

BC <0.0001

A2 0.8151

B2 <0.0001

C2 0.0651

Sum of Square=406; Mean of square= 45.1; degree of freedom=9; Fmodel=71.8; R2=0.9282; Ajusted R2=0.9153 +2.8-4 mm

p-value Prob>F

<0.0001

0.0005

<0.0001

0.0088

0.7543

0.0051

0.8912

<0.0001

0.3082

Sum of Square=1695; Mean of square= 188; degree of freedom=9; Fmodel=38.0; R2=0.8720; Ajusted R2=0.8495

Table 5. Experimental conditions and contact angle results. Conditions NaO concentration (%) 2 2 2 2 2 2 2 2 10 10 10 10 10 10 10 10

Contac angle (°)

Temperature (°)

Treatment time (min)

PET

PMMA

PVC

20 20 40 40 70 70 80 80 20 20 40 40 70 70 80 80

2.5 30 2.5 30 2.5 30 2.5 30 2.5 30 2.5 30 2.5 30 2.5 30

71.6 68.5 69.9 68.1 70.1 58.2 68.4 53.4 70.8 66.2 64.7 51.5 52.6 42.3 45.9 40.1

93.3 91.5 93.0 91.0 91.0 89.0 88.0 85.7 91.8 89.5 90.1 88.6 88.5 86.7 85.2 83.3

75.5 73.0 74.7 68.4 72.1 66.5 67.6 59.7 74.8 70.2 72.9 63.8 65.1 59.2 59.4 55.7

82.7 81.1 81.6 80.6 80.8 78.4 77.7 73.9 80.6 79.5 78.9 77.6 76.5 73.1 71.4 68.2

Previous results illustrated that alkaline treatment had a strong effect on PET floatability, some effect on floatability of PMMA and PVC, but no effect on PS floatability. Thus, the floatability of PET can be significantly reduced in hot alkaline solutions, showing smaller floatability than the other three plastics, particularly than PS. So, one can assume that the alkaline treatment is not efficient to separate PS, PMMA and PVC plastics from each other, but may allow the separation of PET from PS, PMMA and PVC.

3.6 Separation of bi-component mixtures of PET with PS, PMMA and PVC In face of the flotation behavior of single plastic, further alkaline treatment and flotation tests were developed using bi-component plastic mixtures of PET with PS, PMMA and PVC, in equal proportions, for two size fractions (+2-2.8 mm and +2.8-4 mm), in order to render the PET hydrophilic maintaining the other plastic components in a hydrophobic state. The alkaline treatment conditions chosen for each of the three bi-component mixtures (PET/PS, PET/PMMA and PET/PVC) were those that led to maximum flotation differences between PET and the other plastics, to obtain a selective separation. The optimal pretreatment conditions to separate PET/ PS mixture were: NaOH concentration of 10%, at 80 °C, and treatment time of 30 min. In these conditions, PET floatability was minimized, while PS floated recovery was 100%. The results of froth flotation tests are presented Polímeros, 30(3), e2020026, 2020

in Table 6. For PET/PMMA and PET/PVC mixtures, the optimal pretreatment conditions to separate PET from the other plastic were: NaOH concentration of 10%, temperature at 70 °C and treatment time of 20 min. The best result was obtained in the PET/PS mixture separation, having the highest separation efficiency (near 98%) and a sunk with a grade of 100% in PET and a floated with a grade of 98% in PS. On the other side, PET/PMMA mixture had the lowest separation efficiency. These results were consistent with the floatability of plastics observed in the mono-component tests. The influence of the particle size on separation quality of the PET/PS mixture was not evident, since the two size fractions presented similar results (Table 6). The effect of particle size on PET floatability was minimal. Coarse fraction of PET/PMMA mixture had the worst results. The separation was more efficient for the fine fraction because there was a great amount of PMMA recovered in the floated. For the PET/PVC mixture, the quality of separation worsened slightly with the increase of the particles size (Table 6). PVC recovery in the floated decreased with the increase of the particles size, and PET recovery in the floated was not affected by particles size. Floatability of PS, PMMA and PVC increased with decreasing particle size, because their particles have more regular shapes, while the effect of the particle size on PET 7/9

Pita, F., & Castilho, A. Table 6. Results of the flotation tests on the mixtures of PET with PS, PMMA and PVC for two size fractions. Plastic Mixtures

Fraction

PET/PS

+2-2.8

(mm)

+2.8-4 PET/PMMA

+2-2.8 +2.8-4

PET/PVC

+2-2.8 +2.8-4

Recovery (%) Products Floated Sunk Floated Sunk Floated Sunk Floated Sunk Floated Sunk Floated Sunk

Grade (%)

PET

OP*

PET

OP*

1.9 98.1 2.2 97.8 6.2 93.8 5.5 94.5 5.7 94.3 5.2 94.8

100 0 100 0 80.5 19.5 73.8 26.2 97.1 2.9 92.6 7.4

1.9 100 2.2 100 7.2 82.8 6.9 78.3 5.5 97.0 5.3 92.8

98.1 0 97.8 0 92.8 17.2 93.1 21.7 94.5 3.0 94.7 7.2

Separation Efficiency (SE) (%) 98.1 97.8 74.3 68.3 91.4 87.4

OP* denotes the other plastics, namely PS, PMMA or PVC.

floatability was minimal because their particles have lamellar shape and low weight. Also, other studies[15,32-34] verified that small particles and particles with lamellar shape are easier to float than coarse particles and particles with regular shape. For PET/PMMA and PET/PVC, the worst results for coarse fraction can be explained by the more regular shapes and higher densities of PMMA and PVC that hinders flotation. Thus, the particles size control is important for flotation separation of plastic mixtures.

4. Conclusions As four plastics (PET, PS, PMMA and PVC) are naturally floatable, it is necessary a selective wetting component to achieve a selective flotation separation of plastic mixtures. The effect of treatment with alkaline solutions of NaOH on the floatability of the four plastics was studied. It was verified that NaOH solutions had a strong influence on PET floatability, medium influence on PVC and PMMA and limited effect on PS. The contact angle measurement confirms the dropping of flotation recovery of PET is ascribed to a sharp decline of contact angle, which implies PET surface becomes hydrophilic. The flotation recovery of PET, PMMA and PVC decreased with increasing NaOH concentration, temperature, and treatment time. Sodium hidroxide treatment shows superior selectivity for PET plastic. Therefore, development of a selective flotation separation technology for PET/PS and PET/PVC mixtures can be successfully achieved. The best result was obtained in the PET/PS mixture separation. PET recovery in the sunk was about 98%, with a grade of 100%; and PS recovery in the floated was 100% with a grade of about 98%, with the following pretreatment conditions: 10% NaOH concentration, temperature at 80 °C, and treatment time of 30 min. For this mixture, the two size fractions (+2-2.8 mm and +2.8-4 mm) presented similar results. PET/PPMA mixture provided inadequate flotation and efficient flotation was obtained with PET/PVC. For these two mixtures, the quality of separation worsened slightly with the increase of the particles size as a consequence of the decrease of the recovery of PMMA and PVC in the floated for the coarser particles. 8/9

One particular advantage of the sodium hidroxide pretreatment is that in the flotation stage none wetting agents are used, and only a frother is required.

5. Acknowledgements This work was supported by the Portuguese Foundation for Science and Technology (FCT-MEC) through national funds and, when applicable, co-financed by FEDER in the ambit of the partnership PT2020, through the following research projects: UID/Multi/00073/2013 of the Geosciences Center of the University of Coimbra.

6. References 1. PlasticsEurope. (2019). Plastics - the Facts 2019. An analysis of European plastics production, demand and waste data. Brussels, Belgium. Retrieved in 2020, January 5, from https:// www.plasticseurope.org/en/resources/publications/1804plastics-facts-2019 2. Geyer, R., Jambeck, J. R., & Law, K. L. (2017). Production, use, and fate of all plastics ever made. Science Advances, 3(7), e1700782. http://dx.doi.org/10.1126/sciadv.1700782. PMid:28776036. 3. Kangal, M. O. (2010). Selective flotation technique for separation of PET and HDPE used in drinking water bottles. Mineral Processing and Extractive Metallurgy Review, 31(4), 214-223. http://dx.doi.org/10.1080/08827508.2010.483362. 4. Carvalho, M. T., Ferreira, C., Santos, L. R., & Paiva, M. C. (2012). Optimization of froth flotation procedure for poly (ethylene terephthalate) recycling industry. Polymer Engineering and Science, 52(1), 157-164. http://dx.doi.org/10.1002/pen.22058. 5. Guney, A., Poyraz, M. I., Kangal, O., & Burat, F. (2013). Investigation of thermal treatment on selective separation of post consumer plastics prior to froth flotation. Waste Management, 33(9), 1795-1799. http://dx.doi.org/10.1016/j. wasman.2013.05.006. PMid:23747135. 6. Shen, H., Pugh, R. J., & Forssberg, E. (2002). Floatability, selectivity and flotation separation of plastics by using a surfactant. Colloids and Surfaces. A, Physicochemical and Engineering Aspects, 196(1), 63-70. http://dx.doi.org/10.1016/ S0927-7757(01)00706-3. 7. Basarová, P., Bartovská, L., Korínek, K., & Horn, D. (2005). The influence of flotation agent concentration on the wettability Polímeros, 30(3), e2020026, 2020

Separation of PET from other plastics by flotation combined with alkaline pretreatment and flotability of polystyrene. Journal of Colloid and Interface Science, 286(1), 333-338. http://dx.doi.org/10.1016/j. jcis.2005.01.016. PMid:15848435. 8. Takoungsakdakun, T., & Pongstabodee, S. (2007). Separation of mixed post-consumer PET-POM-PVC plastic waste using selective flotation. Separation and Purification Technology, 54(2), 248-252. http://dx.doi.org/10.1016/j.seppur.2006.09.011. 9. Abbasi, A., Salarirad, M. M., & Ghasemi, I. (2010). Selective separation of PVC from PET/PVC mixture using floatation by tannic acid depressant. Iranian Polymer Journal, 19(7), 483-489. 10. Yenial, U., Kangal, O., & Güney, A. (2013). Selective flotation of PVC using gelatin and lignin alkali. Waste Management & Research, 31(6), 613-617. http://dx.doi. org/10.1177/0734242X13476748. PMid:23439876. 11. Yenial, U., Burat, F., Yüce, A. E., Güney, A., & Kangal, M. O. (2013). Separation of PET and PVC by flotation technique without using alkaline treatment. Mineral Processing and Extractive Metallurgy Review, 34(6), 412-421. http://dx.doi. org/10.1080/08827508.2012.702705. 12. Wang, C. Q., Wang, H., Fu, J. G., & Liu, Y. N. (2015). Flotation separation of waste plastics for recycling: a review. Waste Management, 41, 28-38. http://dx.doi.org/10.1016/j. wasman.2015.03.027. PMid:25869841. 13. Pita, F., & Castilho, A. (2019). Plastics floatability: effect of saponin and sodium lignosulfonate as wetting agents. Polímeros: Ciência e Tecnologia, 29(3), e2019035. http:// dx.doi.org/10.1590/0104-1428.01419. 14. Drelich, J., Payne, T., Kim, J. H., Miller, J. D., Kobler, R., & Christiansen, S. (1998). Selective froth flotation of PVC from PVC/PET mixtures for the plastics recycling industry. Polymer Engineering and Science, 38(9), 1378-1386. http:// dx.doi.org/10.1002/pen.10308. 15. Fraunholcz, N. (2004). Separation of waste plastics by froth flotation, review, part I. Minerals Engineering, 17(2), 261-268. http://dx.doi.org/10.1016/j.mineng.2003.10.028. 16. Burat, F., Güney, A., & Kangal, M. O. (2009). Selective separation of virgin and post-consumer polymers (PET and PVC) by flotation method. Waste Management, 29(6), 1807-1813. http://dx.doi.org/10.1016/j.wasman.2008.12.018. PMid:19155169. 17. Caparanga, A. R., Basilia, B. A., Dagbay, K. B., & Salvacion, J. W. L. (2009). Factors affecting degradation of polyethylene terephthalate (PET) during pre-flotation conditioning. Waste Management, 29(9), 2425-2428. http://dx.doi.org/10.1016/j. wasman.2009.03.025. PMid:19394808. 18. Carvalho, M. T., Durão, F., & Ferreira, C. (2010). Separation of packaging plastics by froth flotation in a continuous pilot plant. Waste Management, 30(11), 2209-2215. http://dx.doi. org/10.1016/j.wasman.2010.05.023. PMid:20576423. 19. Nagy, M., Škvarla, J., & Sisol, M. (2011). A possibility of using the flotation process to separate plastics. Annals of Faculty Engineering Hunedoara International Journal of Engineering, 9(3), 275-278. 20. Saisinchai, S. (2014). Separation of PVC from PET/PVC mixtures using flotation by calcium lignosulfonate depressant. Engineering Journal, 18(1), 45-53. http://dx.doi.org/10.4186/ ej.2014.18.1.45. 21. Wang, C. Q., Wang, H., Liu, Q., Fu, J. G., & Liu, Y. N. (2014). Separation of polycarbonate and acrylonitrile-butadiene-styrene waste plastics by froth flotation combined with ammonia pretreatment. Waste Management, 34(12), 2656-2661. http:// dx.doi.org/10.1016/j.wasman.2014.09.002. PMid:25266156.

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22. Güney, A., Özdilek, C., Kangal, O., & Burat, F. (2015). Flotation characterization of PET and PVC in the presence of different plasticizers. Separation and Purification Technology, 151, 47-56. http://dx.doi.org/10.1016/j.seppur.2015.07.027. 23. Wang, C. Q., Wang, H., & Liu, Y. N. (2015). Separation of polyethylene terephthalate from municipal waste plastics by froth flotation for recycling industry. Waste Management, 35, 42-47. http://dx.doi.org/10.1016/j.wasman.2014.09.025. PMid:25449606. 24. Guo, J., Li, X., Guo, Y., Ruan, J., Qiao, Q., Zhang, J., Bi, Y., & Li, F. (2016). Research on Flotation Technique of separating PET from plastic packaging wastes. Procedia Environmental Sciences, 31, 178-184. http://dx.doi.org/10.1016/j.proenv.2016.02.024. 25. Wang, C., Wang, H., Liu, Y., & Huang, L. (2016). Optimization of surface treatment for flotation separation of polyvinylchloride and polyethylene terephthalate waste plastics using responses surface methodology. Journal of Cleaner Production, 139, 866-872. http://dx.doi.org/10.1016/j.jclepro.2016.08.111. 26. Wang, C. Q., Wang, H., Gu, G. H., Lin, Q. Q., Zhang, L. L., Huang, L. L., & Zhao, J. Y. (2016). Ammonia modification for flotation separation of polycarbonate and polystyrene waste plastics. Waste Management, 51, 13-18. http://dx.doi. org/10.1016/j.wasman.2016.02.037. PMid:26965210. 27. Wang, J., Wang, H., Wang, C., Zhang, L., Wang, T., & Zheng, L. (2017). A novel process for separation of hazardous poly(vinyl chloride) from mixed plastic wastes by froth flotation. Waste Managemen, 69, 59-65. http://dx.doi.org/10.1016/j. wasman.2017.07.049. PMid:28801216. 28. Negari, M. S., Ostad Movahed, S., & Ahmadpour, A. (2018). Separation of polyvinylchloride (PVC), polystyrene (PS) and polyethylene terephthalate (PET) granules using various chemical agents by flotation technique. Separation and Purification Technology, 194, 368-376. http://dx.doi. org/10.1016/j.seppur.2017.11.062. 29. Wang, H., Zhang, Y., & Wang, C. (2019). Surface modification and selective flotation of waste plastics for effective recycling: a review. Separation and Purification Technology, 226, 75-94. http://dx.doi.org/10.1016/j.seppur.2019.05.052. 30. Zhang, Y., Jiang, H., Wang, H., & Wang, C. (2020). Separation of hazardous polyvinyl chloride from waste plastics by ﬂotation assisted with surface modiﬁcation of ammonium persulfate: process and mechanism. Journal of Hazardous Materials, 389, 121918. http://dx.doi.org/10.1016/j.jhazmat.2019.121918. PMid:31879107. 31. Schulz, N. F. (1970). Separation efficiency. Transactions of Society for Mining, Metallurgy, and Exploration, Inc., 247, 81-87. 32. Shen, H., Forssberg, E., & Pugh, R. J. (2001). Selective flotation separation of plastics by particle control. Resources, Conservation and Recycling, 33(1), 37-50. http://dx.doi. org/10.1016/S0921-3449(01)00056-8. 33. Pita, F., & Castilho, A. (2017). Separation of plastics by froth flotation: the role of size, shape and density of the particles. Waste Management, 60, 91-99. http://dx.doi.org/10.1016/j. wasman.2016.07.041. PMid:27478025. 34. Marques, G. A., & Tenório, J. A. S. (2000). Use of froth flotation to separate PVC/PET mixtures. Waste Management, 20(4), 265-269. http://dx.doi.org/10.1016/S0956-053X(99)00333-5. Received: Mar. 23, 2020 Revised: Aug. 07, 2020 Accepted: Aug. 14, 2020

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ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.05820

Water vapor permeation and morphology of polysulfone membranes prepared by phase inversion Luis Guilherme Macedo Baldo1 , Marcelo Kaminski Lenzi1  and Daniel Eiras1*  Programa de Pós-graduação em Engenharia Química – PPGEQ, Departmento de Engenharia Química, Universidade Federal do Paraná – UFPR, Curitiba, PR, Brasil

*eiras@ufpr.br

Abstract The aim of this work was to study the effect of different variables in the morphology and water vapor permeation of asymmetric membranes. Different from other works on vapor induced phase inversion this work focus on the formation of a dense skin capable of separating small molecules like gases and on the transport properties of water vapor instead of liquid water. It also correlates the morphologies with the permeability. The results show that higher polymer concentrations lead to denser skin and lower permeability. Water vapor transmission rates varied from 30 to 48 g/m2.h depending on membrane morphology. They also show that for membranes with the same type of skin layer the permeability depends on the sub-layer. Finally, the results suggest that different mechanisms were responsible for the formation of the membranes. Keywords: asymmetric membranes, vapor induced phase inversion, water vapor permeation. How to cite: Baldo, L. G. M., Lenzi, M. K., & Eiras, D. (2020). Water vapor permeation and morphology of polysulfone membranes prepared by phase inversion. Polímeros: Ciência e Tecnologia, 30(3), e2020027. https://doi. org/10.1590/0104-1428.05820

1. Introduction Membranes have been extensively studied and applied for several separations like gas and vapor separations, water production and purification, carbon capture and sequestration among others[1,2]. The application of membrane separation processes requires membranes with high flux and selectivity, characteristics that can be tuned not only by changing the properties of the materials that form the membrane, but also by changing the membrane geometry and morphology. In terms of membrane morphology, asymmetric membranes are usually desired. In the case of gas and vapor separations, a dense separation top layer is supported by a porous layer that is not supposed to impose any resistance to mass transfer. These membranes are produced by the phase inversion process that promotes phase separation of the polymer solution by contacting it with a non solvent (usually water)[3,4]. Phase inversion can be triggered by several different processes that include, for example, non solvent induced phase separation (NIPS)[5], vapor induced vapor separation (VIPS)[6] or temperature induced phase separation (TIPS). Due to its importance in membrane production, phase inversion process has been studied to determine the influence of different variables in the final morphology[7-10]. In these studies, different formation mechanisms were proposed to explain the formation of macrovoids, a cellular morphology, and interconnected structure and the dense layer. Among those mechanisms, two of the most recent studies have relied on rheological characterization of the polymer solutions to

Polímeros, 30(3), e2020027, 2020

correlate the final morphology with chain entanglement[11] and solvent power[12]. Depending on the rheological behavior of the solutions that is related to chain entanglement and solvent quality, the final membranes can have macrovoids or a cellular morphology. Because both studies focus on NIPS spinodal decomposition is not discussed. Other researchers have focused on VIPS trying to correlate the different variables to the final morphology and properties of the membranes[6,13-18]. The results showed that solution concentration, relative humidity[16] and time of exposure to the vapor[15] can influence the final morphology. It has also been reported that the choice of solvent can influence the final morphology by reducing the effect of coalescence of pores in membranes formed by spinodal decomposition[13]. The mechanism proposed by these authors also depends on rheological properties. Although most of these authors are interested in the porous structure and the application of membranes for liquid separations, their work shows that the coalescence of the pores in membranes formed by spinodal decomposition can create a very thin dense skin that can be applied for gas and vapor separations[13,15,19]. Recently, Dai et al.[20] proposed a process that combines VIPS and NIPS to produce membranes for gas separation. By combining the two processes they prevented the formation of pin-holes on the dense layer that can compromise the selectivity of the membrane. The authors argued that vapor induced phase separation can create a sublayer that reduces the mass transfer of both solvent from the membrane and non-solvent from the coagulation bath and protect the dense layer.

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Baldo, L. G. M., Lenzi, M. K., & Eiras, D. In this context, this work has the objective to characterize membranes prepared by vapor induced phase inversion combined with non-solvent induced phase inversion using different combinations of three variables: polymer concentration, water concentration in the solution and time inside the coagulation bath. The membranes were characterized by scanning electron microscopy (SEM) and water vapor permeation. The results show a correlation between the polymer concentration and water concentration in the solution with the final morphology. It is also possible to correlate the morphology and WVT . The morphology indicates that different mechanisms could explain the formation of the final morpohology. Finally, the results indicate that VIPS has a great potential to be applied in the production of asymmetric membranes with a dense skin for gas separation.

2. Materials and Methods 2.1 Materials Polysulfone Udel P-3500® was kindly supplied by SolvayBrazil. 1-methyl-2-pirrolidone (NMP-anhydrous, 99.5%) from Sigma-Aldrich was used as solvent for polysulfone. Ethanol (PA, 99.5%) and hexane (PA, 99.5%) from Synth were used for solvent exchange and tapped water was used as non-solvent and additive.

WVT =

2.2 Methods Polysulfone/NMP solutions were prepared by dissolving different concentrations of the polymer in the solvent. Polysulfone and NMP were mixed in a laboratory glass bottle that was closed and placed on a roller mixer. Complete dissolution was observed after 2 or 3 days depending on polymer concentration. The polymer concentrations used were 15, 25 and 35 wt%. Some solutions were prepared using water as additive. The water concentrations were 0, 2.5 and 5 vol.% in relation to the solvent initial volume. Table 1 shows the compositions and process conditions used to produce each sample. The membranes were prepared using a combination of non-solvent induced phase separation and vapor induced Table 1. Polymer concentrations and process variables used for membrane preparation. Sample 1 2 3 4 5 6 7 8 9 10 11 12 13

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phase separation. The polymer solutions were cast on a glass plate using a casting knife. Then the membrane was exposed to water vapor for 5 minutes inside an oven at 40 °C. The relative humidity inside the oven was kept between 40 and 60% during exposure. After the water vapor exposure, the membrane was immersed in the coagulation bath to complete phase inversion. The membranes were kept inside the coagulation bath for different periods of time varying from 10 to 60 minutes. Afterwards, the membranes were dried by soaking the membrane in ethanol for 24 hours and then soaking in hexane followed by drying in a vaccum oven at 68 °C for 24h. The morphology of the membranes was characterized by scanning electron microscopy analysis (SEM) in a JEOL JSM 6360-LV microscope. Criofractured samples were tested to evaluate the porous structure of the sublayer while the surface of the membranes was analyzed to verify the intregrity of the top layer. Water vapor permeation tests were conducted according to ASTM E96/E96M standard[21]. In a typical procedure, a cylindrical cup filled with water was covered by the membranes. The initial weight of the cups were measured and compared to their weight at different periods of time. Water vapor transmission (WVT) was calculated with Equation 1.

Polysulfone (PSF) concentration (wt%) 35 35 15 15 35 35 15 15 25 25 25 25 25

Time in coagulation bath (min.)

Water concentration (vol%)

60 10 60 10 35 35 35 35 60 60 10 10 35

2.5 2.5 2.5 2.5 5.0 0.0 5.0 0.0 5.0 0.0 5.0 0.0 2.5

G t. A

(1)

In Equation 1 G is the weigth change, t is the time and A is the permeation area.

3. Results and Discussions 3.1 Effect of water vapor activity Water vapor activity can be calculated by the ratio between the partial pressure of water and the equilibrium vapor pressure which makes the activity equivalent to the relative humidity. The activity of water vapor can influence the phase inversion in two different forms. First, the onset of VIPS is influenced by water vapor activity. It has been reported that phase separation induced by water vapor is observed for relative humidities of 65% or higher (aw ≥ 0.65)[16]. Second, the activity gradient through the membrane thickness influences mass transport of water and solvent exchange. Because the activity of water in the vapor phase is considered equivalent to the activity in liquid water, the differences between VIPS and NIPS are due to differences in mas transfer coefficients[6,16,17,19]. Different activity gradients can lead to different morphologies like symmetrical cellular morphology, asymmetrical cellular morphology or finger like pores. In the present work the relative humitidity range was between 40-60% which represents activities of 0.40-0.60. The change in humidity is due to the loss of water vapor that results from opening the oven to cast membranes and some variations in the oven temperature. Despite the lower activity there is evidence of phase separation by VIPS which could be due to the presence of water in the polymer solutions. Based on the morphologies of the membranes, we believe that the water vapor activity inside the oven might have influenced the onset Polímeros, 30(3), e2020027, 2020

Water vapor permeation and morphology of polysulfone membranes prepared by phase inversion of VIPS and the morphology of samples 4 and 10 leading to NIPS (Sample 10) or a combination of NIPS and VIPS (Sample 4). In the other samples it is more likely that the morphology was influenced by the mass transport of the non-solvent than the onset of phase separation. In that case, the range of relative humidity had a minor influence or no influence in the final morphology.

3.2 Morphology and phase inversion mechanism. 3.2.1 Typical morphologies obtained in this work. Figures 1-3 show the typical morphologies that were obtained in this work. Figure 1 shows the morphology of samples resulting from spinodal decomposition. This morphology is obtained when the membrane composition crosses the spinodal line[6,15], being an interconnected combination of pores and fibrils that looks like a deformed material. It is also obtained in the early stages of vapor induced phase inversion and usually coalesce over time to reduce the surface tension of the porous structure[13,15]. In the case of the membranes in Figure 1, coalescence is similar to previously reported results[13,15,16] considering the long time of vapor exposure (5 minutes) and it probably explains the lack of pore conectivity. Figure 2 presents the surface of these membranes, which show one dense surface and three porous surfaces. The dense surface is obtained due to coalescence of the previous spinodal structure that starts on the air (vapor)/ membrane surface and is favored by the increase of the

concentration of polymer in solution[15]. In Figure 2, the denser morphology was obtained from a 35 wt% solution. Figure 3 shows a cellular morphology and some interconnected pores that are characteristic of spinodal decomposition. The cellular morphology results from the slow diffusion of non-solvent through the cross section of polymer solution that makes the concentration gradient negligible in the cross section[6,17]. In the case of the membranes in Figure 3, the final morphology can be the result of spinodal decomposition that takes place in the some parts of the membrane and nucleation and growth to form the cellular morphology. The membrane surfaces in Figure 3 shows that a very thin dense layer was created which can reduce the diffusion of water vapor and create a cellular morphology specially at longer distances from the air interface[6]. Figure 4 shows typical symmetric cellular morphology. The symmetric cellular morphology without interconnected pores indicate that phase separation takes place while the solution composition is in the metastable region between spinodal and binodal[5,15,22]. In this case, the morphology results from nucleation and growth. At the same time, the dense layer in Figure 4d and the small pores in Figure 4c suggest that spinodal decomposition and coalescence took place closer to the air (water vapor) interface. Figure 5 shows the morphology of two membranes with macrovoids in their cross section. The morphology in Figure 5a seems to result from spinodal decomposition that has not reached the end of membrane thickness combined

Figure 1. Cross-section of spinodal morphologies. (a) Sample 5 (35 wt%PSF/35 min./5 vol.%water); (b) Sample 7 (15 wt%PSF/35 min./5 vol.%water); (c) Sample 8 (15 wt%PSF/35 min./0 vol.%water); (d) Sample 9 (25 wt%PSF/1 min./5 vol.%water). Polímeros, 30(3), e2020027, 2020

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Baldo, L. G. M., Lenzi, M. K., & Eiras, D.

Figure 2. Surface morphologies of membranes shown in Figure 1. (a) Sample 5 (35 wt%PSF/35 min./5 vol.%water); (b) Sample 7 (15 wt%PSF/35 min./5 vol.%water); (c) Sample 8 (15 wt%PSF/35 min./0 vol.%water); (d) Sample 9 (25 wt%PSF/1 min./5 vol.%water).

Figure 3. Morphologies of typical asymmetric cellular morphology. (a) and (c) Sample 6 (35 wt%PSF/35 min./0 vol.%water); (b) and (d) Sample 12 (25 wt%PSF/10 min./0 vol.%water). 4/8

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Water vapor permeation and morphology of polysulfone membranes prepared by phase inversion

Figure 4. Symmetric cellular morphology. (a) and (c) Sample 11 (25 wt%PSF/10 min./0 vol.%water); (b) and (d) Sample 13 (25 wt%PSF/35 min./1 vol.%water).

Figure 5. Cross section and surface of membranes having macrovoids. (a) and (c) Sample 4 (15 wt%PSF/10 min./2,5 vol.%water); (b) and (d) Sample 10 (25 wt%PSF/1 min./0 vol.%water). Polímeros, 30(3), e2020027, 2020

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Baldo, L. G. M., Lenzi, M. K., & Eiras, D. with phase inversion due to non solvent induced phase separation that resulted in elongated finger-like pores[6,10-12,15]. The fact that the surface of the membrane still have pores corroborates with the statement that spinodal decomposition took place. In Figure 5b, the morphology is almost entire formed by non solvent induced phase separation, i.e the water vapor exposure did not result in phase inversion and the final morphology was formed after immersion in water. The dense skin without apparent pores or pin holes is characteristic of immersion precipitation membranes. 3.2.2 Effect of polymer concentration on membrane morphology The polymer concentration in solution is an important variable in phase inversion processes[4,10-12]. Increasing the concentration increases the viscosity and reduces the rate of solvent/non solvent exchange[10]. The effect of viscosity relates to chain entanglement and relaxation times[11]. For PSF/ NMP systems it has been reported that there is a transition in zero shear rate viscosity with the increase in concentration. This transition is observed around 23 wt% of PSF in NMP and is characterized by a critical entanglement concentration above which the polymer behaves more like a solid than like a liquid[11]. According to that theory, phase separation will take place only when the polymer behaves like a liquid which means that for concentrated solutions the final morphology will depend on the relation between the relaxation time and the time for phase inversion. In the case of vapor induced phase separation at higher concentrations the distance between the binodal line and spinodal line increases which means that it takes more water vapor to take the solution concentration to the decomposition zone and for spinodal decomposition to start[15]. As a result spinodal decomposition might not be observed in concentrated solutions or more likely it will be observed on the surface of the solution that is close to water vapor. This last statement explains the effect of polymer concentration in the morphologies in Figure 6.

For the membrane prepared from a 15 wt% solution, it is possible to observe a deformed morphology that is characteristic of spinodal decomposition[6,15,17,18]. For higher concentrations the final morphology is a cellular morphology that is usually associated with nucleation and growth[6,15]. When analyzing the surfaces of these membranes it is clear that for the membrane prepared from the 15 wt% solution there is a series of pores that are formed due to coalescence of the spinodal morphology. For higher concentrations, it is possible to observe a denser layer which can only be formed by VIPS during coalescence of spinodal morphology[6,15]. The formation of a denser skin influences the final morphology of VIPS membranes because it reduces water vapor diffusion to the inside of the solution and keeps the solution in the meta stable zone for longer periods of time which favors the nucleation and growth mechanism over spinodal decomposition[15]. 3.2.3 Effect of water on membrane morphology The addition of water to the polymer solution is not recommended in real applications because it makes it difficult to dissolve the polymer but it can be used to evaluate other effects of the phase inversion process in membrane morphology. As a non solvent water increases the viscosity of the solution and also favors the formation of entanglement which reduces the critical entanglement concentration and increases relaxation time[11]. It is also expected that the presence of water in the polymer solutions reduces the amount of water vapor that is necessary to induce spinodal decomposition and also reduces the mass transfer rate of water vapor to the solution. Figure 7 shows the morphology of the membranes prepared with the addition of 5 vol% of water. The amount of water that was added to these solutions is enough to bring the concentration very close to the binodal line which means that a small amount of vapor would be

Figure 6. Effect of polymer concentration on membrane morphology. (a) and (d) Sample 8 (15 wt%PSF/35 min./0 vol.%water); (b) and (e) Sample 12 (25 wt%PSF/10 min./0 vol.%water); (c) and (f) Sample 6 (35 wt%PSF/35 min./0 vol.%water). 6/8

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Water vapor permeation and morphology of polysulfone membranes prepared by phase inversion

Figure 7. Effect of water in the morphology. (a) and (d) Sample 7 (15 wt%PSF/35 min./5 vol.%water); (b) and (e) Sample 9 (25 wt%PSF/1 min./5 vol.%water); (c) and (f) Sample 5 (35 wt%PSF/35 min./5 vol.%water). Table 2. Water vapor transmission (WVT) of Polysulfone membranes. WVT ± Standard Deviation

Sample Number

Morphology (Cross-section)

Morphology

Lacy -Spinodal Decomposition

Dense (very thin)

41 ± 3.4

Dense

27 ± 2.6

Porous

44 ± 3.5

Porous

44 ± 0.3

Porous

45 ± 2.1

Porous

41 ± 4.8

(Top-Layer)

9 2 6

Cellular Nucleation and Growth

(g/m2.h)

Dense (pin-holes)

Porous

41 ± 0.5

Dense (pin-holes)

37 ± 2.0

Dense

30 ± 2.0

Porous

48 ± 4.8

Dense

40 ± 2.1

4 10

Macrovoids NIPS

necessary to trigger phase separation. As a result of the incorporation of water to the solution we can observe the spinodal morphology in all the samples. These results suggest that the incorporation of water favors spinodal decomposition but it also reduces, but does not prevent, the effect of coalescence that is usually observed after long periods of exposure to water vapor specially in concentrated solutions. Although it is possible to observe that the effect of coarsening is reduced, the pore connectivity is still compromised which is expected for PSF/NMP solutions because the addition of water to the solution does not increase the viscosity enough to substantially reduce the domain growth rate[11,12]. From the surfaces of the membranes, it is observed that higher polymer concentration leads to denser skin layer which is a Polímeros, 30(3), e2020027, 2020

result of coarsening. Although that might seem controverse, it is important to understand that coarsening is more likely to take place closer to the vapor surface which explains why there are still signs of lacy morphology in the cross section of the membrane. The morphology in Figure 7f indicates that it is possible to obtain a dense skin layer using VIPS which is an important discovery for applications like gas and vapor separations.

3.3 Water vapor transmission of PSF membranes Table 2 shows water vapor transmission (WVT) of the membranes. The results are organized to group the membranes with similar morphology. The results suggest that the membranes with similar morphology have approximate the same WVT. It is possible to deduce from the resuls that the presence of a dense skin-layer decreases WVT. For the sample that presents a lacy like morphology and a porous surface WVT is approximately 45 g/m2.h (samples 7, 8 and 9) while sample 5 that has a dense skin-layer has a WVT of 27 g/m2.h. Samples 6, 7 and 10 are other examples of membranes with dense layer and lower permeability. Besides the effect of the dense layer the results show that the porous structure influences WVT. Based on the results the WVT increases in the sequence cellular morphology<lacy morphology<macropores. The different morphologies and their properties are an important result for the application of VIPS to produce asymmetric membranes for gas and vapor permeation. Usually, this method is used to produce porous membranes that are applied in liquid separation but the results suggest that it might be used to produce asymmetric membranes.

4. Conclusions Asymmetric membranes were produced by a combination of two different phase inversion methodologies and a variety of membrane morphologies were obtained depending on polymer and 7/8

Baldo, L. G. M., Lenzi, M. K., & Eiras, D. water concentration in the solution that formed the membranes. The results show that vapor induced phase separation (VIPS) can be applied to produce asymmetric membranes with a dense skin layer and that higher polymer concentrations favors the densification of the skin. The incorporation of water in the polymer solution seem to change the phase separation mechanism from nucleation and growth to spinodal decomposition for all the samples. Water vapor permeation tests show that there is a correlation between water vapor transmission and the membrane morphology. The results suggest that different phase inversion mechanisms result from different compositions in the solution and create different morphologies.

5. Acknowledgements The authors acknowledge the support from the Coordenação de Aperfeiçoamento de Pessoal de Nível Superior (CAPES) for the scholarship. From the Conselho Nacional de Pesquisa (CNPq) project 420696/2018-0, and Solvay Specialty Polymers for supplying the polymer Udel P-3500® that was used as matrix for membrane formation.

6. References 1. Kárászová, M., Zach, B., Petrusová, Z., Červenka, V., Bobák, M., Šyc, M., & Izák, P. (2020). Post-combustion carbon capture by membrane separation, review. Separation and Purification Technology, 238, 116448-116456. http://dx.doi.org/10.1016/j. seppur.2019.116448. 2. Nunes, S. P., Culfaz-Emecen, P. Z., Ramon, G. Z., Visser, T., Koops, G. H., Jin, W., & Ulbricht, M. (2020). Thinking the future of membranes: perspectives for advanced and new membrane materials and manufacturing processes. Journal of Membrane Science, 598, 117761-117788. http://dx.doi. org/10.1016/j.memsci.2019.117761. 3. Li, Y., Cao, B., & Li, P. (2019). Effects of dope compositions on morphologies and separation performances of PMDA-ODA polyimide hollow fiber membranes in aqueous and organic solvent systems. Applied Surface Science, 473, 1038-1048. http://dx.doi.org/10.1016/j.apsusc.2018.12.245. 4. Kausar, A. (2017). Phase inversion technique-based polyamide films and their applications: a comprehensive review. PolymerPlastics Technology and Engineering, 56(13), 1421-1437. http://dx.doi.org/10.1080/03602559.2016.1276593. 5. Hołda, A. K., & Vankelecom, I. F. J. (2015). Understanding and guiding the phase inversion process for synthesis of solvent resistant nanofiltration membranes. Journal of Applied Polymer Science, 132(27), 1-17. http://dx.doi.org/10.1002/app.42130. 6. Ismail, N., Venault, A., Mikkola, J. P., Bouyer, D., Drioli, E., & Tavajohi Hassan Kiadeh, N. (2020). Investigating the potential of membranes formed by the vapor induced phase separation process. Journal of Membrane Science, 597, 117601-117636. http://dx.doi.org/10.1016/j.memsci.2019.117601. 7. Ismail, A. F., & Lai, P. Y. (2003). Effects of phase inversion and rheological factors on formation of defect-free and ultrathin-skinned asymmetric polysulfone membranes for gas separation. Separation and Purification Technology, 33(2), 127-143. http://dx.doi.org/10.1016/S1383-5866(02)00201-0. 8. Guillen, G. R., Pan, Y., Li, M., & Hoek, E. M. V. (2011). Preparation and characterization of membranes formed by nonsolvent induced phase separation: a review. Industrial & Engineering Chemistry Research, 50(7), 3798-3817. http:// dx.doi.org/10.1021/ie101928r. 9. Xiang, J., Hua, X., Dong, X., Cheng, P., Zhang, L., Du, W., & Tang, N. (2019). Effect of nonsolvent additives on PES ultrafiltration 8/8

membrane pore structure. Journal of Applied Polymer Science, 136(15), 1-8. http://dx.doi.org/10.1002/app.47525. 10. Smolders, C. A., Reuvers, A. J., Boom, R. M., & Wienk, I. M. (1992). Microstructures in phaseinversionmembranes. 1. Formation of macrovoids. Journal of Membrane Science, 73(23), 259-275. http://dx.doi.org/10.1016/0376-7388(92)80134-6. 11. Hung, W. L., Wang, D. M., Lai, J. Y., & Chou, S. C. (2016). On the initiation of macrovoids in polymeric membranes: effect of polymer chain entanglement. Journal of Membrane Science, 505, 70-81. http://dx.doi.org/10.1016/j.memsci.2016.01.021. 12. Mousavi, S. M., & Zadhoush, A. (2017). Investigation of the relation between viscoelastic properties of polysulfone solutions, phase inversion process and membrane morphology: the effect of solvent power. Journal of Membrane Science, 532, 47-57. http://dx.doi.org/10.1016/j.memsci.2017.03.006. 13. Tsai, J. T., Su, Y. S., Wang, D. M., Kuo, J. L., Lai, J. Y., & Deratani, A. (2010). Retainment of pore connectivity in membranes prepared with vapor-induced phase separation. Journal of Membrane Science, 362(1-2), 360-373. http://dx.doi. org/10.1016/j.memsci.2010.06.039. 14. Peng, Y., Dong, Y., Fan, H., Chen, P., Li, Z., & Jiang, Q. (2013). Preparation of polysulfone membranes via vapor-induced phase separation and simulation of direct-contact membrane distillation by measuring hydrophobic layer thickness. Desalination, 316, 53-66. http://dx.doi.org/10.1016/j.desal.2013.01.021. 15. Su, Y. S., Kuo, C. Y., Wang, D. M., Lai, J. Y., Deratani, A., Pochat, C., & Bouyer, D. (2009). Interplay of mass transfer, phase separation, and membrane morphology in vapor-induced phase separation. Journal of Membrane Science, 338(1-2), 17-28. http://dx.doi.org/10.1016/j.memsci.2009.03.050. 16. Chae Park, H., Po Kim, Y., Yong Kim, H., & Soo Kang, Y. (1999). Membrane formation by water vapor induced phase inversion. Journal of Membrane Science, 156(2), 169-178. http://dx.doi.org/10.1016/S0376-7388(98)00359-7. 17. Tsai, H. A., Kuo, C. Y., Lin, J. H., Wang, D. M., Deratani, A., Pochat-Bohatier, C., Lee, K. R., & Lai, J. Y. (2006). Morphology control of polysulfone hollow fiber membranes via water vapor induced phase separation. Journal of Membrane Science, 278(12), 390-400. http://dx.doi.org/10.1016/j.memsci.2005.11.029. 18. Tsai, H. A., Lin, J. H., Wang, D. M., Lee, K. R., & Lai, J. Y. (2006). Effect of vapor-induced phase separation on the morphology and separation performance of polysulfone hollow fiber membranes. Desalination, 200(1-3), 247-249. http://dx.doi.org/10.1016/j.desal.2006.03.313. 19. Lee, H. J., Jung, B., Kang, Y. S., & Lee, H. (2004). Phase separation of polymer casting solution by nonsolvent vapor. Journal of Membrane Science, 245(1-2), 103-112. http://dx.doi. org/10.1016/j.memsci.2004.08.006. 20. Dai, Y., Li, Q., Ruan, X., Hou, Y., Jiang, X., Yan, X., He, G., Meng, F., & Wang, Z. (2019). Fabrication of defect-free matrimid® asymmetric membranes and the elevated temperature application for N2/SF6 separation. Journal of Membrane Science, 577, 258265. http://dx.doi.org/10.1016/j.memsci.2019.01.050. 21. American Society for Testing and Materials – ASTM. (2000). ASTM E96-00: standard test methods for water vapor transmission of materials (pp. 907-914). West Conshohocken: ASTM. http://dx.doi.org/10.1520/E0096-00. 22. Sadrzadeh, M., & Bhattacharjee, S. (2013). Rational design of phase inversion membranes by tailoring thermodynamics and kinetics of casting solution using polymer additives. Journal of Membrane Science, 441, 31-44. http://dx.doi.org/10.1016/j. memsci.2013.04.009. Received: May 21, 2020 Revised: July 20, 2020 Accepted: Aug. 17, 2020 Polímeros, 30(3), e2020027, 2020

ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.03420

Rheological and thermal properties of EVA-organoclay systems using an environmentally friendly clay modifiera Reinaldo Yoshio Morita1, Juliana Regina Kloss2, Ronilson Vasconcelos Barbosa3, Bluma Guenther Soares4,5, Luis Carlos Oliveira da Silva4 and Ana Lúcia Nazareth da Silva4,6*  Coordenação de Engenharia de Bioprocessos e Biotecnologia, Universidade Tecnológica Federal do Paraná – UTFPR, Dois Vizinhos, PR, Brasil 2 Departamento de Química e Biologia, Universidade Tecnológica Federal do Paraná – UTFPR, Curitiba, PR, Brasil 3 Departamento de Química, Universidade Federal do Paraná – UFPR, Curitiba, PR, Brasil 4 Instituto de Macromoléculas Professora Eloisa Mano – IMA, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brasil 5 Programa de Engenharia Metalúrgica – PEMM, Instituto Alberto Luiz Coimbra de Pós-graduação e Pesquisa de Engenharia – COPPE, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brasil 6 Programa de Engenharia Ambiental, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brasil a Selected paper presented at the 15th Brazilian Polymer Conference – (15thCBPol) held in Bento Gonçalves, Brazil, on 27–31 October, 2019. 1

*ananazareth@ima.ufrj.br

Abstract EVA systems, using an environmentally friendly organoclay modified with a non-ionic and free of ammonium ions modifier (BN-AM), were prepared in a single-screw laboratory extruder and characterized by rheological, morphological and thermal properties. WAXD analysis suggested that the nanocomposites with 1.5 wt% of BN-AM presented an exfoliated structure, while the rheological results showed that the nanocomposites with BN-AM organoclay tended to present a more pronounced shear thinning behavior when compared to EVA and the nanocomposites with the traditional organoclay. The SEM/EDS analysis by using elemental mappings showed good dispersion of the organoclays (BN-AM and BN-CT) in the EVA matrix. Thermogravimetry analysis showed an improvement in thermal stability of EVA when the non-ionic modifier was used instead of the traditional one. In general, it was concluded that the addition of low content of BN-AM organoclay in EVA matrices is a promising option for the production of nanocomposites. Keywords: EVA nanocomposite, organoclay, rheological property, thermal property, morphology. How to cite: Morita, R. Y., Kloss, J. R., Barbosa, R. V., Soares, B. G., Silva, L. C. O., & Silva, A. L. N. (2020). Rheological and thermal properties of EVA-organoclay systems using an environmentally friendly clay modifier. Polímeros: Ciência e Tecnologia, 30(3), e2020028. https://doi.org/10.1590/0104-1428.03420

1. Introduction Poly(ethylene-co-vinyl acetate) (EVA) clay nanocomposite thermoplastics have several industrial applications, such as greenhouse films[1,2], wires and cables[3,4]. The versatility of EVA nanocomposites is due to a synergic effect between the acetate group on the chains and the modified clay. Several studies have demonstrated improvements in mechanical properties[5-7], flame retardant[8-10], and gas barrier for organoclays[11-13]. Organoclay (organically modified clay) has been used since 1990 to prepare polymeric nanocomposites[5,14]. Mineral clays, especially montmorillonite (MMT), are natural hydrophilic agents present in bentonite. A main feature of MMT is swelling in water, due to the presence of mono- or divalent cations (Na+ or Ca2+) in the galleries, usually following an ion-exchange reaction involving

Polímeros, 30(3), e2020028, 2020

an organic modifier such as an alkylammonium or alkyl phosphonium surfactant. These organically modified clays contain organic species in the galleries, which change the basal spacing and the polymer-clay affinity, and they are importance to the morphology and properties of micro and nanocomposites[15-18]. Polymer clay nanocomposites prepared via a melting process are industrially convenient because this process is environmentally friendly and low-cost. However, it is not always possible to obtain a sufficient interaction between the polymer and the organoclay to achieve the desired level of exfoliated structures. Previous studies have indicated that the nature and alkyl-chain length of the modifier affect the structure of nanocomposites[19,20]. The rheological behavior of nanocomposites may reveal information about the relationship

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

Morita, R. Y., Kloss, J. R., Barbosa, R. V., Soares, B. G., Silva, L. C. O., & Silva, A. L. N. of the polymer chain and organoclays structure. Lee and Han[21] evaluated the rheological behavior of nanocomposites of systems based on two organoclays treated with a surfactant having polar and non-polar group and different gallery distances, in addition, they have chosen a polymer matrix having a wide range of polarity from partly polar EVA to highly polar poly(ethylene-co-vinyl alcohol) EVOH. The experimental results showed significant differences in the degree of exfoliation of organoclay and rheological responses in terms of the compatibility between polymer matrix and the surfactant residing at the surface of organoclay and the gallery distance of the organoclays. The authors pointed out that polymer melt viscosity plays an important role in dispersing the organoclay. The mainstream techniques for preparing polymerclay nanocomposites are extrusion and injection molding, which generate a high-shear environment for delaminating or exfoliating the clay platelets in the polymer matrix. The modifier may decompose during the process, at the typical melt-processing temperatures of some matrices. Usually, the onset temperature for MMT modified by alkylammonium ions is around 180 °C[22]. Therefore, other types of modifiers, such as non-ionic modifiers, are being used[23,24]. In a previous study, we examined the mechanical and thermal properties of the EVA nanocomposites containing non-ionic modified clays. The nanocomposites were prepared using two different natural clays and clay modifiers, ammonium ion and a non-ionic modifier. The results clearly showed that the non-ionic modifier behaved similarly to the ammonium-salt modifier. However, the organically modified clays proved to be more compatible with the polymer matrix than did the natural clay[25]. Carli et al.[26] studied nanocomposites reinforced by a nonionic organoclay, which showed improvement in properties. The authors compared the morphology and the thermal and mechanical properties of nanocomposites prepared using commercial organically modified montmorillonite, raw montmorillonite and non-ionic organoclay. They found that the nanocomposite using non-ionic surfactants was more chemically stable than the commercial cationic organoclays. The novelty of the present work is the addition of a new organoclay synthetized with a non-ionic surfactant, free of ammonium salt, which makes it more environmentally friendly[27-29]. BN-AM was added to an EVA matrix and the system characterized by WAXD, rheological and thermal analyzes and the results obtained were compared to similar systems containing a traditional organoclay.

2. Materials and Methods 2.1 Materials EVA (3019 PE, melt flow index 2.5 g 10 min-1, at 190 °C, 2.16 kg, 18 wt% vinyl acetate content) was supplied by Braskem. The non-ionic organoclay was prepared by modifying the sodium bentonite Vulgel CN 45 with a nonionic surfactant containing 18 carbon atoms and supplied by Ioto International[28,29]. The organoclay modified with cetyltrimethylammonium bromide (CTAB) was prepared 2/9

Table 1. Identification of the types of modifier and organoclays. Bentonite Vulgel CN 45

Type of Modifier non-ionic cetyltrimethylammonium bromide

Designated BN-AM BN-CT

by the ion-exchange method, using the same sodium bentonite (Table 1). Both organoclays and sodium bentonite Vulgel CN 45 had their physical and chemical properties characterized and disclosed in previous work[27]. In that work, the authors reported that the BN-AM showed the lower iron content and it was considered, among the other studied organoclays, the more promising one for nanocomposite use in nanocomposite preparation. Thus, in the present study, the BN-AM organoclay was chosen to be added to EVA matrix and the materials produced were compared to EVA/BN-CT compositions.

2.2 Methods 2.2.1 Composite preparation by melt processing EVA composites with 1.5 and 3.0 wt% were prepared with a MH Equipment MH 50H model intensive homogenizer operating at 3,600 rpm for 10 s. And, melt mixture was processed in a single-screw laboratory extruder, BGM EL-45 model, L/D = 40:1 and D = 25 mm. The temperature profile used was 135/140/150 °C at 200 rpm. After, the extruded was pelletized. The pellets obtained were injection molded using a Romi Primax 65R injector with a temperature profile ranging from 160 to 180 °C in the 4 zones with a mold for tensile test specimen (ASTM D638 Type 1). The injection pressure was 300 bar and the holding pressure was 150 bar, for 2 s. 2.2.2 Characterization The dynamic melt rheological properties of the neat EVA and the EVA composites were measured in order to gain a fundamental understanding of the processability and the relationship between the structure and the properties of the materials. Melt rheological measurements were performed on a TA Instruments TRIOS Discovery HR-1 rheometer in parallel-plate geometry, using 25 mm-diameter parallel plates. All tests were conducted at 180 °C. The linear viscoelastic zone was assessed by performing strain sweep tests from 0.1 to 100% at 1 Hz. Frequency sweep tests from 0.01 to 600 rad s-1 were performed at 1% strain under nitrogen atmosphere. The rheological analysis was conducted from a piece cut out from the tensile test specimen. Wide angle X-ray diffraction analysis (WAXD, Shimadzu XRD-6000) was operated at 40 kV/30 mA with Cu Kα radiation (wavelength, λ=0.154 nm). The samples were scanned with the diffraction angle 2θ, ranging from 2° to 50° at a scan rate of 1.0° min-1 in 2θ. Scanning electron microscopy (SEM) and energy dispersive spectroscopy (EDS) of the cryogenic fractured surfaces of the samples were performed on a JEOL, 1200 EX model and Bruker, Quantax 70 model. The thermal stability of the composites was evaluated by thermogravimetric analysis (TGA), with a Netzsch TG 209 thermobalance operating from 20 to 860 °C, with the following atmosphere programming: 20 to 560 Polímeros, 30(3), e2020028, 2020

Rheological and thermal properties of EVA-organoclay systems using an environmentally friendly clay modifier °C with nitrogen and 560 to 860 °C with synthetic air, using a heating rate of 10 °C min-1.

3. Results and Discussions 3.1 WAXD analysis WAXD analysis revealed an intercalated/exfoliated morphology. In a WAXD pattern, the peak that corresponds to the clay basal spacing (d001) is between 2° and 10°. The complete disappearance of this peak indicates that the stacked layers of the clay were exfoliated by the polymer chains, and the shift of the peak to a lower 2θ angle indicates that the stacked layers of clay were intercalated with the EVA chains. The preservation of the diffraction peak in the nanocomposites indicates that a significant clay content remained in layered form after processing. Figure 1 shows the WAXD patterns of the BN-AM organoclay and the EVA composites. The BN-AM organoclay showed diffraction peaks at 2θ that corresponded to d-spacing 19.6, 15.5 and 13.2 Å. Especially in the EVA composite containing 1.5 wt% BN-AM,

these peaks disappeared, indicating a favorable disordered structure of the organoclay in the polymer matrix, which may account for the tendency to produce exfoliated-clay domains. In the case of the higher organoclay content (3.0 wt%), a single displacement peak occurred, indicating that most of the organoclay platelets were evenly stacked[30,31]. This partial destructuring of the organoclay in the matrix may be associated with the models proposed by Lagaly[31] for the possible structural arrangements of the modifiers in the interlayer region of the clays, since BN-AM presents a mixture of mono-, bi-layer lateral arrangements are confirmed by the absence of a single diffraction peak of up to 10º (2θ) (Figure 1). Thus, the increase in the organoclay content has somehow hindered the process of intercalation of the polymer chains in the interlamelar region of the clay, not favoring exfoliation in the EVA matrix. In comparison of the patterns of the BN-CT organoclay (Figure 2) revealed a diffraction peak in the composite containing 1.5 wt%, corresponding to 18.2 Å. Only a slight shift of the diffraction peak occurred in this composition. This behavior does not mean that there was good intercalation of the BN-CT layer with the polymer chains, as probably

Figure 1. WAXD patterns of BN-AM organoclay, neat EVA and EVA/BN-AM composites.

Figure 2. WAXD patterns of BN-CT organoclay, neat EVA, and EVA/BN-CT composites. Polímeros, 30(3), e2020028, 2020

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Morita, R. Y., Kloss, J. R., Barbosa, R. V., Soares, B. G., Silva, L. C. O., & Silva, A. L. N. tactoids also appeared in the structure. These structures are termed “immiscible” because the organoclay platelets exist even after melting processing. Therefore, only EVA/BN-CT microcomposites were formed[5].

3.2 Small amplitude oscillatory shear flow: frequency sweep and SEM/EDS morphology Oscillatory shear-flow analyses were performed in two configurations: (i) frequency sweep and (ii) stress relaxation. The experiments were conducted as a function of time at 180 °C and a strain of 1%, defined from strain sweep test conducted to determine the linear viscoelastic region of the materials. The variation of the complex shear viscosity (η*) as a function of the frequency of the neat EVA and the EVAorganoclay composites (EVA/BN-CT and EVA/BN-AM) is shown in Figure 3. As seen in Figure 3, the neat EVA and all the composites showed a shear-thinning behavior, with an increase in viscosity in the low-frequency region. Figure 3a shows that when the BN-CT organoclay was added to the EVA matrix, this matrix behaved similarly to neat EVA, however, when the BN-AM modified clay was incorporated into the EVA matrix (EVA/BN-AM 1.5), the material produced tended to show a stronger frequency-thinning comparing to neat EVA and EVA/BN-CT 1.5 composite. When 3.0 wt% BNAM was added to the EVA matrix (Figure 3b), the final composite tended to show an even more shear-thinning behavior when compared to EVA/BN-AM 1.5 composite. The frequency dependency of the storage modulus (G’) of the materials obtained in the frequency (ω) sweep analysis is shown in Figure 4.

According to Jiang et al.[32], information on the G’ value variation with frequency can indicate polymer chain structure and dynamics. Based on the present experimental data, neat EVA deviated from standard terminal behavior, showing G’ α ω0.73 (R2=0.97). When 1.5 wt% and 3.0 wt% contents of BN-CT organoclay were added to EVA matrix, the relationships between G’ and ω were also G’ α ω0.73 (R2=0.97), the same relationship found to neat EVA, indicating a similar rheological behavior between EVA/BN-CT systems and neat EVA. When 1.5 wt% and 3.0 wt% contents of BN-AM organoclay were added to EVA matrix, the relationships between G’ and ω were G’ α ω0.68 (R2=0.97) and G’ α ω0.70 (R2=0.96), respectively, indicating an improved solid-like behavior. The variation of dynamic modulus, G’ and G”, was also evaluated. Table 2 presents the values of the crossover point, at which G’ = G”, obtained from G’ and G” versus the frequency curves. The results were evaluated by the displacement of the G’ x G” crossover point, which allows the flow behavior of polymeric materials to be predicted. The data in Table 2 obtained from the results for G’ = G” versus frequency show that the BN-AM compositions showed the lowest modulus values at the crossover point compared with the other EVA nanocomposites. For the EVA/ BN-AM composites, a decrease in the G’ = G” points were observed, but with similar frequency values at the crossover point compared to the neat EVA. This may be related to the more shear-thinning behavior observed. However, these data are not in accordance with the exfoliation-like morphology observed in the EVA/BN-AM 1.5 composite, as showed in Figure 1. It was expected that a more elastic behavior at lower frequency range would be observed and, as the frequency increased, a more pronounced shear thinning had occurred. Table 2. Dynamic modulus and frequency values at the G’/G” crossover point for neat EVA and EVA/organoclay composites. Sample code EVA EVA/BN-CT 1.5 EVA/BN-CT 3.0 EVA/BN-AM 1.5 EVA/BN-AM 3.0

Figure 3. Frequency sweep results for neat EVA and EVA/ organoclay systems (a) 1.5 wt% organoclay and (b) 3.0 wt% organoclay. 4/9

Modulus at crossover point

Crossover point

G’ = G” (104 Pa)

ωc (rad s-1)

1.81 2.21 1.87 1.41 1.35

10 16 16 10 10

Figure 4. Storage modulus G’ versus frequency curves for neat EVA and EVA/organoclay systems. Polímeros, 30(3), e2020028, 2020

Rheological and thermal properties of EVA-organoclay systems using an environmentally friendly clay modifier The addition of 3.0 wt% of the BN-AM organoclay further decreased the G’ = G” values, indicating that a nanocomposite with more frequency-thinning was produced. This behavior might be related to the presence of a mixture of BN-AM arrangements in the EVA matrix, which probably favored the chain flow and it is in accordance with Figure 1. The displacement of the G’ x G” crossover point results also show that the addition of BN-CT organoclay in EVA matrix led to a pronounced increase in viscous behavior of the final composite. It is well known that nanofilled polymers present a viscosity increase with filler content, although in some cases a decrease of the viscosity can occur. According to La Mantia et al.[33], the decrease of the viscosity can be explained by the occurrence of two possible mechanisms: (i) lubricant effect of the platelets when the clay particles are completely exfoliated or (ii) the occurrence of a low compatibility between polymer matrix and organoclay. It was also found from Figure 1 that the peaks disappeared in

EVA nanocomposite with 1.5 wt% of BN-AM, indicating the presence of exfoliated-clay domains. However, the possibility of organoclay agglomeration in the polymer matrix should also be considered. If this happened, a small clay content would occupy much less area in the polymer matrix, and so, the results observed in WAXD pattern would be leading to erroneous conclusions. SEM/EDS analysis is a characterization tool useful to investigate the dispersion of nanofiller in a polymer matrix[34,35]. Thus, SEM/EDS analysis was performed to evaluate the dispersion of the organoclays in the EVA matrix. The dispersion of the organoclay (BN-AM 1.5 wt% composite) in EVA matrix were evaluated by SEM/EDS analysis (Figure 5). As can be seem in Figure 5, BN-AM organoclay is well dispersed in EVA matrix, thus, the hypothesis of the occurrence of clay agglomeration, which could lead to

Figure 5. SEM/EDS micrographs and EDS results of BN-AM 1.5 composite.

Figure 6. SEM/EDS micrographs and EDS results of BN-AM 3.0 composite. Polímeros, 30(3), e2020028, 2020

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Morita, R. Y., Kloss, J. R., Barbosa, R. V., Soares, B. G., Silva, L. C. O., & Silva, A. L. N. misinterpretations in WAXD’s analysis can be discarded. Comparing SEM/EDS analysis of BN-AM 1.5 composite with that of BN-AM 3.0 composition (Figure 6), it can also be observed that even at higher BN-AM content, the dispersion of the filler in EVA matrix seems to be good. On the other hand, WAXD analysis showed a stacked morphology of the organoclay platelets on this composition (BN-AM 3.0).

which showed that neat EVA and BN-CT composites (BN-CT 1.5 and BN-CT 3.0) present similar flow behavior. These results are also in accordance with the WAXD analysis, which indicated that a clay fraction might remain in layered form in the BN-CT composites.

The results show that the tendency of the increase of the viscous behavior can be related to the lubricant effect of the clay platelets and that this effect should occur in a more pronounced way when the organoclay platelets are evenly stacked (BN-AM 3.0 composite).

At a constant deformation, ε0, and at some initial point in time, t=0, a slow decay of stresses over time σ(t) can be observed. This phenomenon, called stress relaxation, can be represented by a continuous spectrum. The relaxation properties of a polymer melt are given by the relaxation modulus function, G(t), which is related to the continuous spectrum (H(λ)) as follows:

Figure 7 and 8 show SEM/EDS analysis of BN-CT 1.5 and BN-CT 3.0 composites, respectively. The micrographs show that BN-CT organoclay are also well dispersed in the matrix, although the composition with 1.5 wt% of BN-AM seems to be better dispersed in the matrix when than the BN-CT 1.5 composite. For composites based on BN-CT organoclay it seems that despite the good dispersion of BN-CT in EVA matrix, probably there is a poor interaction between nanofiller and matrix. This hypothesis corroborates the rheological results,

3.3 Stress relaxation

∞

( )

G ( t ) = ∫ H( λ ) exp -t  d(lnλ ) λ   -∞

(1)

where λ is the relaxation time[36,37]. The stress relaxation experiments were conducted at a constant strain of 1% and at a temperature of 180 °C. Figure 9 illustrates the variation of the relaxation modulus over time of the neat EVA and the EVA/organoclay composites.

Figure 7. SEM/EDS micrographs and EDS results of BN-CT 1.5 composite.

Figure 8. SEM/EDS micrographs and EDS results of BN-CT 3.0 composite. 6/9

Polímeros, 30(3), e2020028, 2020

Rheological and thermal properties of EVA-organoclay systems using an environmentally friendly clay modifier As seen in Figure 9, the G(t) values of EVA dropped rapidly and vanished for a long period at around 500 s (liquid-like behavior), whereas the composites tended to extend for a long period and reached a finite pseudo-plateau (Figure 9a). A different behavior was observed for the BN-AM composites with lower clay content (1.5 wt%). A transition zone appeared, followed by a discreet plateau zone in which the modulus was nearly constant. At longer times, flow occurred and the G(t) curves moved into a “terminal zone”, where the modulus relaxed to zero after a sufficiently long time. WAXD analysis indicated that this composition seems to have an exfoliated morphology. At short times, the network-like structure should inhibit polymer chain flow, but at longer times, this structure should disappear and flow would occur. Figure 9b shows a different behavior. In the composite with 3.0 wt% BN-AM, the G(t) values dropped more rapidly compared to the neat EVA and the composite with BN-CT organoclay. This may be associated with the lubricant effect showed in the EVA/BN-AM 3.0 composite, as mentioned before.

3.4 Thermogravimetry analysis The TGA thermograms for the EVA, BN-AM and BN-CT composites show a profile well established by two subsequent mass losses (Figure 10). In all composites, the presence of the organoclays changed the initial temperature of the first polymer-degradation step. The values of the BN-CT composites decreased compared to EVA, indicating a slight change in the thermal stability of the polymer. This degradation step is attributed to deacetylation of the EVA chain, from 200 to 400 °C[38,39]. Based on the literature data, modifiers based on alkyl ammonium ions degrade at a lower temperature, and this temperature is easily achieved during preparation by the melting method, causing the polymer to degrade[22]. One of the ways to explain the reported degradation of polymers by these chemical species is that the surfactant modifier present in the interlayer region of the clay is released during processing[40]. Addition of BN-AM improved the thermal stability of the first step in EVA degradation, demonstrating that the modifier promotes slightly more stability than alkyl ammonium ion.

Figure 9. Variation of the relaxation modulus over time for neat EVA and EVA/organoclay systems (a) 1.5 wt% organoclay and (b) 3.0 wt% organoclay.

Figure 10. TGA thermograms of neat EVA, BN-AM (a) and BN-CT (b) composites. Mass loss derivative curves (DTG).

Table 3. TGA results for neat EVA and EVA/organoclay composites. 1st degradation step Sample code

TONSET

TMAX

2nd degradation step TENDSET

TONSET

(°C) EVA EVA-BN-CT 1.5 EVA-BN-CT 3.0 EVA-BN-AM 1.5 EVA-BN-AM 3.0

242 236 236 256 256

Polímeros, 30(3), e2020028, 2020

358 356 357 370 368

TMAX

TENDSET

(°C) 399 397 400 400 400

399 397 400 400 400

481 460 458 475 475

524 523 523 559 548

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Morita, R. Y., Kloss, J. R., Barbosa, R. V., Soares, B. G., Silva, L. C. O., & Silva, A. L. N. The TENDSET values of the second degradation step confirm this behavior (Table 3). This degradation step corresponds to cleavage of the polymer backbone and according to Beltrán et al.[41] when the clay mineral layers are exfoliated, they reduce the volatilization of the degradation products due to the barrier properties of the organoclay, which leads to an increase in thermal stability. TGA results of the BN-AM composites can be associated with better dispersion and interaction of the organoclay with the polymer in agreement with the WAXD patterns and rheological results.

4. Conclusions We studied the effect of the modification of clay on the rheological properties of the EVA matrix, using a new nonionic modifier agent (AM). WAXD diffractograms showed that only the nanocomposites based on EVA/BN-AM, with 1.5 wt%, did not show a diffraction peak, indicating that in this composition, stacked layers of clay were exfoliated by EVA chains when AM was used as a modifier, instead of the conventional ammonium salt modifier. However, when the EVA/BN-AM organoclay content of the EVA matrix was increased, small diffraction peaks started to appear, indicating that in this composition, a certain amount of EVA/BN-AM organoclay might have remained in layered form after processing. These results are in accordance with the observed rheological behavior. The thermogravimetry analysis showed an increase in the initial temperature of the polymer degradation process, which may be associated with an improved thermal stability of the EVA/BN-AM composite. Finally, this study allowed the evaluation of the rheological and thermal properties of EVA systems containing an alternative to the organoclay containing a new non-ionic modifier and not based on quaternary ammonium, and concluded that the addition of low BN-AM content is a promising option for the production of nanocomposites.

5. Acknowledgements The authors are grateful for financial support from the National Council for Scientific and Technological Development (CNPq, 305007/2018-1), the Universidade Tecnológica Federal do Paraná (UTFPR), and IOTO International Company.

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Rheological and thermal properties of EVA-organoclay systems using an environmentally friendly clay modifier the morphology and properties of EVA/clay nanocomposites. Polymer, 44(26), 7953-7961. http://dx.doi.org/10.1016/j. polymer.2003.10.046. 20. Ugel, E., Giuliano, G., & Modesti, M. (2011). Poly(ethylene– co–vinyl acetate)/clay nanocomposites: effect of clay nature and compatibilising agents on morphological thermal and mechanical properties. Soft Nanoscience Letters, 01(04), 105-119. http://dx.doi.org/10.4236/snl.2011.14018. 21. Lee, K. M., & Han, C. D. (2003). Rheology of organoclay nanocomposites: effects of polymer matrix/organoclay compatibility and the gallery distance of organoclay. Macromolecules, 36(19), 7165-7178. http://dx.doi.org/10.1021/ma030302w. 22. Cui, L., Khramov, D. M., Bielawski, C. W., Hunter, D. L., Yoon, P. J., & Paul, D. R. (2008). Effect of organoclay purity and degradation on nanocomposite performance, Part 1: surfactant degradation. Polymer, 49(17), 3751-3761. http:// dx.doi.org/10.1016/j.polymer.2008.06.029. 23. Wang, G., Wang, S., Sun, Z., Zheng, S., & Xi, Y. (2017). Structures of nonionic surfactant modified montmorillonites and their enhanced adsorption capacities towards a cationic organic dye. Applied Clay Science, 148, 1-10. http://dx.doi. org/10.1016/j.clay.2017.08.001. 24. Zhuang, G., Zhang, Z., & Jaber, M. (2019). Organoclays used as colloidal and rheological additives in oil-based drilling fluids: an overview. Applied Clay Science, 177, 63-81. http:// dx.doi.org/10.1016/j.clay.2019.05.006. 25. Morita, R. Y., Kloss, J. R., & Barbosa, R. V. (2014). Characterization of mechanical and thermal properties of poly(ethylene-co-vinyl acetate) with differents bentonites. Macromolecular Symposia, 343(1), 88-95. http://dx.doi.org/10.1002/masy.201300198. 26. Carli, L. N., Daitx, T. S., Guégan, R., Giovanela, M., Crespo, J. S., & Mauler, R. S. (2014). Biopolymer nanocomposites based on poly(hydroxybutyrate-co-hydroxyvalerate) reinforced by a non-ionic organoclay. Polymer International, 64(2), 235-241. http://dx.doi.org/10.1002/pi.4781. 27. Morita, R. Y., Kloss, J. R., & Barbosa, R. V. (2015). Caracterização de bentonitas sódicas: efeito o tratamento com surfactante orgânico livre de sal de amônio. Revista Virtual de Química, 7(4), 1286-1298. http://dx.doi.org/10.5935/1984-6835.20150071. 28. Iodice, B., Torrens, G. L., Kloss, J. R., Reis, D. M., & Jarek, F. (2010) BR nº PI 10013121. Processo de obtenção de nanoargila modificada para a produção de nanocompósitos poliméricos e nanoargila modificada. 29. Iodice, B., Morita, R. Y., Kloss, J. R., Torrens, G. L., & Barbosa, R. V. (2013) WO 2013/185196 A1. Use of an ammonium salt-free organophilic nanostructured clay in polyethylene. 30. Pistor, V., Lizot, A., Fiorio, R., & Zattera, A. J. (2010). Influence of physical interaction between organoclay and poly(ethylene-co-vinyl acetate) matrix and effect of clay content on rheological melt state. Polymer, 51(22), 5165-5171. http:// dx.doi.org/10.1016/j.polymer.2010.08.045. 31. Lagaly, G. (1986). Interaction of alkylamines with different types of layered compounds. Solid State Ionics, 22(1), 43-51. http://dx.doi.org/10.1016/0167-2738(86)90057-3. 32. Jiang, L., Zhang, J., & Wolcott, M. P. (2007). Comparison of polylactide/nano-sized calcium carbonate and polylactide/

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montmorillonite composites: reinforcing effects and toughening mechanisms. Polymer, 48(26), 7632-7644. http://dx.doi. org/10.1016/j.polymer.2007.11.001. 33. La Mantia, F. P., Scaffaro, R., Ceraulo, M., Mistretta, M. C., Dintcheva, N. T. Z., & Botta, L. (2016). A simple method to interpret the rheological behavior of intercalated polymer nanocomposites. Composites. Part B, Engineering, 98, 382388. http://dx.doi.org/10.1016/j.compositesb.2016.05.045. 34. Lotti, C., Isaac, C. S., Branciforti, M. C., Alves, R. M. V., Liberman, S., & Bretas, R. E. S. (2008). Rheological, mechanical and transport properties of blown films of high density polyethylene nanocomposites. European Polymer Journal, 44(5), 1346-1357. http://dx.doi.org/10.1016/j. eurpolymj.2008.02.014. 35. Son, D., Cho, S., Nam, J., Lee, H., & Kim, M. (2020). X-raybased spectroscopic techniques for characterization of polymer nanocomposite materials at a molecular level. Polymers, 12(5), 1053. http://dx.doi.org/10.3390/polym12051053. PMid:32375363. 36. Malkin, A. Y. (1994). Rheology fundamentals. Toronto: ChemTec Publishing. 37. Oliveira, A. G., Moreno, J. F., de Sousa, A. M. F., Escócio, V. A., Guimarães, M. J. O. C., & da Silva, A. L. N. (2020). Composites based on high-density polyethylene, polylactide and calcium carbonate: effect of calcium carbonate nanoparticles as co-compatibilizers. Polymer Bulletin, 77(6), 2889-2904. http://dx.doi.org/10.1007/s00289-019-02887-9. 38. Allen, N. S., Edge, M., Rodriguez, M., Liauw, C. M., & Fontan, E. (2000). Aspects of the thermal oxidation of ethylene vinyl acetate copolymer. Polymer Degradation & Stability, 68(3), 363-371. http://dx.doi.org/10.1016/S0141-3910(00)00020-3. 39. Rimez, B., Rahier, H., Van Assche, G., Artoos, T., Biesemans, M., & Van Mele, B. (2008). The thermal degradation of poly(vinyl acetate) and poly(ethylene-co-vinyl acetate), Part I: experimental study of the degradation mechanism. Polymer Degradation & Stability, 93(4), 800-810. http://dx.doi. org/10.1016/j.polymdegradstab.2008.01.010. 40. He, H., Ding, Z., Zhu, J., Yuan, P., Xi, Y., Yang, D., & Frost, R. L. (2005). Thermal characterization of surfactant-modified montmorillonites. Clays and Clay Minerals, 53(3), 287-293. http://dx.doi.org/10.1346/CCMN.2005.0530308. 41. Beltrán, M. I., Benavente, V., Marchante, V., Dema, H., & Marcilla, A. (2014). Characterization of montmorillonites simultaneously modified with an organic dye and an ammonium salt at different dye/salt ratios: properties of these modified montmorillonites EVA nanocomposites. Applied Clay Science, 97-98, 43-52. http://dx.doi.org/10.1016/j.clay.2014.06.001. Received: Mar. 23, 2020 Revised: Aug. 10, 2020 Accepted: Aug. 20, 2020

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ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.05320

Improved durability of Bisphenol A polycarbonate by bilayer ceramic nano-coatings alumina-zinc oxide Abdellah Moustaghfir1* , Agnes Rivaton2, Bénédicte Mailhot2 and Michel Jacquet2 Laboratory of a Research Odontologicals, Biomaterials, and Nanotechnology, Faculty of Dentistry, Mohammed V University in Rabat, Rabat, Morocco 2 Institut de Chimie de Clermont-Ferrand, Université Clermont Auvergne, Aubière, France

*abdellah.moustaghfir@um5.ac.ma

Abstract Polycarbonate exposed to sunlight yellows, degrades and loses its usable properties. In order to increase its lifetime, it can be coated with nano-ceramic thin layers of ZnO and Al2O3 deposited by sputtering. The role of the ZnO is to absorb the UV photons that can damageable for the polycarbonate. However, one of the limitations in the use of ZnO is the photocatalytic oxidation that could occur at interface ZnO/PC as a consequence of the photocatalytic activity of this oxide. Insertion of Al2O3 between PC and ZnO could be a way to inhibit this interfacial oxidation. The photooxidation of the ceramic/polymer assemblies, in condition of artificial accelerated ageing, was measured by infra-red and UV-vis spectroscopies. The results show that the photocatalytic activity of ZnO occurring in monolayer coated substrates can be significantly reduced by insertion of Al2O3 and that, in addition, Al2O3 decreases the permeability to oxygen of the coating. Keywords: photoprotection, photoageing, polycarbonate, thermooxidation, thin films. How to cite: Moustaghfir, A., Rivaton, A., Mailhot, B., & Jacquet, M. (2020). Improved durability of Bisphenol A polycarbonate by bilayer ceramic nano-coatings alumina-zinc oxide. Polímeros: Ciência e Tecnologia, 30(3), e2020029. https://doi.org/10.1590/0104-1428.05320

1. Introduction Bisphenol-A polycarbonate (PC) has an excellent toughness, a high transparency and is lightweight. As a consequence, PC is often used as a lighter and tougher substitute for glass or metals in a wide range of applications, including building and automobiles[1]. Exposing this polymer to weathering drastically changes their properties. As a consequence of the ability of the polymer to absorb UV radiations of the sunlight, several chemical reactions are induced causing embrittlement and colour changes[2,3]. The chemical mechanisms which are responsible for the UV light induced degradation of PC[4,5] are well known. When sunlight falls onto PC, the monomer units absorb energy of short wavelength in the near-UV range: this is particularly so for the ester and the carbonate groups in α position to the aromatic rings. The absorbed energy provokes the rupture of covalent bond triggering photolytic reaction (without oxygen intervention) and photooxydative reaction (with oxygen fixation) reactions. Both mechanisms can take place in the environment and are closely intertwined[6,7] (Figure 1). In order to suppress these damages, several methods can be used: - An effective way of protecting polymers from the effects of photodegradation is to add to the polymers UV-absorbing additives. These non-coloured UV-filters absorb the incident UV light that is damageable for the

Polímeros, 30(3), e2020029, 2020

polymer and are transparent in the visible domain[8]. The UV energy absorbed by the stabiliser must be dissipated without producing reactive species. However, the inhibition of the ageing by UV-absorbers is not fully efficient when photodegradation results from direct absorption of sunlight radiations by the intrinsic chromophoric groups of polymer[9]. This is the case for aromatic polymers. This effect is a direct result of the Beer-Lambert law, and a simple calculation shows that, in the first layers, the competition of absorption between the polymer and the additive is in favour of the polyme[10]. Consequently, an efficient photoprotection of the sample surface (below 100 µm) cannot be achieved. Another drawback of the use of UV-absorbers is their fatigue as these compounds are known to also act as anti-oxidant[11]: their degradation gives rise to the formation of coloured by-products; - Protecting the surface by coating the substrate with a thick varnish containing UV-absorbers is another solution to achieve photoprotection. This permits solving the problem of the competition of absorption between the polymer and the UV-absorber. However, the photooxidation of the varnish may occur inducing the fatigue of the UV-absorbers[12,13]. Consequently, the ability of the surface layer to totally absorb the incident near-UV light is reduced and a loss of adhesion of the coating occurs; a significant decrease of the durability of the substrate is then observed[14];

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

Moustaghfir, A., Rivaton, A., Mailhot, B., & Jacquet, M. The photochemical behaviour of the polymer and ceramic/polymer assemblies was tested by exposing the samples in an artificial accelerated photo-ageing device. The chemical evolution of the polymer was assessed by infra-red and UV-vis analysis.

2. Materials and Methods

Figure 1. Brief illustration of PC photodegradation mechanism.

- We have developed another method to stabilize polymer materials which consists in depositing onto the surface of the polymer a thin ceramic layer transparent in the visible range. The stabilizing action of the ceramic layer is associated with its ability to physically screen out the incident radiation, reducing therefore the rate of the photochemical processes. In addition, the coating acts as a barrier to oxygen and consequently the rate of photooxidation is decreased. This technique has been successfully employed for PET[15], PEEK[16], PEN[17] and PMMA[18]. The ceramic used was zinc oxide (ZnO) for good chemical stability and interesting optical properties because it’s transparent in the visible range and absorb UV radiation below 380 nm[19,20]. Coatings were elaborated by Radiofrequency Magnetron Sputtering, which permits working at low temperature (generally lower than 60 °C). Moreover, the deposits obtained with this technique have a higher density and a better adhesion to the substrate than those elaborated by other methods such as vacuum evaporation[21-23].

This method based on an inert outer filter constitutes an interesting alternative to the more classical methods. However, it is known that ZnO manifest a photocatalytic activity[24]. The Photoactivity due from the development of active species that can cause photooxidative degradation of the polymer films through reaction with oxygen and/ or water[25]. Most authors have proposed the HOO• and HO• radicals as the predominant reactive species[26,27]. The photoactivity of these pigments could then constitute a limitation to their use as coatings on polymeric substrates because they could induce a photocatalytic oxidation of the polymer at the interface. In this paper, we report the results of a preliminary investigation of the possibility of reducing the photocatalysed oxidation by insertion of an inactive Al2O3 layer (no photocatalytic activity at the opposite of ZnO) between the polymer and the upper ceramic coating. The feasibility of achieving bilayer ceramic coatings on the polymeric substrate was carried out on the polymer film of Bisphenol-A polycarbonate PC. PC/Al2O3/ZnO and PC/ZnO/Al2O3 assemblies ware successfully obtained. 2/7

PC film (50 μm thick) was supplied by Technifilm (ref. Makrofol D.E. 6-2). This polymer does not contain antioxidant. It was ultrasonically cleaned in ethanol before coating. The ceramic coatings of Al2O3 and ZnO were carried out in a sputtering unit (Alcatel SCM 450) equipped with a radiofrequency generator operating at 13.56 MHz. Bulk ZnO and Al2O3 targets (purity 99.9%; diameter 100 mm) fixed on cooled magnetron effect cathodes were used as starting materials. The substrates were situated at a distance of 90 mm from the targets. The sputtering chamber was evacuated below a pressure of 10-4 Pa before admitting the sputtering gas which was either pure argon or argon-oxygen mixtures. The plasma composition was controlled by a mass flow meter. The thickness of the deposits was determined by mean of a Jobin-Yvon ELLISEL single wavelength ellipsometre, using the 632.8 nm wavelength. PC films irradiations were carried out in SEPAP 12.24 units -Atlas- at a temperature of 60 °C with medium-pressure mercury lamps. This medium-acceleration photoaging device has been described previously[4]. It allows irradiation at wavelengths above 300 nm so that the ageing is representative of the behaviour under natural exposure. The ceramic coating was deposited on both sides of the films. Thermooxidation experiments were carried out in a ventilated oven at temperatures of 170 °C. The evolutions of UV-vis and infrared spectra were recorded respectively on a Shimadzu UV-2101 PC equipped with an integrating sphere and on a Nicolet Magna-IR 760 FTIR spectrophotometer.

3. Results and Discussions 3.1 Photooxidation of PC Since polycarbonate light absorption is up to 330 nm, this polymer is directly reachable to UV light that is present in terrestrial solar radiation. The resulting photodegradation of PC at λ > 300 nm leads to noticeable modifications of the UV-vis and infrared spectra of irradiated films that are worthy to be recalled. Figure 2 shows the effect of irradiation on the UV-vis absorption spectrum of PC. For short exposure time, the maxima observed to develop around 320 and 355 nm are related to photo-Fries rearrangement of the carbonate units[4]. As irradiation proceeds, these bands are rapidly overlapped by an unstructured absorption at wavelengths below 500 nm. This absorption, attributed to a mixture of colored species formed in ring oxidation[4-6], produces the yellowing of the irradiated film. Photooxidation of PC leads in parallel to notable modifications of the IR spectra of the samples. In the hydroxyl absorption region (3800-3000 cm-1), a broad Polímeros, 30(3), e2020029, 2020

Improved durability of Bisphenol A polycarbonate by bilayer ceramic nano-coatings alumina-zinc oxide increase of absorbance centred around 3400 cm-1 is observed (Figure 3) and attributed to the formation of alcohols, acids and hydroperoxides[4-6]. The rates of photooxidation and photolysis can be characterised by determining the concentration of the stable

photoproducts detected by UV-vis and infrared analysis, and that accumulates during irradiation: - The photo-yellowing of polycarbonate can be assessed by measuring the absorption increase at a wavelength of 400 nm; - Measuring the increase of absorbance in the hydroxyl region can be used to characterise accurately the progress of the photodegradation reactions in PC.

3.2 Nano-ceramic coating on PC

Figure 2. Evolution of the UV-visible spectrum of PC films subjected to irradiation.

Figure 3. Evolution of the IR spectrum of PC (a) in the 4000-400 cm-1 and (b) in the 3800-3000 cm-1 –hydroxyl region-, during irradiation. Polímeros, 30(3), e2020029, 2020

Spectra reported on Figure 4 show the absorption properties of ZnO and Al2O3 mono or bilayer coatings in the UV-vis region. ZnO presents a high absorption in the area from 300 to 400 nm. This absorption clearly depends on the thickness of the coated layer. On the opposite, Al2O3 is totally transparent in this region. The bilayer Al2O3/ZnO coating absorbs more than ZnO/Al2O3 in this region. In a former paper[19], we have shown that ZnO coating is efficient in protecting PC against photodegradation. The photostability of coated PC increases with the thickness of the layer. As an example, Figure 5 shows the increase of absorbance at 400 nm and in the hydroxyl domain versus the irradiation time of unprotected PC and of three different polymer/ceramic assemblies: 50 nm ZnO, 100 nm ZnO and 600 nm ZnO. One can note that the rate of photooxidation decreases when the thickness of the coating increases and that the photoprotective effect is fully efficient when the thickness is 600 nm. This photostabilising effect can be easily explained. An increase of the thickness of ZnO coating not only increases the screening effect of the ceramic for the photons damageable for the polymer, but also gives ZnO grain with higher size and density[19]. Therefore, the permeability to oxygen of the ceramic coating is decreased and consequently the rate of oxidation of the polymer is also reduced. Then the feasibility of Al2O3 deposit on PC was experimented. In a first time, Al2O3 coatings of 100 nm and 600 nm have been deposited on PC. The rate of yellowing and the rate of formation of hydroxyl products are shown in Figure 5. Analysis of the curves given in this figure provides the following comment: at low conversion degree (below 100 h), the photodegradation of PC is not inhibited by Al2O3 whatever

Figure 4. UV–visible spectra of reference and ceramic(s) coated PC films before irradiation versus different thicknesses. 3/7

Moustaghfir, A., Rivaton, A., Mailhot, B., & Jacquet, M. its thickness; but for longer exposure duration, Al2O3 causes a decrease in the rate of oxidation of PC. These results can be interpreted through simultaneous intervention of photooxidative and photolytic reactions[28]. The UV-visible spectrum in Figure 4 obviously indicates that the polycarbonate cannot be protected against solar radiation by alumina coatings. As a consequence, photo-Fries rearrangement of PC, which does not implicate oxygen, is not inhibited by the presence of the alumina thin film. However, the ulterior oxidative reactions induced by the photo-Fries products, are prevented because Al2O3 behaves and acts as an oxygen inhibitor. To put it differently, alumina does not inhibit photolytic processes (without oxygen intervention) but decreases the degree of oxidative reactions as this covering is impermeable to oxygen. In a second time, polycarbonate coated with bilayers ZnO and Al2O3 samples were experienced. To find out some degradation and yellowing after an adapted exposure time, we reduced the thickness of ZnO and Al2O3 coatings to 50 nm. To obtain a precise proof of ZnO photocatalytic activity action at the polycarbonate interface, another test was

carried out in which Al2O3 was deposited above ZnO. The rate of photo-yellowing and the rate of formation of hydroxyl products are shown in Figure 5. Analysis of the curves presented in this figure indicates that Al2O3/ZnO coatings have a greater photostabilsation efficacy than ZnO alone. This effect can be explained by the Al2O3 coating which eliminates the photocatalytic activity of ZnO, which in turn promotes PC degradation at the interface between PC and ZnO. Whereas the photooxidation and photo-yellowing of bilayer PC/ZnO/Al2O3 coating is well noted and thus less photostabilizing and so less photoprotective than PC/Al2O3/ZnO coatings. The scanning electron microscopy was used to evaluate the superficial changes of PC coated with ZnO (50 nm and 600 nm) and those coated with bilayer ZnO (50 nm) and Al2O3 (50 nm). Figure 6 shows the micrographs of the layer surface obtained. The sputtered thin films have a granular microstructure. The grain size is about 40 to 130 nm diameter. The growth of the deposit is columnar type. The size of the columns, whose ends appear at the surface of the deposit, increases with the thickness. The ceramic coating PC/Al2O3-ZnO exhibits an increase in the size of columns which corresponds to a coalescence of grains forming a homogeneous surface. Therefore, PC / Al2O3-ZnO nanocoatings are dense and exhibit the best barrier properties against the diffusion of gases, especially oxygen. In conclusion, PC/Al2O3/ZnO coatings exhibit the best photoprotective efficiency as this assemblage integrates ZnO capacity to absorb in the UV-vis band with Al2O3 oxygen barrier property without the photocatalytic effect. Table 1 summarizes the photoprotection of polycarbonate by zinc oxide and/or alumina ceramic nano-coatings.

3.3 Thermooxidative ageing

Figure 5. Increase of the absorbance at (a) 400 nm and (b) 3470 cm-1 versus irradiation time.

In order to determine whether ceramic deposits could inhibit PC thermooxidation because of their barrier effect to oxygen diffusion. PC uncoated and PC covered with ceramic coatings were placed at a temperature of 170 °C in a ventilated oven for a period of approximately 400 days. Thermooxidation kinetics of the samples was followed to observe the appearance and development of thermooxidized photoproducts. Figure 7 shows the evolution of the UV-visible absorption spectra of virgin PC and PC coated with bilayer ceramic Al2O3/ZnO during the thermooxidation. The absorbance increases without a specific maximum, in contrast to photooxidation where the formation of photo-Fries products resulted in the appearance of an absorption band at 320 nm. The yellowing that develops in thermooxidation is approximately equivalent for coated PC films as for virgin PC. Figures 8 and 9 represent the IR spectra of these thermooxidized samples in the domain of carbonylated and hydroxylated products. The bands at 1724 and 1690 cm-1 are

Table 1. Photoprotective ceramic coatings on PC. Sample

4/7

Protection against:

Interfacial photocatalytic activity

PC/ZnO

Photon yes

Oxygen yes

PC/Al2O3

yes

PC/Al2O3-ZnO

yes

Polímeros, 30(3), e2020029, 2020

Improved durability of Bisphenol A polycarbonate by bilayer ceramic nano-coatings alumina-zinc oxide

Figure 6. Scanning Electron Micrographs of polycarbonate film surfaces coated by (a) ZnO-50 nm, (b) ZnO-600 nm (c) ZnO-50 nm/Al2O3-50 nm and (d) Al2O3-50 nm/ZnO-50 nm.

Figure 7. Evolution of the UV-visible spectra of (a) virgin PC and (b) PC coated with Al2O3/ZnO bilayer during thermooxidation at 170 °C; (c) Change in absorbance determined from UV-vis (400 nm) spectra vs exposure time. Polímeros, 30(3), e2020029, 2020

Figure 8. Evolution of the spectra in the region of carbonylated product of (a) virgin PC and (b) PC coated with Al2O3/ZnO bilayer during thermooxidation at 170 °C. 5/7

Moustaghfir, A., Rivaton, A., Mailhot, B., & Jacquet, M. attributed respectively to aliphatic and aromatic ketones[2]; as for the doublet at 1840/1860 cm-1, also observed in photooxidation (Figure 3a), indicates the oxidation of the aromatic ring[2]. One of the differences between thermo- and photo-oxidation is the non-accumulation in thermooxidized films of saturated aliphatic acids detected at 1713 cm-1 in photooxidation. In the hydroxyl domain, the formation of hydroxylated products (alcohols and phenols) characterized by absorption bands at 3560 and 3514 cm-1 is observed. On the other hand, there is no appearance of a wide absorption band between 3200-3300 cm-1 corresponding to carboxylic acids. This observation is in good agreement with the analysis of the carbonyl domain. The advancement state of polycarbonate thermooxidation included in the different systems can be characterized by the formation of the hydroxylated and carbonyl products. Figure 10 shows that the rate of thermooxidation is slowed down by the presence of a ceramic coating. The most effective coating is the one with bilayer Al2O3/ZnO: the role of the ceramic is to limit the diffusion of oxygen responsible for the thermal ageing of the PC, and the barrier effect increases with the thickness of the surface deposit.

4. Conclusions

Figure 9. Evolution of the spectra in the region of hydroxylated products of (a) virgin PC and (b) PC coated with Al2O3/ZnO bilayer during thermooxidation at 170 °C.

Coating polycarbonate with ZnO thin layers reduces the rate of photodegradation of the polymer as a result of the screen effect role of the ceramic. Meanwhile, our results show that there coatings have some photocatalytic activity. The light excitation of the thin ZnO layers generates the formation of activated species (mainly OH•, HO2• radicals) which are susceptible to initiate photooxidative reactions at the surface of the coated polymer. The insertion of an Al2O3 thin layer between the ceramic and the polymer provides a higher photoprotective efficiency which can be attributed to the suppression of the photocatalysed degradation of the polymer at the interface PC/ZnO. In addition, the Al2O3 deposit acts as a barrier to oxygen limiting therefore oxidative degradation involved in the mechanism of photodegradation of the coated polymer. The results show that it is also possible to deposit Al2O3 above ZnO obtaining therefore a hard upper layer. In term of photostabilisation, this bilayer is also more efficient than ZnO alone as both the screening effect of ZnO and the impermeability of Al2O3 are involved. As a conclusion, the durability of PC can be drastically improved by a bilayer Al2O3-ZnO ceramic nano-coatings which ensures a good photoprotection and significant thermooxidative inhibition.

5. References

Figure 10. Kinetics of appearance of thermooxidation products in (a) hydroxyl region (b) carbonyl region. 6/7

1. Alavi Nikje, M. M., & Askarzadeh, M. (2013). Green and inexpensive method to recover Bisphenol-A from polycarbonate wastes. Polímeros: Ciência e Tecnologia, 23(1), 29-31. http:// dx.doi.org/10.1590/S0104-14282013005000019. 2. Wu, D., Zhang, D., Liu, S., Jin, Z., Chowwanonthapunya, T., Gao, J., & Li, X. (2020). Prediction of polycarbonate degradation in natural atmospheric environment of China based on BPANN model with screened environmental factors. Chemical Polímeros, 30(3), e2020029, 2020

Improved durability of Bisphenol A polycarbonate by bilayer ceramic nano-coatings alumina-zinc oxide Engineering Journal, 399, 125878. http://dx.doi.org/10.1016/j. cej.2020.125878. 3. Motta, A., La Mantia, F. P., Ascione, L., & Mistretta, M. C. (2020). Theoretical study on the decomposition mechanism of bisphenol A polycarbonate induced by the combined effect of humidity and UV irradiation. Journal of Molecular Graphics & Modelling, 99, 107622. http://dx.doi.org/10.1016/j. jmgm.2020.107622. PMid:32344302. 4. Rivaton, A. (1995). Recent advances in bisphenol-A polycarbonate photodegradation. Polymer Degradation & Stability, 49(1), 163-179. http://dx.doi.org/10.1016/0141-3910(95)00069-X. 5. Rivaton, A., Mailhot, B., Soulestin, J., Varghese, H., & Gardette, J.-L. (2002). Comparison of the photochemical and thermal degradation of bisphenol-A polycarbonate and trimethylcyclohexanepolycarbonate. Polymer Degradation & Stability, 75(1), 17-33. http://dx.doi.org/10.1016/S0141-3910(01)00201-4. 6. Rivaton, A., Mailhot, B., Soulestin, J., Varghese, H., & Gardette, J.-L. (2002). Influence of the chemical structure of polycarbonates on the contribution of crosslinking and chain scissions to the photothermal ageing. European Polymer Journal, 38(7), 1349-1369. http://dx.doi.org/10.1016/S00143057(01)00307-X. 7. Pickett, J. A. (2011). Influence of photo-Fries reaction products on then photodegradation of bisphenol-A polycarbonate. Polymer Degradation & Stability, 96(12), 2253-2265. http:// dx.doi.org/10.1016/j.polymdegradstab.2011.08.016. 8. Mohammed, A., El-Hiti, G., Yousif, E., Ahmed, A. A., Ahmed, D. S., & Alotaibi, M. H. (2020). Protection of poly(vinyl chloride) films against photodegradation using various valsartan tin complexes. Polymers, 12(4), 969. http://dx.doi.org/10.3390/ polym12040969. PMid:32326307. 9. Lungulescu, E. M., & Zaharescu, T. (2016). Stabilization of polymers against photodegradation. In D. Rosu & P. M. Visakh (Eds.), Photochemical behavior of multicomponent polymeric-based materials (Advanced Structured Materials, Vol. 26, pp. 165-192). Cham: Springer. http://dx.doi.org/10.1007/978-3-319-25196-7_6. 10. Diepens, M. (2009). Photodegradation and stability of bisphenol a polycarbonate in weathering conditions. Eindhoven: Technische Universiteit Eindhoven. https://doi.org/10.6100/IR642300. 11. Allen, N. S., Luc-Gardette, J., & Lemaire, J. (1983). Photostabilising action of ortho-hydroxy benzophenones in polypropylene film: influence of processing and wavelength of irradiation. Polymer Photochemistry, 3(4), 251-265. http:// dx.doi.org/10.1016/0144-2880(83)90034-9. 12. Claudé, B., Gonon, L., Verney, V., & Gardette, J.-L. (2001). Consequences of photoageing on the durability of plastic glasses for automotive applications. Polymer Testing, 20(7), 771-778. http://dx.doi.org/10.1016/S0142-9418(01)00022-8. 13. Alsadi, J. (2020). Systematic review: impact of processing parameters on dispersion of polycarbonate: composites, and pigment characterized by different techniques. Materials Today: Proceedings, 27(4), 3254-3264. http://dx.doi.org/10.1016/j. matpr.2020.05.027. 14. Saron, C., Felisberti, M. I., Zulli, F., & Giordano, M. (2007). Effects of bismuth vandate and anthraquinone dye on the photodegradation of polycarbonate. Journal of the Brazilian Chemical Society, 18(5), 900-910. http://dx.doi.org/10.1590/ S0103-50532007000500005. 15. Awitor, K. O., Rivaton, A., Gardette, J.-L., Down, A. J., & Johnson, M. B. (2007). Photo-protection and photo-catalytic activity of crystalline anatase titanium dioxide sputter-coated on polymer films. Thin Solid Films, 516(8), 2286-2291. http:// dx.doi.org/10.1016/j.tsf.2007.08.005. 16. Giancaterina, S., Ben Amor, S., Baud, G., Gardette, J.-L., Jacquet, M., Perrin, C., & Rivaton, A. (2002). Photoprotective Polímeros, 30(3), e2020029, 2020

ceramic coatings on poly(ether ether ketone). Polymer, 43(24), 6397-6405. http://dx.doi.org/10.1016/S0032-3861(02)00499-8. 17. Guedri-Knani, L., Gardette, J. L., Jacquet, M., & Rivaton, A. (2004). Photoprotection of poly(ethylene-naphthalate) by zinc oxide coating. Surface and Coatings Technology, 180-181(7175), 71-75. http://dx.doi.org/10.1016/j.surfcoat.2003.10.039. 18. Chodun, R., Skowronski, S., Okrasa, B., Wicher, K., Nowakowska-Langier, K., & Zdunek, K. (2019). Optical TiO2 layers deposited on polymer substrates by the Gas Injection Magnetron Sputtering technique. Applied Surface Science, 466, 12-18. http://dx.doi.org/10.1016/j.apsusc.2018.10.003. 19. Moustaghfir, A., Tomasella, E., Rivaton, A., Mailhot, B., Jacquet, M., Gardette, J.-L., & Cellier, J. (2004). Sputtered zinc oxide coatings: structural study and application to the photoprotection of the polycarbonate. Surface and Coatings Technology, 180-181, 642-645. http://dx.doi.org/10.1016/j. surfcoat.2003.10.109. 20. Ghamsari, M. S., Alamdari, S., Han, W., & Park, H. H. (2016). Impact of nanostructured thin ZnO film in ultraviolet protection. International Journal of Nanomedicine, 12, 207-216. http:// dx.doi.org/10.2147/IJN.S118637. PMid:28096668. 21. Mosbah, A., Moustaghfir, A., Abed, S., Bouhssira, N., Aida, M. S., Tomasella, E., & Jacquet, M. (2005). Comparison of the structural and optical properties of zinc oxide thin films deposited by d.c. and r.f. sputtering and spray pyrolysis. Surface and Coatings Technology, 200(1-4), 293-296. http:// dx.doi.org/10.1016/j.surfcoat.2005.02.012. 22. Juarez, T., Schroer, A., Schwaiger, R., & Hodge, A. M. (2018). Evaluating sputter deposited metal coatings on 3D printed polymer micro-truss structures. Materials & Design, 140, 442-450. http://dx.doi.org/10.1016/j.matdes.2017.12.005. 23. Andrade, J. E., Machado, R., Macêdo, M. A., & Cunha, F. G. C. (2013). AFM and XRD characterization of silver nanoparticles films deposited on the surface of DGEBA epoxy resin by ion sputtering. Polímeros: Ciência e Tecnologia, 23(1), 19-23. http://dx.doi.org/10.1590/S0104-14282013005000009. 24. Zailan, S. N., Bouaissi, A., Mahmed, N., & Abdullah, M. M. A. (2020). Influence of ZnO nanoparticles on mechanical properties and photocatalytic activity of self-cleaning ZnO-based geopolymer paste. Journal of Inorganic and Organometallic Polymers and Materials, 30(6), 2007-2016. http://dx.doi. org/10.1007/s10904-019-01399-3. 25. Kamalian, P., Khorasani, S. N., Abdolmaleki, A., Karevan, M., Khalili, S., Shirani, M., & Neisiany, R. E. (2020). Toward the development of polyethylene photocatalytic degradation. Journal of Polymer Engineering, 40(2), 181-191. http://dx.doi. org/10.1515/polyeng-2019-0230. 26. Lemaire, J. (1982). The photocatalyzed oxidation of polyamides and polyolefins. Pure and Applied Chemistry, 54(9), 1667-1682. http://dx.doi.org/10.1351/pac198254091667. 27. Serpone, N. (2000). Photocatalysis. In R. E. Kirk & D. F. Othmer (Eds.), Kirk-Othmer encyclopedia of chemical technology (Vol. 19, pp. 1-17). New York: John Wiley & Sons. https://doi.org /10.1002/0471238961.1608152019051816.a01. 28. Rivaton, A., Gardette, J.-L., Morlat-Therias, S., Mailhot, B., Tomasella, E., Awitor, O., Komvopoulos, K., & Fabbri, P. (2009). Enhancement of photoprotection and mechanical properties of polymers by deposition of thin coatings. In J. W. Martin, R. A. Ryntz, J. Chin & R. A. Dickie (Eds.), Service life prediction of polymeric materials (pp. 327-343). Boston: Springer. http://dx.doi.org/10.1007/978-0-387-84876-1_22. Received: May 08, 2020 Revised: Aug. 14, 2020 Accepted: Aug. 27, 2020 7/7

ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.06820

Layered cryogels laden with Brazilian honey intended for wound care Gabriela de Souza dos Santos1 , Natália Rodrigues Rojas dos Santos1 , Ingrid Cristina Soares Pereira1 , Antonio José de Andrade Júnior1 , Edla Maria Bezerra Lima2 , Adriana Paula Minguita2 , Luiz Henrique Guerreiro Rosado3 , Ana Paula Duarte Moreira4 , Antonieta Middea5 , Edlene Ribeiro Prudencio6 , Rosa Helena Luchese6  and Renata Nunes Oliveira3*  Departamento de Engenharia Química, Instituto de Tecnologia, Universidade Federal Rural do Rio de Janeiro – UFRRJ, Seropédica, RJ, Brasil 2 Empresa Brasileira de Pesquisa Agropecuária – Embrapa Tecnologia de Alimentos, Rio de Janeiro, RJ, Brasil 3 Programa de Pós-graduação em Engenharia Química, Instituto de Tecnologia, Universidade Federal Rural do Rio de Janeiro – UFRRJ, Seropédica, RJ, Brasil 4 Programa de Engenharia de Materiais e Metalurgia, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brasil 5 Centro de Tecnologia Mineral, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brasil 6 Departamento de Engenharia de Alimentos, Instituto de Tecnologia, Universidade Federal Rural do Rio de Janeiro – UFRRJ, Seropédica, RJ, Brasil 1

*renatanunes.ufrrj@gmail.com

Abstract PVA cryogels are well established as candidate biomaterials for wound healing applications but are not themselves biodegradable or antimicrobial. Blending PVA with NaCMC (CMC) or gelatin (G) can increase the gel’s ability to swell and would introduce a degree of biodegradability. The incorporation of appropriate amounts of a natural antimicrobial/healing agents, such as honey (H), would contribute to the gels properties. The present work addresses the development and characterization of layered gels (PVA-H, PVA-CMC-H and PVA-G-H, with empty PVA, PVACMC, PVA-G gels presented as controls). The gels were characterized by FTIR, DSC, in vitro analysis of swelling and microbiological (S. aureus) effects. Addition of gelatin, NaCMC and honey to PVA diminished the PVA chains’ ability to pack into crystallites. Samples containing honey swelled less and presented higher weight loss/biodegradability than samples without honey. Only the honey-laden PVA-CMC and PVA-G presented activity against S. aureus. Keywords: layered hydrogel, PVA, NaCMC, gelatin, honey. How to cite: Santos, G. S., Santos, N. R. R., Pereira, I. C. S., Andrade Júnior, A. J., Lima, E. M. B., Minguita, A. P., Rosado, L. H. G., Moreira, A. P. D., Middea, A., Prudencio, E. R., Luchese, R. H., & Oliveira, R. N. (2020). Layered cryogels laden with Brazilian honey intended for wound care. Polímeros: Ciência e Tenologia, 30(3), e2020031. https:// doi.org/10.1590/0104-1428.06820

1. Introduction Infection associated with wounds to the skin affect 14 million people per year in the USA. Such infections are characterized by colonization by gram-positive bacteria such as S. aureus in the early stages of healing, which are later replaced by gram-negative organisms[1]. Requirements for successful healing include a mechanically stable and moist environment, a capacity for absorbing wound exudate, and antimicrobial properties which act against the development of infection. Hydrogels have been manufactured in layers to mimic the layers and function of the skin. For example, chitosang-poly(ethylene glycol) hydrogel reinforced with chitosanalginate were designed to mimic the micro-environment relevant to skin tissue engineering[2]. A hydrogel composed

Polímeros, 30(3), e2020031, 2020

of layers of alginate/chitosan/poly(-glutamic acid) increased wound epithelialization and collagen regeneration[3]. Layered hydrogels based on PVA (polyvinyl alcohol) have also been developed: 2-layered skin equivalent PVA or PVAcellulose fibre blocks were prepared by freeze-thawing and presented roughness and elasticity similar to human skin[4]. A PAA layer has been added to a PVA layer, with the PAA solution poured on the top of PVA swollen layer, where the chains of both polymers would entangle and form hydrogen bonding at the interface. The presence of PAA led to increased swelling, adhesion and biocompatibility while the presence of PVA underpinned the mechanical properties of the gel[5]. Hydrogels are potential materials for burn dressings.

1/10

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

Santos, G. S., Santos, N. R. R., Pereira, I. C. S., Andrade Júnior, A. J., Lima, E. M. B., Minguita, A. P., Rosado, L. H. G., Moreira, A. P. D., Middea, A., Prudencio, E. R., Luchese, R. H., & Oliveira, R. N. PVA is a hydrophilic biocompatible polymer used to manufacture hydrogels. PVA hydrogels are capable of presenting physical properties similar to human tissue, such as elasticity[6]. Hydrogels can be based on chemical crosslinking or physical crosslinking, where chemical crosslinking is based on the use of crosslinking agents, such as glutaraldehyde, or based on the use of radiation[7,8]. Physically crosslinked PVA hydrogels can be prepared by freezing and then thawing an aqueous solution of PVA. PVA chains entangle when in solution. When frozen, ice crystals are formed. By phase separation (ice/PVA), PVA chains are pushed together forming crystallites, which are responsible for the high mechanical properties of these gels. These gels have tissue-like elasticity, toughness and are non-toxic[9]. Blended PVA gels incorporating several additives, such as L. bulgaricus extract[10] and neomycin sulfate[11], have been studied in the literature with respect to antibacterial and healing properties. Sodium carboxymethyl cellulose (NaCMC) is a hydrophilic polymer, a polyelectrolyte, derived from cellulose. The charges of this polyelectrolyte elongate the polymer chains, increasing the water uptake capability of the network, with counterions favouring water entrance[12]. NaCMC is able to form hydrogels by blending it with hydroxyethyl cellulose and chemically crosslinking with divynilsulphone[12]; NaCMC can be combined with protein sericin, crosslinked by freeze-thawing, with glutaraldehyde and AlCl3 to produce hydrogels for wound dressing[13]. NaCMC can also be blended with polyethylene glycol and crosslinked with citric acid for dressing purposes, showing no cytotoxicity[14]. In order to present antibacterial properties, NaCMC hydrogels can be loaded with drugs. For example, NaCMC gels can be crosslinked with citric acid and loaded with MCM-41 mesoporous silica nanoparticles containing tetracycline (antibiotic), where higher particle content led to higher S. aureus inhibition[15]. In another study, NaCMC was blended to PVA to physically (freeze-thawing) form hydrogels loaded with sodium fucidate, where the addition of NaCMC increased the gels swelling capacity, vapor transmission rate and porosity[16]. PVA-NaCMC-sodium fucidate hydrogel showed faster healing compared to a PVA-NaCMC gel[17]. Gelatin is a water-soluble protein obtained from the thermal denaturation of collagen. Gelatins with high levels of amino acids present high gel strength and melting point[18]. Gelatin is biocompatible, biodegradable, it has the ability to form films, and is generally low cost. In order to avoid dissolution in aqueous fluids, gelatin hydrogels have to be crosslinked with glutaraldehyde, genipin or lactose. Gelatin hydrogels crosslinked with lactose have been investigated as wound dressing materials[19]. Chitosan/gelatin crosslinked hydrogels were loaded with phenols from Hamamelis virginiana and showed activity against P. aeruginosa and S. aureus[20]. Superabsorbent gels were prepared by esterification of PVA with gelatin[21]. Chitosan/gelatin/PVA irradiated hydrogels for wound dressing applications presented swelling capacity, adequate vapor transmission rate and mechanical properties[22]. PVA/chitosan/gelatin hydrogels loaded with PCL microspheres containing bFGF and BSA (protein) showed no cytotoxicity, and an improved healing rate, while chitosan improved the antimicrobial properties[23]. 2/10

Honey is a bee-based product which can present more than 200 constituent components, mainly carbohydrates (most of them monosaccharides, e.g. fructose, glucose, sucrose), proteins, organic acids, vitamins and phenolic compounds. Honey’s antimicrobial characteristics are related to the H2O2 produced by enzymatic activity and to the action of complex phenols and organic acids (flavonoids)[24,25]. The honey composition is influenced by the flower types visited by bees, the soil composition, bee species and climatic conditions. Brazilian honey presents a considerable variability, since the local flora and climate change all over the country. The honey’s consistency, color, odor, taste and scent are constant physical-chemical characteristics in monoflower honeys. Honey’s phenolic substances and flavonoids are responsible by pharmacological properties, e.g. antimicrobian and antioxidant activities[26]. Wounded rats treated with Brazilian honey (from southwest region of Brazil) presented fastest recovery than the control group, it also treated the wounds infection[27]. Brazilian honey, also from the southwest, was tested against S. aureus microorganisms and it showed antimicrobial activity, due to its phenolic compounds[28]. Brazilian honey shows potential to be used as wound healing agent. Honey has been successfully incorporated in hydrogels. For example, a chitosan dressing containing 75% honey enhanced tissue regeneration[29]. A pectin-honey hydrogel promoted faster wound healing than a control[30]. PVA crosslinked with borax and loaded with honey promoted the proliferation of cells, with swelling, and permeability properties considered adequate for moderate exudative wounds[31]. Gelatin-chitosan hydrogels loaded with compounds from manuka honey presented antibacterial activity, a high wound healing rate and the absorption of exudates[32]. Gelatin-chitosan-honey hydrogel was effective against S. aureus and E. coli and stimulated burns healing[33]. PVP-agar-Peg gels containing 6% of honey acceletrated rats wounds contraction[34]; PVA/ chitosan/montmorillonite gels loaded with 15% Iranian honey showed potential to absorb exudate and efficient wound healing[35]; acrylamide hydrogels were loaded with different amounts of Indian honey (5-15%), but the gel containing 10% honey was considered the best one, due to optimized tensile strength and swelling capacity[36]; carbopol 934 and chitosan gels containing 75% Egyptian honey presented high burning healing rate[32]. The goal of the present work was to develop layered PVA, PVA-honey (PVA-H), PVA-NaCMC (PVA-CMC), PVA-NaCMC-honey (PVA-CMC-H), PVA-gelatin (PVA-G), PVA-gelatin-honey (PVA-G-H) hydrogels and characterise their physico-chemical and functional properties in order to gain insights into their potential development as woundcare biomaterials.

2. Materials and Methods 2.1 Sample preparation The samples were manufactured by the dissolution of PVA (Sigma-Aldrich, Mw: 85,000-124,000 Da, 99+% hydrolyzed) at 90 °C under mechanical stirring for 4h according to Table 1. For each PVA blend without honey, the amount of water used was split in half to separately dissolve Polímeros, 30(3), e2020031, 2020

Layered cryogels laden with Brazilian honey intended for wound care Table 1. Samples composition. Sample PVA (3 layers) PVA-CMC (3 layers) PVA-G (3 layers) PVA-H – 1st layer PVA-H – 2nd layer PVA-H – 3rd layer PVA-CMC-H – 1st layer PVA-CMC-H – 2nd layer PVA-CMC-H – 3rd layer PVA-G-H – 1st layer PVA-G-H – 2nd layer PVA-G-H – 3rd layer

PVA 10g 8g 8g 10g 10g 10g 8g 8g 8g 8g 8g 8g

NaCMC 2g 2g 2g 2g -

the PVA and either NaCMC (Sigma-Aldrich, average Mw ~250,000, degree of substitution 0.9) or Gelatin (from bovine skin, Sigma-Aldrich, Type B). The resulting solutions were then mixed together after each solution had reached room temperature (under stirring). The NaCMC was dissolved in water under stirring at room temperature for 2h. Gelatin was dissolved at 60 °C under stirring for 4h. The samples were prepared in three steps using 24-well plates. The composition of each layer of the samples without honey was similar (PVA or PVA-CMC or PVA-G). The first layer was obtained by pouring 1 mL of the polymer solution or the blend solution per well followed by a freeze-thaw cycle (each freeze-thawing cycle was to submit the sample to 1h at -18 °C and 30 min at room temperature). The second layer (1 mL/well) was subsequently added and freezethawed, followed by the addition and freeze-thawing of the third layer. The first layer was therefore subjected to 3 freeze-thawing cycles, the second layer to 2 cycles and the third layer to 1 cycle. For the samples that were to contain honey (a commercial Brazilian honey from Southeast region, characterized as “silvestre” honey), the polymers were dissolved in the designated amounts of water from Table 1 following the procedures previously described. Honey was then added to the polymer solution or to the blend solution, according to the quantities indicated in Table 1, at room temperature under mechanical stirring for 5 minutes. The first layer of these samples was either pure PVA or a PVA based polymer blend without honey and had experienced 3 freeze-thawing cycles. The second layer was the polymer/blend with 10% honey, submitted to 2 freeze-thawing cycles; and the third layer was polymer/blend with 5% honey, submitted to 1 freeze-thawing cycle. All the samples were dried in oven at 50 °C for 24h. Regarding honeys’ thermal degradation, honey samples treated at 23 °C for short times did not degrade, but the same behavior was not observed for samples heated at 95 °C, although antiocidant activity might increase with the heating temperature[34]. Heat treated honey might present hydroxymethylfurfural, which should be lower than 80 mg/kg in tropical honeys, since it is toxic and carcinogenic. In addition, tropical honeys can present fermentation when heat treated, but it can be prevented by heating honey at 60-70 °C for 10 min or at 60-65 °C for 30 min[35]. The drying temperature was low to preserve honey activities, although honey samples autoclaved at 121 °C could maintain its properties[27]. Polímeros, 30(3), e2020031, 2020

gelatin 2g 2g 2g 2g

honey 10 mL 5 mL 10 mL 5 mL 10 mL 5 mL

H2O 100 mL 100 mL 100 mL 100 mL 90 mL 95 mL 100 mL 90 mL 95 mL 100 mL 90 mL 95 mL

2.2 Physico-chemical analysis The physic-chemical analysis of the samples was performed by Fourier Transformed Infrared Spectroscopy (FTIR), equipment VERTEX-70 (UFRRJ), 32 scans per sample, from 4000 cm-1 to 600 cm-1. Samples representing each layer were prepared separately and submitted to the proper number of freeze-thawing cycles separately and then dried to be evaluated by FTIR. The multi-layered samples were also evaluated by FTIR.

2.3 Thermal analysis The samples were analysed by Differential Scanning Calorimetry (DSC), where approximately 5 mg of each sample composition was submitted to 10 °C/min of heat flow from room temperature to 240 °C using a DSC Q200 TA Instruments equipment (EMBRAPA Agroindustry). The transition temperatures of the samples (glass transition temperature - Tg, melting temperature - Tm) were obtained using the second cycle of heating to overcome the samples’ thermal history. The samples degree of crystallinity (Xc) was calculated by X c = 100

∆H m

∆H m100

(1)

The ΔHm is the enthalpy of melting related to the samples at the PVA melting temperature and ΔHm100 is the enthalpy of melting related to PVA 100% crystalline, 138.6 J/g[28].

2.4 In vitro analysis The swelling tests proceeded in saline solution at room temperature. For each sample composition, 5 samples were cut (samples weight was standardized among the ones with the same composition) and immersed in saline solution (20 mL per sample, used to mimic the body fluids). The samples were removed from the saline solution at regular intervals (30 min, 1h, 2h, 3h, 4h, 24h, 48h, 72h and 96h). The adsorbed fluid was removed using filter paper and the samples were weighed and then returned to the media. The swelling degree (SD) was calculated according to SD = 100

(WS − WDB )

WDB

(2)

3/10

Santos, G. S., Santos, N. R. R., Pereira, I. C. S., Andrade Júnior, A. J., Lima, E. M. B., Minguita, A. P., Rosado, L. H. G., Moreira, A. P. D., Middea, A., Prudencio, E. R., Luchese, R. H., & Oliveira, R. N. WDB is the weight of dried samples before swelling; WS is the weight of the swelled samples. At the end of the 4 days of immersion, the samples were dried in oven at 50 °C for 24h and weighed to calculate the sample gel fraction (GF) GF = 100

WDB

(3)

WD is the weight of the dried samples after swelling) and weight loss (WL)/biodegradability[34], WL = 100

(WDB − WD )

WDB

(4)

For microbiological analysis, the ASTM E2180- 07 standard was adapted. A suspension of Staphylococcus aureus cells (ATCC 6538) was prepared (guaranteeing that the suspension reached 0.5 in the MacFarland scale), which correspond to 108 colony forming unit (UFC)/mL. The agar paste received 1 mL of the mentioned suspension, where the agar paste total amount of microrganisms were 106 UFC/mL. Each sample was placed in an empty well of 24 wells flat bottom polystyrene plate, and each well received 200 μL of the agar paste containing S. aureus (where duplicates for each samples composition were evaluated). The plates were incubated at 30 °C for 24h. After incubation, the media in contact with samples were placed in Falcon tubes. After that, different amounts of buffer solution were added to prepare decimal dilutions of the media in contact with samples and to subsequently, count the colony forming units (cfu) by the micro dropping technique. The bacterial calculation was performed using the optical microscope (equipment Olympus). The reduction calculation was based on ASTM Standard 20170504.

3. Results and Discussions It was not possible to distinguish the samples layers, indicating that the system might behave homogeneously. All samples, containing honey or not, presented homogeneous morphology, similar to PVA-H, Figure 1.

Figure 1. PVA-H sample section. 4/10

3.1 Physico-chemical analysis The FTIR spectra of all samples are shown in Figure 2. Each layer of PVA presented similar FTIR spectra of the three-layered sample. Nonetheless, the band at 1142 cm-1, related to the formation of PVA crystallites (intra- and intermolecular hydrogen bonding between the chains originated by hydrophilic forces)[7], presented lower transmittance with increasing numbers of freeze-thawing cycles, indicating the contribution of the freeze-thawing process to the formation of PVA crystals. There were (in all PVA layers) bands at: 3273 cm-1 (OH hydrogen bonds)[7]; 2941 cm-1 and 2909 cm-1, ν(C–H) alkyl[7]; 1652 cm-1, ν(C=O-, from residual aldehyde)[35,36]; 1563 cm-1 and 1237 cm-1, ν(C=C-)[35]; 1411 cm-1, δ(-CH)[37]; 1380 cm-1, ω(C-H)[38]; 1329 cm-1, δ(CH+OH)[35]; 1089 cm-1, ν(C-O)[39]; 916 cm-1, δ(-CH2)[39]; and 835 cm-1, ρ(-CH)[39].

The previously reported PVA bands were also observed in the PVA-CMC as well as the overlapping bands of NaCMC, showing the dispersion of NaCMC in the PVA matrix[40]; a shoulder at 2855 cm-1, ν(C-H)[41], 1585 cm-1 (non-hydrated C=O of COO- group)[42]; 1415 cm-1, ν(COO-); 1325 cm-1, δ(C-H) of methyl groups; 1080 cm-1 ν(C-O)[43]. The band at 1652 cm-1 (C=O) is absent in the sample PVA-CMC[35]. The band at 1560 cm-1 (carbonyl group) in the PVA-CMC sample has a different shape compared to the same band in the PVA sample, indicating a change in the balance of free associated carbonyl groups. It could relate to the polymers miscibility, although displacement of the bands related to the crystallinity of PVA (1142 cm-1, ν(C-C)) and of NaCMC (1372 cm-1, δ(CH)) and displacement of the NaCMC β-glucosidic groups were not identified. As expected, the PVA crystallinity index (ratio between the absorbance of the band at 1142 cm-1 and the band at 2905 cm-1) is lower in the blend, IPVA is ~1.45 and IPVA-CMC is ~1.36, where the presence of NaCMC diminishes the PVA chains ability to pack into crystals[43]. The PVA-CMC layers revealed that the increase of freeze-thawing cycles displaced the band at 1593 cm-1 (NaCMC’s ν(-COO)[44]) to 1566 cm-1, which indicate the effectiveness of the “crosslinking”[45]. In addition, the two bands at 1087 cm-1 (PVA’s δ(C-O-H)[46]) and at 1060 cm-1 (NaCMC’s ν(OCH–O–CH2)[47]) cannot be distinguished in layers submitted to more freeze-thawing cycles, where only one band at 1087 cm-1 can be observed. The PVA-G sample presented some differences when compared to the PVA sample, e.g. the band at 1646 cm-1 was more intense in the PVA-G sample, which could be due to the gelatin ν(C=O)[38]. There was a band at 1169 cm-1 in the PVA-G sample (absent in the PVA sample); the PVA bands at 1142 cm-1 and at 1089 cm-1 were displaced to 1136 cm-1 and to 1101 cm-1 in the PVA-G sample, respectively. The gelatin bands displacement towards higher wavenumbers could be related to reaction products, e.g. from 1680 cm-1 to 1758 cm-1 would represent the formation of esterified product[21]. There were bands at 640 cm-1 and at 612 cm-1 in the PVA-G sample, while there was no band in this region in the PVA sample. A band at 670 cm-1 would be attributed to the gelatin γ(N-H), although bands observed at lower wavelengths were not reported to gelatin[38]. The PVA-G layers revealed similar bands to those of the whole PVA-G sample. Polímeros, 30(3), e2020031, 2020

Layered cryogels laden with Brazilian honey intended for wound care

Figure 2. FTIR spectra of: (a) PVA layers; (b) PVA-CMC layers; (c) PVA-G layers; (d) PVA 3-layered sample, honey and PVA-H 3-layered sample; (e) PVA-CMC 3-layered sample, honey and PVA-CMC-H 3-layered sample; (f) PVA-G 3-layered sample, honey and PVA-G-H 3-layered sample; (g) PVA-H layers; (h) PVA-CMC-H layers; (i) PVA-G-H layers.

The PVA-H sample presents bands related to PVA, e.g. 3275 cm-1, 2940 cm-1, 2910 cm-1, 1564 cm-1, 1379 cm-1, 1239 cm-1, 1143 cm-1. Nonetheless, some of its bands can be related to the presence of honey in the samples, e.g. at 1646cm-1 (ν(C–H) of carboxylic acids, ν(NH3) of free amino acids, water δ(OH)[48]), 1417 cm-1 (δ(O–H) of the C–OH group and δ(C–H) of the alkenes[48]), 775 cm-1 (saccharide configuration; anomeric region of carbohydrates vibration or δ(C–H)[49]). In addition, some of the honey bands are slightly displaced in the PVA-H sample, e.g. the bands at 1054 cm-1 (ν(C–O) of the C–OH group and carbohydrate structure’s ν(C–C)[48]); 1030 cm-1 (vibration of the C–OH group, carbohydrate structure’s ν(C-C) and ν(C-O), phenol’s C-O vibration[50]); 898 cm-1; 874 cm-1, displaced by ~9 cm-1; 819 cm-1, these last bands related to saccharide configuration and to anomeric region of carbohydrates vibration or δ(C–H)[49]. Some of the PVA-H bands are located in between PVA bands position and honey bands position, e.g. the band at 1336 cm-1, which is in between the honey band at 1345 cm-1 (flavanol’s and phenol’s δ(O-H), δ(C-O), δ(C-H) and δ(C=C)[50]) and the PVA band at 1329 cm-1[35]); there is a shoulder at 1097 cm-1, between the PVA band at 1089 cm-1 (ν(C-O)[39]) and the honey band at 1100 cm-1 (ν(C–O) related to the C–O–C linkage[48]); 917 cm-1, which is between the PVA band at 916 cm-1[39] and the honey band at 918 cm-1 (carbohydrate’s δ(C–H)[48]; α and β anomers’ νas(C-O-C)[51]). These bands displacement and the bands located in between the original PVA and honey Polímeros, 30(3), e2020031, 2020

bands could indicate physical interaction (Van der Waals or hydrogen bonding) between PVA and honey. Since the layers of the samples presented varied composition, they were examined separately. The layer containing 5% honey presented bands similar to the PVA-H whole sample previously described, but the bands at 1336 cm-1 and at 898 cm-1 (honey’s bands[52]) were absent and a shoulder at 831 cm-1 (PVA’s band[39]) was present; the band that should be at 1054-1030 cm-1 was displaced towards the PVA band (1089 cm-1[39]), it was at 1083 cm-1. The layer containing 10% honey also present the same bands as the PVA-H whole sample, but the bands at 2940 cm-1 and at 2910 cm-1 were absent. Nonetheless, there were bands at 2930 cm-1 (honey’s band[48]), 1338 cm-1, 1259 cm-1 (honey’s band[48]), 1189 cm-1 (honey’s band[50]). A slight displacement of PVA and of honey bands in the PVA-H layers could be observed with the increased honey content, which could indicate the effect of honey in the movement of PVA’s functional groups and vice versa. The increased level of honey in the layer also revealed more honey bands in the FTIR spectrum, as expected. The PVA-CMC-H sample presented the same bands as the PVA-CMC and that were previously described. In addition, there were bands related to honey, e.g. 2923 cm-1, 1417 cm-1, 1056 cm-1, 1035 cm-1 and 819 cm-1. These bands are slightly displaced in the PVA-CMC-H sample. The layer containing 10% honey and the layer with 5% 5/10

Santos, G. S., Santos, N. R. R., Pereira, I. C. S., Andrade Júnior, A. J., Lima, E. M. B., Minguita, A. P., Rosado, L. H. G., Moreira, A. P. D., Middea, A., Prudencio, E. R., Luchese, R. H., & Oliveira, R. N. honey presented similar bands to PVA-CMC-H whole sample previously described, although there were slight displacements of the bands with the addition of honey, e.g. the band at 1088 cm-1[39] in PVA-CMC layer is broadened and shifted to 1074 cm-1 in the layer containing 10% honey. There is a band at ~1592 cm-1 in the PVA-CMC-H layers that is absent in the PVA-CMC layer and in the honey spectrum. The presence of a new band could indicate new bonding between the polymers and honey[53].

3.3 In vitro analysis

The PVA-G-H sample presented the previously described PVA bands; some gelatin bands, e.g. at 1551 cm-1 and at 1337 cm-1[54]; and some honey bands previously described, e.g. at 1052, 1029, 896, 853, 817, 775 cm-1. Band displacement or the formation of bands related to a chemical reaction or interaction between honey and the polymers was not observed[21].

There was a significant difference of the equilibrium of swelling degree (SD) between all the samples, although the SD of the samples PVA and PVA-CMC-H can be considered the same. In general, it can be observed that samples containing honey swelled less than samples without honey[58], probably because honey occupies network pores created by the ice during freeze-thawing, filling space that could have been available to the saline solution[59]. PVA samples swelled less than PVA-G samples (which have a low degree of crystallinity compared to PVA), and these in turn swelled less than PVA-CMC samples (which have an even lower degree of crystallinity). Samples that are less crystalline present more amorphous phase and thereby stretch more with fluid ingress, and therefore swell more[60,61].

3.2 Thermal analysis DSC analysis of the whole three-layer samples revealed that the addition of the polysaccharide (NaCMC) diminished the sample Tg and raised the sample Tm, compared to PVA, while Xc( PVA−CMC ) was lower than the Xc( PVA) (Table 2). It suggests that the NaCMC chains do not contribute to polymer chain entanglement, probably working as a plasticizer in the amorphous region of the blend. In addition, the NaCMC diminished the degree of crystallinity, but more “perfect” crystals were formed (with a higher Tm). A similar effect was observed when gelatin was added to PVA, but the Tm of PVA and PVA-G was similar, indicating that gelatin contributed only as a plasticizer and to diminish the sample’s degree of crystallinity. The addition of honey to PVA diminished considerably the Tg, Tm and the Xc of the samples, indicating that honey is a physical barrier to the interaction of polymer chains. The addition of both honey and either polysaccharide or protein to PVA led to the absence of crystallinity (there is no peak related to crystal formation). Nonetheless, compared to PVA-H, PVA-CMC-H and PVA-G-H both presented higher Tg, indicating that in the presence of the coupled materials the amorphous chains need more energy to gain movement. The Tg and Tm altering with the addition of NaCMC or honey could indicate miscibility and interaction between the materials[55,56]. The addition of gelatin altered the PVA Tg and Xc also due to the miscibility of the polymers which disorganize the chains packing[38].

All the samples’ swelling degree tests presented a similar trend (Figure 3), where there is a stretching of the network by the fluid’s initial diffusion, followed by a plateau in the curves, where the polymers network reaches stability/ equilibrium, when the elastic forces and the osmotic forces on the network are balanced[57]. Although some samples presented significant differences between the first hours of swelling (p < 0.05), the samples reached swelling equilibrium in 24 h of immersion.

The gel fraction (GF) of all the samples without honey (Figure 3) can be considered similar (p > 0.05). Although the addition of natural polymers altered the degree of crystallinity of the samples, the amorphous chain entanglements and the crystallites formed in freeze-thawing maintained the samples’ structural integrity[62]. The addition of honey decreased the gel fraction of the samples (p < 0.05). The gel fraction property is related to the hydrogel’s crosslinked polymers chains that remain insoluble when immersed in aqueous fluid. The PVA-H GF is significantly higher than the PVA-CMC-H GF which is, in turn, significantly higher than the PVA-G-H, GF (p < 0.05). The GF of the honey samples diminish in accordance with the decrease in the samples’ degree of crystallinity, indicating that the samples crystallites work as physical crosslinking points. The addition of the natural polymers, as well as the addition of honey, diminishes the gels’ degree of crystallinity, diminishing the gels’ GF[63,64]. The samples without honey presented similar weight loss (p > 0.05), Figure 3. The PVA samples weight loss (WL) could be due to biodegradation in saline: “hydrolytic cleavage of hydrogen bonding among -OH groups of PVA chains”, as discussed by Kamoun et al.[65]. The weight

Table 2. Data regarding the samples degree of crystallinity (Xc), glass transition temperature (Tg), melting temperature (Tm), swelling degree (SD), gel fraction (GF), weight loss (WL), S. aureus growth. DSC Samples

PVA PVA-H PVA-CMC PVA-CMC-H PVA-G PVA-G-H

87 80 69 85 79 102

Tm (°C)

6/10

Swelling Xc

142 ± 16 79 ± 8 363 ± 16 113 ± 22 216 ± 7 49 ± 11

89 ± 1 45 ± 2 88 ± 5 36 ± 6 86 ± 1 22 ± 1

(%) 218 150 223 218 -

21 5 19 13 -

S. aureus WL

Growth

Reduction

(CFU/g)

(%)

10 ± 1 54 ± 2 11 ± 5 63 ± 6 13 ± 1 77 ± 1

5,5 x 106 5,7 x 106 7,7 x 105 5,4 x 105 8,7 x 106 7,4 x 106

(%)

29.9 14.9

Polímeros, 30(3), e2020031, 2020

Layered cryogels laden with Brazilian honey intended for wound care

Figure 3. Samples (a) Swelling Degree and (b) Weight Loss and gel fraction.

loss of the PVA blends could be related to PVA hydrolytic cleavage and also to the solubility of the natural polymers in aqueous media[66-68]. The samples containing honey presented significantly higher weight loss/biodegradation than the samples without honey. The WL followed the trend: PVA-H < PVA-CMC-H < PVA-G-H (p < 0.05), hypothesised to be due to degradation (due to chains leaching out by the saline solution) and honey delivery of both samples[69]. It is worth noting that only the three-layered samples were evaluated. In this study, the PVA hydrogel can be considered a negative control (non-bactericidal) against gram-positive bacteria (S. aureus), allowing bacteria to grow in its presence[70]. The addition of honey inhibited the bacterial growth in the presence of the samples for PVA-CMC and PVA-G samples, Table 2. S. aureus attaches to proteins on the surface of collagen matrices and gelatin is a hydrolysed form of collagen[71]. Honey itself has a bactericidal effect due to its low water activity, the presence of H2O2 and its low pH[72]. The antibacterial effect on S. aureus in wound healing hydrogels would depend on exceeding the minimum inhibitory concentration (MIC) of the honey with respect to the amount of honey in, and released from, the samples[73]. Several hydrogels containing honey present antibacterial activity[32,74]. The samples of the present work had 200 mg/mL of honey, but the MIC for this Brazilian honey might not be have been reached for the PVA-honey samples. The PVA-CMC-H and PVA-G-H samples presented activity against S. aureus relative to their counterparts with no honey[59]. It was observed a bactericidal effect. It is known that even low concentrations of honey can stimulate wound healing[75].

4. Conclusions It was observed that the addition of some materials to PVA diminishes the gels crystallinity and gel fraction (related to the degree of crosslinking of the gels), altering the samples ability to swell, where the samples with honey presented lower fluid uptake than samples without it, since Polímeros, 30(3), e2020031, 2020

honey can occupy and obstruct pores. Nonehteless, the samples containing honey presented significantly higher biodegradation (hydrolytic degradation/weight loss) than the samples without honey. The samples with NaCMC or gelatin and honey were the ones that presented the highest activity against S. aureus relative to their honey-free counterparts, showing potential to be used as wound care materials.

5. Acknowledgements The authors thank Conselho Nacional de Desenvolvimento Científico e Tecnológico - CNPq (funding/project 405922/2016-7) for the financial support, Fundação de Amparo à Pesquisa do Estado do Rio de Janeiro - FAPERJ, CAPES (“This study was financed in part by the Coordenação de Aperfeiçoamento de Pessoal de Nível Superior - Brasil (CAPES) - Finance Code 001”), professor Garrett Brian McGuinness and professor Mauricio Mancini, CETEM/UFRJ and EMBRAPA.

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Polímeros, 30(3), e2020031, 2020

ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.05920

Reactive compatibilization effect of graphene oxide reinforced butyl rubber nanocomposites Sathishranganathan Chinnasamy1, Rajasekar Rathanasamy1* , Harikrishna Kumar Mohan Kumar1, Prakash Maran Jeganathan2, Sathish Kumar Palaniappan3 and Samir Kumar Pal3 Department of Mechanical Engineering, Kongu Engineering College, Tamil Nadu, India Department of Food Science and Nutrition, Periyar University, Salem, Tamil Nadu, India 3 Department of Mining Engineering, Indian Institute of Technology, Kharagpur, West Bengal, India 1

*rajasekar.cr@gmail.com

Abstract The objective of this work is to develop graphene oxide (GO) incorporated butyl rubber (IIR) nanocomposites by three different methods: direct addition approach (DAAM), single step method (SSM) and two step method (TSM). Chlorobutyl rubber was used as a compatibilizer in SSM and TSM. Mechanical properties of developed nanocomposites was increased and gas permeability co-efficient was decreased up on addition of GO content in IIR matrix. Maximum technical properties was achieved for the nanocomposite with 1.6 wt.% of GO in all methods was achieved due to better interfacial bonding with IIR matrix. When GO content increases above 1.6 wt.% in IIR matrix leads to agglomeration which resulted in deterioration of mechanical properties. HR-TEM studies revealed that nanocomposites prepared by TSM shows exfoliated structure of GO in IIR matrix due to homogenous distribution when compared to the nanocomposites prepared with DAAM and SSM. Keywords: butyl rubber, graphene oxide, nanocomposite, mechanical. How to cite: Chinnasamy, S., Rathanasamy, R., Kumar, H. K. M., Jeganathan, P. M., Palaniappan, S. K., & Pal, S. K. (2020). Reactive compatibilization effect of graphene oxide reinforced butylrubber nanocomposites. Polímeros: Ciência e Tecnologia, 30(3), e2020032. https://doi.org/10.1590/0104-1428.05920

1. Introduction Nanocomposites are currently used in several fields and new applications are being continuously developed like innerwalls of tires, thin film capacitors for computer chips, O-rings, ball bladders, fiber optic compounds, hand gloves, impellers blades and food packaging, etc[1]. Nanocomposites are materials which fused with nano-sized elements in the matrix to increase the macroscopic properties. The introduction of nanoscale, distribution of filler or measured nanostructures in the base matrix with high surface to volume ratio can enhance the physico-mechanical properties and unique functional behaviours of base material. Rubber nanocomposites is one of the composites which were developed by several researchers around the world using different rubbers such as natural rubbe r(NR)[2-4], butadiene rubber (BR), butyl rubber (IIR)[5], chlorobutyl rubber (CIIR)[6,7], epoxidizednatural rubber (ENR)[8], styrene butadiene rubber (SBR), styrene butadiene styrene rubber (SBS) and nitrile butadiene rubber (NBR) for various applications. Even though IIR has its own merits and its property was enhanced with reinforcement of various nanosized fillers such as carbon black (CB)[9], silica, carbon nanotubes[10], graphene[11,12], graphene oxide (GO)[13,14] and nanoclay[15]. Amongst, GO is one of the filler utilized to improve the mechanical and gas barrier properties of IIR rubber compounds. This work aims to develop

Polímeros, 30(3), e2020032, 2020

IIR nanocomposites by reinforcing GO in presence and absence of compatibilizer using DAAM, SSM and TSM[16]. Mechanical and gas permeability properties of the prepared nanocomposites were evaluated and the properties of the nanocomposites was compared against the preparation method. CIIR was used as a compatibilizer in this work to achieve uniform dispersion of GO in the IIR matrix.

2. Materials Butyl rubber and chlorobutyl with 1.25% of chlorine was purchased from Laxness and Bayer, Mumbai. Graphene oxide with a bulk density of 1.8 g/cm3 and soluble in polar solvents was procured from Sigma Aldrich chemicals, Mumbai. Remaining chemicals like sulphur, zinc oxide, stearic acid[17], N-cyclohexyl-2-benzothiazyl sulphenamide (CBS) and tetramethylthiuram disulphide (TMTD)[18] were obtained from Loba Chemicals, Chennai.

3. Development of Composite Nanocomposite were prepared with and without compatibilizer[19]. Nanocomposite without compatabilizer was developed by directly mixing of GO with rubber

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Chinnasamy, S., Rathanasamy, R., Kumar, H. K. M., Jeganathan, P. M., Palaniappan, S. K., & Pal, S. K. through DAAM. Nanocomposites with compatibilizer has been prepared by two approaches namely, SSM and TSM. CIIR[20-22] was employed as a compatibilizer for both methods and the composition/ratio for the preparation of the composites was given in Table 1. In SSM, two roll mill was utilized to reinforce GO in IIR matrix in presence of compatibilizer. Curing and vulcanizing agents were also added along with rubber and GO in two roll mill. Two roll mill was operated at atmospheric temperature and speed ratio of rotors was maintained at 1:14. Compression molding process was used for vulcanization and specimen preparation. In TSM, Toluene was used as a solvent (150 ml for CIIR and 200 ml for GO) to dissolve and disperse CIIR and GO separately in specific ratio (1:4 w/w). Magnetic stirrer (800 rpm) was used to dissolve CIIR and GO separately until complete dissolution of CIIR and GO was attained. Dissolved CIIR solution was taken in the beaker, kept it in magnetic stirrer, dispersed GO was added directly in CIIR solution at the rate of 5 ml/min until GO solution was completely poured in CIIR solution and stirring process was continued till uniform mixture of CIIR and GO was obtained. Then, mixture was poured in glass plate and kept in plain surface at room temperature without any disturbance until film like material was obtained. Two roll mixing mill[24] was used to mix the acquired film, pure IIR and other ingredients (curing and vulcanizing agents) according to the ratio mentioned in Table 1. Resultant material was molded and vulcanized using compression moulding process. [23]

nanocomposites from BOC Gases, India. Nitrogen gas was passed at a constant pressure of 3.5 bar at 35 °C for a time period of 30 min through the prepared nanocomposite. For isothermal measurement conditions the apparatus was placed in a thermostatically controlled environment.Time lag method is used to calculate gas barrier measurement. Mean permeability coefficient (P)[5] was determined with steady state gas pressure increment(dp/dt) in the calibrated volume (V). The permeability coefﬁcients were calculated from Equation 1 P=

VdTo  dp  Api po T  ds S

(1)

where, P, V, d, To, A, Po, Pi, T and dp/ds)s denotes mean permeability coefficient, calibrated volume, thickness of the film, standard temperature, effective permeation area, standard pressure, upstream side of the film with a gas pressure of 3.5 bar, temperature of measurement and steady state gas pressure increment.

4.3 High Resolution Transmission Electron Microscopy (HR-TEM) Morphology of the prepared composites was analyzed using HR-TEM (JEOL, USA). The thickness of the sample was maintained as 80 nm to record the images of the samples under HR-TEM[29].

5. Results and Discussion 5.1 Mechanical properties

4. Characterization of Nanocomposites 4.1 Mechanical properties Mechanical properties of the samples were determined based on the ASTM D-412-06 standard using a universal tensile testing machine. Specimens for tensile and tear testswere obtained from the molded slabs[25]. Values oftensile strength, elongation at break and tear strength were recorded.

4.2 Gas Permeability Co-efficient (GPC) measurement GPC[26-28]of the developdcompositeswere studiedusing automated diffusion permeameter. Nitrogen gas (XL grade) was purchased to measure the gas permeability for prepared

Mechanical property is one of the dynamic property of any composites for its applications. Hence, in this study, mechanical properties such as tensile strength (TS), tear strength (TES), percentage of elogantion(% E) and modulus at 100 and 300%[30] were determined and outcomes were depicted in Figures 1-4. Mechanical properties of nanocomposites increases as weight content of GO increases in IIR matrix. Nanocomposites containing IIR with GO content of 1.6 wt.% (D4, S4 and T4) exhibited higher TS, TES and modulus at 100 and 300% when compared to pure IIR in all three preparation methods (DAAM, SSM and TSM). TS, TES and modulus at 100 and 300% decreases when GO content increases to 2 wt.% in IIR matrix. While

Table 1. Formulation of rubber nanocomposites. Components

IIR (g) CIIR (g) GO (g) CIIR-GO films Stearic Acid CBS Zinc Oxide TMTD Sulphur

100 2 1 3 1 1.5

IIR Nanocomposites(with compatibilizer) Single step mixing method Two step mixing method parts per hundred rubber of weight (phr) D5 S1 S2 S3 S4 S5 T1 T2 T3 T4 100 98 96 94 92 90 98 96 94 92 1.6 3.2 4.8 6.4 8 2 0.4 0.8 1.2 1.6 2 2 4 6 8 2 2 2 2 2 2 2 2 2 2 1 1 1 1 1 1 1 1 1 1 3 3 3 3 3 3 3 3 3 3 1 1 1 1 1 1 1 1 1 1 1.5 1.5 1.5 1.5 1.5 1.5 1.5 1.5 1.5 1.5

IIR Nanocomposites(without compatibilizer) D1 100 0.4 2 1 3 1 1.5

D2 100 0.8 2 1 3 1 1.5

D3 100 1.2 2 1 3 1 1.5

D4 100 1.6 2 1 3 1 1.5

T5 90 10 2 1 3 1 1.5

CN, D1 to D5, S1 to S5 and T1 to T5 indicates the specimen code for the prepared composites.

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Reactive compatibilization effect of graphene oxide reinforced butyl rubber nanocomposites

Figure 1. Tensile strength of pure and GO filled uncompatibilized and compatibilied IIR nanocomposites.

%E was reduced up on addition of GO content in IIR matrix. TS, TES and modulus at 100 and 300% of nanocomposites prepared with SSM (S4) was increased by 56%, 7.6%, 48% and 59% respectively and %E was reduced by 37% when compared with pure IIR. Similarly, mechanical properties of nanocomposites prepared with TSM (T4) were enhanced by 63%, 23%, 70% and 59% respectively and % E was reduced by 40% when compared to pure IIR. Addition of compatibilier through solution mixing helps in achieving uniform dispersion of GO in IIR matrix. Uniform distribution of GO in IIR matrix leads to form intercalated structure and strong interfacial bonding. So, nanocomposites prepared by TSM shows superior mechanical properties when compared to nanocomposites prepared by other two methods (DAAM and SSM). Mechanical properties starts to deteriorate when reinfocement of GO content increases to 2 wt.% in IIR matrix (D5, S5 and T5) for the nanocomposites prepared by all three methods. Increment in the GO content above 1.6 wt. % in the developed nanocomposites (all methods) leads to form agglomeration of GO in the matrix which produced the stress concentration on the matrix and decreased the TS, TES, modulus.

5.2 Gas permeability co-efficient

Figure 2. Tear Strength of pure and GO filled uncompatibilized and compatibilied IIR nanocomposites.

Gas barrier properties of the prepared nanocomposites were evaluated by determining the GPC. GPC values of the nanocomposites prepared using all the three methods were depicted in Figure 5. GPC of nanocomposites prepared with all three methods (DAAM, SSM and TSM) decreases upon addition of GO content in IIR matrix. Nanocomposites prepared with 1.6 wt.% of GO (D4, S4 and T4) in all three methods shows lower GPC when compared to other samples. GPC of the nanocomposites D4, S4 and T4 was reduced by 25%, 31% and 36% respectively when compared to pure IIR. Lower GPC was observed for T4 due to even distribution of GO in IIR matrix. Even distribution of GO in IIR matrix was achieved through TSM which leads to strong interfacial bonding thereby reduces the passage of nitrogen gas. GPC was increased when GO content increases from 1.6 wt.% to 2wt.% (D5, S5 and T5). It is due to agglomeration of GO in IIR matrix which leads to decrease in interfacial bonding between the matrix.

Figure 3. Elongation of pure and GO filled uncompatibilized and compatibilied IIR nanocomposites.

Figure 4. Modulus (100 and 300%) of pure and GO filled uncompatibilized and compatibilied IIR nanocomposites. Polímeros, 30(3), e2020032, 2020

Figure 5. Permeability coefficient of pure and GO filled uncompatibilized and compatibilied IIR nanocomposites. 3/5

Chinnasamy, S., Rathanasamy, R., Kumar, H. K. M., Jeganathan, P. M., Palaniappan, S. K., & Pal, S. K.

Figure 6. Morphology of nanocomposites at 1.6 wt.% of GO. (a) DAAM; (b) SSM; and (c) TSM.

5.3 HR-TEM analysis HR-TEM images of developed nanocomposites at 1.6 wt.% of GO were depicted in Figure 6. In DAAM and SSM, interactions between IIR and GO were weak due to lower interfacial bonding caused discontinuous phase of GO in IIR matrix (Figure 6a, b) which affects the properties of nanocomposites. Even distribution of GO in IIR matrix leads to better interaction and continuous phase during mixing process of TSM (Figure 6c) and leads exfoliation of GO in matrix of nanocomposites.

6. Conclusion Graphene oxide nanocomposites was successfully developed in presence and absence of compatibilizer using three different methods (DAAM, SSM and TSM). Morphological studies, mechanical properties and GPC of prepared nanocomposites were examined and properties were compared to identify the suitable method for preparing nanocomposites. Nanocomposites (T4) prepared using TSM with compatibilizer exhibited higher mechanical properties (TS, TES and modulus at 100% and 300%) and less GPC when compared to nanocomposite (D4 and S4) prepared with other methods. Enhancement in properties was achieved due to even dispersion of GO in IIR matrix and confirmed thorugh HR-TEM analysis. Increase in weight content of GO in IIR matrix from 1.6 wt.% to 2 wt.% leads to aggromeralation which resulted in deterioration of TS, TES and modulus of nanocomposite (D5, S5 and T5) prepared in all three methods. From the above results two step method (TSM) in presence of compatibilizer was found to be a suitable method to develop graphene oxide nanocomposites.

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22. Liu, C., Fan, J., & Chen, Y. (2019). Design of regulable chlorobutyl rubber damping materials with high-damping value for a wide temperature range. Polymer Testing, 79, 106003. http://dx.doi.org/10.1016/j.polymertesting.2019.106003. 23. Wu, J., Xing, W., Huang, G., Li, H., Tang, M., Wu, S., & Liu, Y. (2013). Vulcanization kinetics of graphene/natural rubber nanocomposites. Polymer, 54(13), 3314-3323. http://dx.doi. org/10.1016/j.polymer.2013.04.044. 24. Potts, J. R., Shankar, O., Murali, S., Du, L., & Ruoff, R. S. (2013). Latex and two-roll mill processing of thermallyexfoliated graphite oxide/natural rubber nanocomposites. Composites Science and Technology, 74, 166-172. http:// dx.doi.org/10.1016/j.compscitech.2012.11.008. 25. Rajasekar, R., Pal, K., Heinrich, G., Das, A., & Das, C. (2009). Development of nitrile butadiene rubber–nanoclay composites with epoxidized natural rubber as compatibilizer. Materials & Design, 30(9), 3839-3845. http://dx.doi.org/10.1016/j. matdes.2009.03.014. 26. Jo, J. O., Saha, P., Kim, N. G., Ho, C. C., & Kim, J. K. (2015). Development of nanocomposite with epoxidized natural rubber and functionalized multiwalled carbon nanotubes for enhanced thermal conductivity and gas barrier property. Materials & Design, 83, 777-785. http://dx.doi.org/10.1016/j. matdes.2015.06.045. 27. Li, L., Zhang, J., Jo, J. O., Datta, S., & Kim, J. K. (2013). Effects of variation of oil and zinc oxide type on the gas barrier and mechanical properties of chlorobutyl rubber/epoxidised natural rubber blends. Materials & Design, 49, 922-928. http:// dx.doi.org/10.1016/j.matdes.2013.02.057. 28. Azizli, M., Naderi, G., Bakhshandeh, G., Soltani, S., Askari, F., & Esmizadeh, E. (2014). Improvement in physical and mechanical properties of IIR/CR rubber blend organoclay nanocomposites. Rubber Chemistry and Technology, 87(1), 10-20. http://dx.doi.org/10.5254/rct.13.87951. 29. Rajasekar, R., Nayak, G., Malas, A., & Das, C. (2012). Development of compatibilized SBR and EPR nanocomposites containing dual filler system. Materials & Design, 35, 878-885. http://dx.doi.org/10.1016/j.matdes.2011.10.018. 30. Kumar, H. K. M., Subramaniam, S., Rathanasamy, R., Pal, S. K., & Palaniappan, S. K. (2020). Substantial reduction of carbon black and balancing the technical properties of styrene butadiene rubber compounds using nanoclay. Journal of Rubber Research, 23(2), 79-87. http://dx.doi.org/10.1007/ s42464-020-00039-7. Received: June 19, 2020 Revised: Sept. 21, 2020 Accepted: Sept. 25, 2020

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ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.04020

Nanoscale morphology, structure and fractal study of kefir microbial films grown in natura Robert S. Matos1,2* , Ellen C. M. Gonçalves1 , Erveton P. Pinto1 , Gerson A. C. Lopes1,3 , Nilson S. Ferreira4  and Cristiane X. Resende1  Laboratório de Biomateriais, Programa de Pós-graduação em Ciência e Engenharia de Materiais, Universidade Federal de Sergipe – UFS, São Cristóvão, SE, Brasil 2 Grupo de Materiais Amazônicos, Departamento de Física, Universidade Federal do Amapá – UNIFAP, Macapá, AP, Brasil 3 Laboratório de Física, Universidade do Estado do Amapá – UEAP, Macapá, AP, Brasil 4 Departamento de Física, Universidade Federal de Sergipe – UFS, São Cristóvão, SE, Brasil 1

*amazonianmaterialsgroup@gmail.com

Abstract Kefir is a natural probiotic produced by kefir grains fermentation. Biofilms produced from fresh kefir grains in natura were studied for presenting structural characteristics that will be of great interest in the area of regenerative medicine. This work presents a study on the surface of kefir biofilms, obtained by the cultivation of kefir grains in commercial white sugar. Four different films were produced, varying the concentration of sugar. The crystallinity of the biofilms was analyzed and revealed that sugar concentration influences biofilm amorphousness. Morphology showed that the biofilms presented excellent superficial adhesiveness. Fractal parameters were studied and revealed that there was homogeneity in the biofilm microtexture. Both fractal succolarity and surface entropy showed that the degree of water penetration and topographic homogeneity of the biofilms was not influenced by sugar concentration. These results show that kefir biofilms have excellent structural and morphological properties to be used in the biomedical field. Keywords: kefir, biofilms, crystallinity, morphology, fractal. How to cite: Matos, R. S., Gonçalves, E. C. M., Pinto, E. P., Lopes, G. A. C., Ferreira, N. S., & Resende, C. X. (2020). Nanoscale morphology, structure and fractal study of kefir microbial films grown in natura. Polímeros: Ciência e Tecnologia, 30(3), e2020033. https://doi.org/10.1590/0104-1428.04020

1. Introduction The use of natural products in the treatment of several diseases has been studied since ancient times. Historically, human disease problems were treated using drugs extracted from classical medicinal plants. Skin disorders were treated with the use of natural plant extracts by traditional peoples a long time ago. However, at the beginning of the last century, there was an increase in the development of synthetic drugs that could definitively treat or cure these problems, but which later came up against the problem of the toxicity of these drugs. Currently, regenerative medicine has again turned to the use of natural products that may be more biocompatible and less cytotoxic to humans[1]. Biopolymers have had high relevance in this aspect because they are biodegradable[2] and accurately extracted from natural products. Kefir is a natural product well-known in traditional medicine as a probiotic drink produced from kefir grains[3]. The kefir grains are small rigid yellowish granules with irregular shapes[4] containing bacteria and yeasts. Moreover, kefir grains also have in their composition proteins and polysaccharides[5]. Kefir grains are composed of a mixture of several bacteria, including various species of lactobacilli, lactococci, leuconostocs and acetobacteria, and yeasts (with

Polímeros, 30(3), e2020033, 2020

or without lactose fermentation)[6]. This ensures that its composition is extremely diversified and rich, presenting great potential for regenerative medicine applications; however, the organization of the microorganisms in the grain is not entirely known[7]. The grains are capable of producing an exopolysaccharide matrix, known as kefiran, which is much studied today. This polysaccharide has been presented as an excellent alternative for application in some fields such as tissue engineering[8], regenerative medicine[9], and drug delivery[10,11]. However, this polysaccharide matrix has also been observed to harbor microorganisms such as bacteria and yeasts[3]. Microbial films can be found in the natural aquatic environment, soil, living tissues, medical devices, or piping systems for potable or industrial water. As a result of the symbiotic cooperation between bacteria, they form a layer on the surface that remains in constant maintenance[12] until the environment is no longer favorable[13]. Thus, bacteria can quickly form complex living communities, preferably in aggregates[14]. Therefore, the biofilms can be understood as a community of microorganisms that tend to adhere to

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Matos, R. S., Gonçalves, E. C. M., Pinto, E. P., Lopes, G. A. C., Ferreira, N. S., & Resende, C. X. wet surfaces, to multiply themselves, and to get soaked in a viscous matrix composed of extracellular polymeric substances (EPS)[15,16]. Thus, it is believed that a similar polysaccharide matrix can be produced by kefir grains in natura, by the formation of a biofilm on the substrate[17,18], which can be a natural juice, a sugar solution, cane molasses or animal milk. Nowadays, the most used substrate for growing water kefir grains is brown sugar because it is a product that undergoes little processing. However, from an economic point of view, this type of sugar is the most unviable and has high-cost compared to commercial white sugar. Furthermore, it worth highlighting that white sugar has fewer nutrients and more sucrose as compared to brown sugar[19,20]. Thus, it is interesting to study the consequences of using commercial white sugar for the production of biofilms for a potential application in regenerative medicine. Some researches have reported that kefiran films can have distinct applications on the food[21] and pharmaceutical industries[3]. The reason for the application kefir biofilms on the biomedical field (e.g., biocurative) is that it presents a uniform distribution of bacteria that can act to protect the wound against pathogens microorganisms, moreover, according to Coma et al.[2] biodegradable films can be functional, acting as antimicrobial or antioxidant agents. Atomic force microscopy (AFM) has been successfully used to study this kind of biological surface as a result of its high sensitivity, e.g., see[3,22]. Several quantitative parameters are used to describe the information contained on the surface. As an example, we can mention the determination of the number of cells present in the biological tissue, the calculation of the shapes of the contours of a cell, or even the determination of the distribution of a specific population of a group of cells[23]. However, more quantitative parameters, such as roughness, asymmetry, and kurtosis, can provide accurate information about the topographic condition of the surface. In this present work, we are introducing three new parameters, such as surface entropy, fractal succolarity, and fractal lacunarity, to study biofilms surface texture. These parameters were successfully presented by Ştefan et al.[24] for the description of the homogeneity behavior of the superficial microtexture. A recent study has suggested that the surface of kefir biofilms has excellent superficial adhesiveness and is semi‑crystalline, even when produced with brown sugar[3]. In this research, we propose the synthesis of these biofilms using white sugar as a substrate because of its lower cost. Thus, this work aims to evaluate the behavior of the structure, morphology, and microtexture on the surface of kefir biofilms. The analyzes were made using different methodologies such as X-ray diffraction (XRD), scanning electron microscopy (SEM), and atomic force microscopy (AFM). These results are significant in evaluating the possible applications of these exopolysaccharide-based biofilms.

2. Materials and Methods 2.1 Materials The kefir grains were obtained at the Drugs Research Laboratory of the Federal University of Amapá, Brazil, and 2/10

kept refrigerated for conservation. The white and brown sugars used in this research were purchased at a local supermarket in the Macapá/AP, Brazil. The grains were submitted to a process of stabilization of the microorganisms during 15 days, for maintenance of the grains viability. In this regard, a stirred white sugar solution (15 g.L-1) was prepared. Then, the grains were added in this solution in a ratio 1:10 (w/v). The stirred solution was changed daily and the temperature kept at 27±2 ºC.

2.2 Biofilms obtention The biofilm cultivation methodology previously proposed by Oliveira et al.[17] was used. The biofilms were produced in a solution containing distilled water, kefir grains, and commercial white sugar. Four different experiments were performed, where 40 g.L-1 of Kefir were inoculated in the following concentrations of white sugar: 20 g.L-1, 40 g.L-1, 60 g.L-1, and 80 g.L-1, respectively. A film grown only in the presence of brown sugar solution (40 g.L-1) was obtained to be used as a basis for comparisons to the analysis of the surface topography and crystallography of the biofilms, following the same methodology proposed by Matos et al. [3] . This film was named control. Afterward, all the produced samples were kept at 25±2 °C for 25 days during its formation and posteriorly deposited on glass slides (Rectangular Microscope Slide Coverglass).

2.2 Characterization of the biofilms 2.2.1 X-Ray Diffraction (XRD) The biofilms were cut out while still on the glass slide in dimensions that varied between 2 and 4 cm2 and kept in a sample holder made of glass and with side coating of aluminum. The measurements were performed at room temperature on a Miniflex II Rigaku diffractometer using Cu K-α (=1.542Å) tube operated at 40 kV and 1.2 mA, a scan range from 5 to 45º (2θ), a step size of 0.02º, and scan‑speed of 1°/min. The contributions of the amorphous and crystalline phases to the diffraction pattern was characterized by the empirical structure factors calculated via Le Bail method[25] because of the large overlapped peaks on XRD patterns. The analysis was performed using the Fullprof program[26], with the space group and unit cell parameters found in the indexing. Peak profiles were modeled using a pseudo-Voigt peak shape function in order to fit the various parameters to the data point such as one scale factor, one zero shift, background, three cell parameters, width of the peaks, and one overall thermal factor. The crystallinity percentage of the biofilms was estimated, separating the integrated intensities from the crystalline and noncrystalline phases of the Le Bail fitted XRD patterns using Equation 1[27]: = Cest . 100 I c / ( I a + I c )

(1)

where Ic and Ia are, respectively, the crystalline and noncrystalline integrated intensities. Polímeros, 30(3), e2020033, 2020

Nanoscale morphology, structure and fractal study of kefir microbial films grown in natura 2.2.2 Atomic Force Microscopy (AFM)

2.2.2.3 Advanced fractal analysis

2.2.2.1 Roughness parameters The topographic analysis of the biofilms was performed using an Atomic Force Microscope (AFM) Nanosurf EasyScan 2 controller in contact mode, with silicon cantilever ContAL-G having a resonance frequency of 13 kHz and elastic constant of 0.2 N/m. The samples were fixed on the sample holder with double-sided tape. The region of analysis was of 30 µ m2, with a scan

rate of 0.7 seconds per line and a contact force of 19.2 nN. Twenty images were taken of each sample to obtain a more representative average of the medium roughness RM, and root-mean-square roughness RRMS, which are obtained using Equation 2 and Equation 3, respectively. Rm = ∬ Z ( x, y ) dx.dy

(2)

where RM is medium roughness evaluated on the entire surface, and Z(x,y) is the height function of the vertical profile. RRMS =

2 Z ( x, y ) dx.dy

(3)

where RRMS is the Root Mean Square Roughness , and both parameters were determined by the software WSxM 5.0 version (Nanotec Eletronica S. L.). [28]

2.2.2.2 Topographic homogeneity In order to obtain the superficial entropy of biofilms, an algorithm developed in R language was used, as in our previous works (for more details see, e.g[3,24]). This algorithm considers that the surface entropy can be calculated from a height matrix that was provided by AFM images. The WSXM software was used to convert the pixels of the AFM images to binary images of heights. The uniform patterns of biofilms were studied as parameters because they represent the uniformity of the surface, based on its height distribution. The algorithm was developed to consider this matrix with an NxN dimension. Thus, Shannon entropy was used to determine uniform and non-uniform height patterns related to the binary matrix, using Equation 4[29]: N N 2 H ( ) = − ∑ ∑ pij log pij =i 1=j 1

(4)

where, pij is the array that converts each pixel value to height. After that, the algorithm normalized the values to find the patterns so that the normalized entropy value varies from 0 to 1 using a normalization obtained by Equation 5: H matr alt =

2 (2) H ( ) − H min 2 2 H − H( ) max

(5)

min

( 2 ) represents the surface with minimum uniform where H max (2) patterns and H min represents the non-uniform pattern surface. We focused our analysis only on the uniform ( 2 ) standard values below represented only by the letter H max H. It is also worth highlighting that this algorithm is not commercially provided.

Polímeros, 30(3), e2020033, 2020

The complexity of a surface can reveal patterns of surface texture. When this parameter is combined with other parameters such as percolation and gap distribution, it can be a powerful tool for analyzing the fractal behavior of a surface. Fractal characterization was performed by calculating the fractal dimension (FD), fractal succolarity (FS), and fractal lacunarity (FL). Fractal dimension is a parameter linked to surface texture and is related to texture homogeneity, as well as fractal lacunarity. However, percolation is related to the homogeneity of fluid entry over the surface texture. FD was measured using the free software Gwyddion 2.47. Fractal succolarity was calculated using the same binary matrices used in the calculation of surface entropy. The adopted procedure was the same one carried out in our previous works, see, e.g[24]. In summary, according to Melo and Conci[30] it is necessary to use Equation 6. FS (T ( k ) , dir ) =

∑ k =1P0 (T ( k ) ⋅ PR (T ( k ) , pc ) n ∑ k =1PR (T ( k ) , pc ) n

(5)

where dir is the water inlet direction, PR is the occupation pressure, T(k) is boxes of equal sizes T(n), pc is the position of the centroid (x, y) of pressure applied to the calculated box, and Po(T(k)) is the occupation percentage. The FS and FL were calculated using an algorithm developed by our group (see, e.g.,[24]). For the calculation of FL, the Box‑Counting‑Differential algorithm[31] was adapted for the programming language FORTRAN 77[22]. Moreover, the same binary matrices were used. The box containing small squares depends on the sample size and has been denoted r x r. The Equation 6 then calculates the lacunarity for the box size: L (r ) =

[ M 1]2

(6)

where M 1 = ∑ sP ( s, r ) and M 2 = ∑ s 2 P ( s, r ) are the first and second moments of the distribution P (s, r), respectively. Fractal lacunarity decreases with increasing the size r of the selected box, according to the power-law of the Equation 7[22]: L (r ) = αrβ

(7)

where the exponent β < 0 can be estimated as the angular coefficient of the log [L(r)] versus log (r) curve. The analysis of microtexture homogeneity was focused on β beta, which is the quantitative parameter of fractal lacunarity. 2.2.3 Scanning Electron Microscopy (SEM) The film morphology was evaluated using scanning electron microscopy (SEM) with a JEOL 5700 microscope. The samples were previously coated with gold, and the images were obtained with an accelerating voltage of 10 kV. 2.2.4 Statistical analysis Analyses of variance (ANOVA) were carried out (p<0.05) to evaluate significant differences between the means for the topographic and fractal parameters. The analyzes were based on twenty different measurements. The Tukey test was applied to discriminate differences between treatments when necessary using OriginPro© 8.0 software (trial version). 3/10

Matos, R. S., Gonçalves, E. C. M., Pinto, E. P., Lopes, G. A. C., Ferreira, N. S., & Resende, C. X.

3. Results and Discussion 3.1 XRD Analysis

Figure 1. XRD patterns of the control film and biofilms grown with kefir in white sugar subtract.

The influence of the white sugar on the crystalline structure of the biofilm was investigated by XRD analysis. The XRD patterns of the control film and biofilms prepared with sugar concentrations of 20 g.L-1, 40 g.L-1, 60 g.L-1, and 80 g.L-1 are shown in Figure 1. All XRD patterns are typical of amorphous-crystalline structure. The Kefir biofilms prepared with a sugar concentration of 20 g.L-1 exhibited four diffraction broad peaks at 2θ≈10.45°, 12.83°, 18.71°, and 20.90°. However, XRD peaks intensities decreased as the concentration of sugar in the films was increased from 40 g.L-1 to 80 g.L-1, indicating that the introduction of sucrose into water kefir EPS induces a decrease in crystallinity. Further, it is also observed that reflection peak at 2θ≈12.83° disappeared and other additional low-intensity peaks are observed at 2θ≈30.13° and 42.16° for sugar concentration higher than 40 g.L-1. Moreover, a comparison of the XRD patterns shows that the angles of the diffraction peaks were not noticeably shifted toward a higher angle and broadened for the biofilms prepared with sugar concentration

Figure 2. Le Bail refined XRD pattern of kefir biofilms prepared with sugar concentrations of (a) 20 g.L-1, (b) 40 g.L-1, (c) 60 g.L-1 and (d) 80 g.L-1. 4/10

Polímeros, 30(3), e2020033, 2020

Nanoscale morphology, structure and fractal study of kefir microbial films grown in natura higher than 20 g.L-1. This can probably indicate increases in the formation of chain regularity in amorphous regions and a decrease in the degree of crystallinity. The XRD peaks indexing were performed with the program McMaille[32], which gives the monoclinic system ( P121 / c1space group) with a≈10.86 Å, b≈14.59 Å, c≈27.48 Å, and α=β=γ=90° and a factor of merit ranging from 20 to 30. This information was selected as the starting model structure for the Le Bail Rietveld refinement of the observed X-ray diffraction profile. The XRD pattern and the corresponding Le Bail refinement result for the Kefir biofilms are shown in Figure 2. Le Bail refinements yielded an acceptable result (χ2≈ 1.5-2.0) and confirmed a significant increasing of the sample amorphous character as sugar concentration is increased from 20 g.L-1 to 80 g.L-1, as given in Table 1. Furthermore, it is also possible to see that the crystallinity percentage decreased almost linearly for sugar concentrations between 20 g.L-1 and 60 g.L-1, indicating that intramolecular and intermolecular bonds of the Kefir structure are weakened as the sugar concentration increases. However, there was a drastic decrease when the concentration was between 60 g.L-1 and 80 g.L-1. This suggests that there was a saturation of the formation of crystalline phases, which can be associated with the formation of the biofilm. Furthermore, it worth mentioning the uncontrolled character of biofilm growth as a consequence of the random process of formation from in vivo biofilm[33]. The obtained results show that Kefir biofilms cultivated in white sugar have excellent flexibility indicating the semi-crystalline character of these materials[34]. In general, a crystalline partial behavior of the biofilms is observed when white sugar concentrations are between 20 g.L-1 and 60 g.L-1 (Table 1). Furthermore, our results also indicate that biofilms structure is relatively balanced, containing amorphous and crystalline phases (57.5 against 42.5%), especially for 20 g.L-1 concentration. These results are quite close to that previously reported by Matos et al.[3] for kefir biofilms grown in brown sugar solutions. Therefore, our results suggest that the kefir biofilms grown in white sugar have a structure similar to that of biofilms grown in brown sugar solutions. However, the presence of more crystalline phases was observed, which was not seen in the study of Matos et al.[3] and Ghasemlou et al.[21]. This result shows that white sugar altered the crystalline structure of biofilms, mostely in the biofilms prepared using lower sugar concentrations. Microorganisms probably consume sugar more quickly at low concentrations and secrete a more crystalline EPS matrix. Thus, one of the causes of the emergence of such phases would be the strengthening of the intramolecular and intermolecular bonds of the biofilm, whose characteristic is very similar to one of the polymers, which spreads the potentiality of application of this material.

3.2 AFM and Advanced fractal analysis Recently, several researchers such as Ghanbarzadeh and Oromiehie[35], Bergo et al.[36], Ghasemlou et al.[21] and Matos et al.[3] have studied films or biofilms with some biotechnological applications such as food packaging and biocurative. However, only in the last two studies, the surfaces characteristics of these films were evaluated by Polímeros, 30(3), e2020033, 2020

Table 1. Crystallographic profile of the kefir biofilms cultivated in white sugar and control film. Concentration Control Film 40 g.L-1 Kefir Biofilms 20 g.L-1 40 g.L-1 60 g.L-1 80 g.L-1

Crystallinity (%)

Amorphicity (%)

0.1

99.9

42.5 34.5 29.6 0.5

57.5 65.5 70.4 99.5

Figure 3. Topographic Images of 3D deflexion (left) and 2D deflexion (right) of the film without kefir. The surface details presented some trace of microorganisms.

AFM. Therefore, considering that topographic analysis can reveal many morphologic characteristics of the materials, in this work, we studied surface adhesion using topographic parameters such as roughness and surface entropy. Figure 3 shows both 3D and 2D (deflection) AFM images of the control film obtained in this analysis. The 3D images show the behavior of the peak distribution in the sample and help to understand the roughness parameter. On the other hand, the biofilms of kefir grown in white sugar present an abundance of microorganism of the microbiota of kefir grains, as presented in the topographic images of AFM, Figure 4. As can be seen, between 20 g.L-1 and 40 g.L-1 samples, there is more abundance of lactobacilli and between 60 g.L-1 and 80 g.L-1, more yeasts. The presence of these microorganisms in biofilms was expected because the kefir microbiota is formed mostly by lactobacilli and yeasts. A similar result was found by Almeida et al.[37] for kefir biofilms prepared with Açaí extract. This fact is fundamental for the analysis of the biotechnological applicability of biofilms. The quantitative parameters related to the topographic irregularities of the biofilms are displayed in Table 2. The statistical analysis through Tukey test revealed that only kefir biofilms with 60 g.L-1 and 80 g.L-1 had RRMS and Rm significant compared to control film. This means that higher concentrations of sugar affected the film roughness. However, no statistical difference was found among the kefir biofilms. These analyzes showed that all biofilms with kefir are relatively similar to the one found by Matos et al.[3] to biofilms with brown sugar. However, microbiologically, the surface of the control film presents traces of some microorganisms that should not exist. Technically, as already mentioned they can be microorganisms that emerge as result of bacteria resulting from the artisanal fabrication process of sugar or 5/10

Matos, R. S., Gonçalves, E. C. M., Pinto, E. P., Lopes, G. A. C., Ferreira, N. S., & Resende, C. X.

Figure 4. Topographic images of 2D deflexion (left) and 3D deflexion (right) with four different concentrations of white sugar: (a) 20 g.L-1, (b) 40 g.L-1, (c) 60 g.L-1, (d) 80 g.L-1. Table 2. Mean values of roughness parameters and topographic homogeneity obtained from atomic force microscopy images; RM, RRMS, and H. The average results were expressed as mean value and standard deviation. Concentration Control Film 40 g.L-1 Kefir Biofilms 20 g.L-1 40 g.L-1 60 g.L-1 80 g.L-1

RM (nm)

RRMS (nm)

327.234±99.100 412.670±101.700 0.988±0.004 406.955±123.700 395.625±99.500 482.755±150.300 480.305±108.200

509.030±155.100 499.010±125.200 600.760±179.900 598.865±136.400

0.982±0.019 0.975±0.021 0.980±0.023 0.985±0.023

Table 3. Mean values (and standard deviation) of FD, FS and β obtained from atomic force microscopy images. The average results were expressed as mean value and standard deviation. Concentration Control Film 40 g.L-1 Kefir Biofilm 20 g.L-1 40 g.L-1 60 g.L-1 80 g.L-1

2.450±0.045

0.689±0.082

0.078±0.010

2.310±0.087 2.370±0.046 2.420±0.033 2.430±0.040

0.536±0.031 0.531±0.049 0.516±0.046 0.532±0.034

0.058± 0.008 0.067± 0.010 0.057± 0.009 0.058± 0.010

when there is improper storage of the product, once it is commercial brow sugar. In Matos et al.[3], where the brown sugar concentration was kept constant at 40 g.L-1, and the kefir concentration was varied, the lower value to roughness found was 671.515 nm, while in the study presented it was 499.010 ± 125.200 nm. This result is too slightly higher than those found by Ghasemlou et al.[21] for kefiran film. This high roughness value is considered a positive factor for surface adhesion because according to Kantorsk and Pagani[38] high roughness is associated with excellent 6/10

adhesiveness. Likewise, high topographic uniformity (H~1) contributes to good uniformity of the surface adhesion. Moreover, Matos et al.[3] have already shown that kefir biofilms that have high surface uniformity would not allow the formation of bubbles when adhered to some surface, reflecting on the interaction of these biofilms with the human skin, like a potential application of these biofilms as natural curatives. Another important conclusion is that sugar concentration does not affect surface homogeneity. Biofilms with high surface uniformity were previously reported, e.g[3,39]. However, most important is the fact that the biofilms displayed microorganisms deposited on the surface, which can be considered excellent for their applicability, in addition to topographic parameters whose values suggest good wettability and surface adhesion. Kefir microbiota is mainly composed of bacteria of the genus Lactobacillus[40,41] and yeasts[42]. Lactobacillus has a cylindrical and elongated shape, while the yeasts usually are blastoconidium shaped[38]. Thus, from the Figure 4 it can be seen that the emergence of yeasts are more evident at the sugar concentrations of 60 (Figure 4c) and 80 g.L-1 (Figure 4d) while Lactobacillus is most observed in 20 (Figure 4a) and 40 g.L-1 (Figure 4b). This can be explained by the fact that an increase in sugar concentration increases fermentation in the sample, which makes the environment conducive to the growth of yeasts. This same behavior was observed by Domingues et al.[43] for sugar cane, from where commercial sugar is derived. Thus, at low concentrations, Lactobacillus can rapidly consume the sugar, preventing the abundant growth of the yeasts. Regarding spatial complexity of the biofilms, the Table 3 shows the values related to the three fractal parameters studied. The Tukey test revealed that FD of kefir biofilms with 60 and 80 g.L-1 was no significant in relation to control film. These higher concentrations showed greater spatial complexities (major irregularities). Hence, kefir biofilms with 20 and 40 g.L-1 exhibited lesser spatial complexities. Polímeros, 30(3), e2020033, 2020

Nanoscale morphology, structure and fractal study of kefir microbial films grown in natura For FS, Tukey test showed that all kefir biofilms were significant in relation to control film. However, between kefir biofilms there was no significant difference, indicating that kefir biofilms prepared with white sugar have similar permeability. Likewise, lacunarity coefficient β was significant for all kefir biofilms in relation to control film. The control film presented the highest β, with the most heterogeneous surface texture. However, among biofilms, only the biofilm with 40 g.L-1 had β significant in relation to the others. As this concentration exhibited one of the highest β values (Table 3), this biofilm had the least homogeneous texture. Therefore, the films that exhibited the most uniform textures had 20, 60, and 80 g.L-1 of white sugar. Fractal analyses confirmed (along with the roughness value) the increase in complexity of patterns existing in the nanotexture of the biofilms. Moreover, the fractal succolarity values did not present significant differences between biofilms, showing that the white sugar concentration does not influence the degree of water penetration on the biofilms’ surface. These statistical results prove that all biofilms have approximate surface percolation, suggesting that the variation in the white sugar concentration does not influence the surface porosity. However, fractal lacunarity did not exhibit a consistent pattern of nanotexture homogeneity according to the increase in sugar concentration in the biofilm structure. This fact can be explained by the high randomness of the formation of membranes in vivo because the mechanisms of the formation of biofilms depend of others factors, in addition to the availability of energy for microbial development. Furthermore, Israelachvili[33] emphasizes that the in vivo membrane formation does not follow a predetermined pattern. This means that the formation of our biofilms was not a predominantly controlled event, which was to be

expected because it was a development in natura. This occurred because they are biofilms produced with living microorganisms, which also justifies their high roughness value. This can also occur because of the increase of the coalescence in EPS matrix as a result of an increase in the sugar concentration once every living organism has its emulsifying system[44] capable of changing the surface properties that makes it a more complex system. Furthermore, it is possible to observe that the microorganisms are randomly arranged on the surface, as we look at Figure 4. In comparison to the XRD results, it is possible to notice that biofilms with high white sugar concentration have a better bacteria development and a better adhesiveness concerning film controlling. What seems more consistent is that biofilms produced with brown sugar develop better than biofilms with white sugar, as reported by Matos et al. [3] , where a better distribution of bacteria and yeasts in all samples was observed. Whether this potential curative can be produced, the difference between the commercial white sugar and brown sugar is the price because commercial white sugar has a lower price than brown sugar, which can cheapen the cost of this product. Although, this research reinforces the possibility of use, as natural dressings, of the kefir biofilms made in white sugar, having insight into his antibacterial activity, considering the maintenance of its antimicrobial potential by preferential deposition of microorganisms, mainly observed for sample with 20 g.L-1.

3.3 SEM analysis Figure 5 shows the most relevant SEM images obtained for the 20 g.L-1 biofilms. The micrographs of this biofilm show structures similar to lactobacillus and yeast, corroborating the results described by AFM. The observed topography shows

Figure 5. Transversal section SEM images of the 20 g.L-1 biofilm: (a) x500, (b) x1.000, (c) x2.000 and (d) x5.000. Polímeros, 30(3), e2020033, 2020

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Matos, R. S., Gonçalves, E. C. M., Pinto, E. P., Lopes, G. A. C., Ferreira, N. S., & Resende, C. X. the presence of structures with apparent unfilled voids on the film’s EPS matrix. Figure 5a shows details of the thickness of the biofilm obtained with an average value of 57.10 µm which classifies it as a thin film, and Figure 5d shows the arrangement and shape of bacteria and yeast. Furthermore, Figure 5b and 5c show how the microorganisms are tangled in the exopolysaccharide matrix. These micrographs show microstructures similar to those obtained by Oliveira et al.[18], where a random arrangement of the microorganisms on the exopolysaccharide matrix of the films was observed, which is a characteristic of the formation of polymicrobial films[45]. This fact, combined with the high roughness observed by AFM and high topographic uniformity, shows that surface adhesion is excellent and uniform. In addition, as there is biological activity on the surface, the interaction between the film and the surface of human skin will likely be beneficial. Considering the pharmacotherapeutic characteristics of kefir, the product generated by these biofilms may act fighting infections and accelerating the skin healing process. This is in agreement with what was discussed by Matos et al.[3] and Oliveira et al. [17] , where both produced brown sugar kefir films and showed similar results. It is important to report that kefiran-based films, which is the polysaccharide present in kefir grains, are highly hydrophilic[21]. This leads to high wettability that allows fluid and gas exchange with the surrounding environment. Moreover, this can help in cellular respiration when the film is acting as a bio-curative for human skin. Aditionally, it is essential to mention that the use of white sugar instead of brown sugar does not significantly affect the roughness, surface uniformity, and probable wettability of the biofilms.

4. Conclusions In this research, we performed studies about the crystallography, topography, fractality, and morphology of kefir biofilms cultivated in commercial white sugar. XRD results demonstrated that the biofilms have a semi-crystalline structure for all the samples (crystallinity about 50%), with resembling structure to other biopolymers like Kefiran. The topological analysis by AFM showed that the biofilms increased their roughness with increasing sugar concentration, indicating that samples have excellent adhesiveness. Surface morphologies presented structures similar to those of bacteria of Lactobacillus genus and yeasts distributed randomly in the samples. The randomness was confirmed with the measures of fractal parameters, where it was seen greater pattern complexity of texture in samples with higher sugar concentration, although a robust pattern of surface texture homogeneity has not been observed. SEM analysis confirmed the presence of microorganisms and microstructural details on the surface of the biofilm, showing that the symbiotic association of bacteria and yeast occurs throughout most of the film. All the results appoint that biofilms developed in the environment with white sugar have high roughness comparing to values referent to kefiran membrane, but less roughness compared to kefir biofilms grown in brown sugar. In addition, they can be used as natural dressings because their adhesiveness was relatively high, and the surface percolation was not affected by the increase of the white 8/10

sugar concentration. Comparatively, biofilms produced with white sugar showed surface characteristics similar to those produced with brown sugar. Therefore, they can be a low-cost alternative of great interest for regenerative medicine using as a natural skin dressing.

5. Acknowledgements The authors acknowledge the access to the experimental facilities at Drugs Research Laboratory of the Federal University of Amapá, Brazil. Ellen C. M. Gonçalves acknowledges the Research Department of the Federal University of Amapá, Brazil, for the scholarship. Nilson S Ferreira acknowledges Conselho Nacional de Desenvolvimento Científico e Tecnológico - CNPq for the financial support under grant No. 309054/2019-2.

6. References 1. Mao, A. S., & Mooney, D. J. (2015). Regenerative medicine: current therapies and future directions. Proceedings of the National Academy of Sciences of the United States of America, 112(47), 14452-14459. http://dx.doi.org/10.1073/ pnas.1508520112. PMid:26598661. 2. Coma, M. E., Peltzer, M. A., Delgado, J. F., & Salvay, A. G. (2019). Water kefir grains as an innovative source of materials: study of plasticiser content on film properties. European Polymer Journal, 120(109234), 1-9. http://dx.doi.org/10.1016/j. eurpolymj.2019.109234. 3. Matos, R. S., Lopes, G. A. C., Ferreira, N. S., Pinto, E. P., Carvalho, J. C. T., Figueiredo, S. S., Oliveira, A. F., & Zamora, R. R. M. (2018). Superficial characterization of kefir biofilms associated with açaí and cupuaçu extracts. Arabian Journal for Science and Engineering, 43(7), 3371-3379. http://dx.doi. org/10.1007/s13369-017-3024-y. 4. Güzel-Seydim, Z. B., Seydim, A. C., Greene, A. K., & Bodine, A. B. (2000). Determination of organic acids and volatile flavor substances in kefir during fermentation. Journal of Food Composition and Analysis, 13(1), 35-43. http://dx.doi. org/10.1006/jfca.1999.0842. 5. Garrote, G. L., Abraham, A. G., & De Antoni, G. L. (2001). Chemical and microbiological characterization of kefir grains. The Journal of Dairy Research, 68(4), 639-652. http://dx.doi. org/10.1017/S0022029901005210. PMid:11928960. 6. Otles, S., & Cagindi, O. (2003). Kefir: a probiotic dairycomposition, nutritional and therapeutic aspects. Pakistan Journal of Nutrition, 2(2), 54-59. http://dx.doi.org/10.3923/ pjn.2003.54.59. 7. Pogačić, T., Šinko, S., Zamberlin, Š., & Samaržija, D. (2013). Microbiota of kefir grains. Mljekarstvo, 63(1), 3-14. http:// dx.doi.org/10.3923/pjn.2003.54.59. 8. Radhouani, H., Bicho, D., Gonçalves, C., Maia, F. R., Reis, R. L., & Oliveira, J. M. (2019). Kefiran cryogels as potential scaffolds for drug delivery and tissue engineering applications. Materials Today Communications, 20(100554), 1-6. http:// dx.doi.org/10.1016/j.mtcomm.2019.100554. 9. Radhouani, H., Gonçalves, C., Maia, F. R., Oliveira, J. M., & Reis, R. L. (2018). biological performance of a promising kefiran-biopolymer with potential in regenerative medicine applications: a comparative study with hyaluronic acid. Journal of Materials Science. Materials in Medicine, 29(124), 1-10. http://dx.doi.org/10.1007/s10856-018-6132-7. PMid:30051294. 10. Blandón, L. M., Islan, G. A., Castro, G. R., Noseda, M. D., Thomaz-Soccol, V., & Soccol, C. R. (2016). Kefiran-alginate gel microspheres for oral delivery of ciprofloxacin. Colloids Polímeros, 30(3), e2020033, 2020

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topography of Ag/DLC composite synthesized by RF-PECVD. Surface Engineering, 36(7), 713-719. http://dx.doi.org/10.10 80/02670844.2019.1710937. 25. Le Bail, A., Duroy, H., & Fourquet, J. L. (1988). Ab-initio structure determination of LiSbWO6 by X-ray powder diffraction. Materials Research Bulletin, 23(3), 447-452. http:// dx.doi.org/10.1016/0025-5408(88)90019-0. 26. Carvajal, J. R. (2001). Recent developments of the program FULLPROF, in commission on powder diffraction (IUCr). Newsletter, 26, 12-19. 27. Stern, P. G., & Segerman, E. (1968). On the structure of polypropylene fibres. Polymer, 9, 471-477. http://dx.doi. org/10.1016/0032-3861(68)90057-8. 28. Khulbe, K. C., Feng, C., & Matsuura, T. (2008). synthetic polymeric membranes: characterization by atomic force microscopy. USA: Springer. http://dx.doi.org/10.1007/9783-540-73994-4. 29. Nosonovsky, M. (2010). Entropy in tribology: in the search for applications. Entropy (Basel, Switzerland), 12(6), 1345-1390. http://dx.doi.org/10.3390/e12061345. 30. Melo, R. H. C., & Conci, A. (2008). Succolarity: defining a method to calculate this fractal measure. In Proceedings of the 15th International Conference on Systems, Signals and Image Processing (pp. 291-294). USA: IEEE. http://dx.doi. org/10.1109/iwssip.2008.4604424. 31. Mandelbrot, B. (1983). The fractal geometry of nature. New York: W. H. Freeman and company. http://dx.doi.org/10.1119/1.13295. 32. Le Bail, A. (2004). Monte carlo indexing with McMaille. Powder Diffraction, 19(3), 249-254. http://dx.doi.org/10.1154/1.1763152. 33. Israelachvili, J. N. (2011). Intermolecular and surface forces. Cambridge: Academic Press. 34. Coutinho, F. M. B., Mello, I. L., & Santa Maria, L. C. (2003). Polietileno: principais tipos, propriedades e aplicações. Polímeros: Ciência e Tecnologia, 13(1), 1-13. http://dx.doi. org/10.1590/S0104-14282003000100005. 35. Ghanbarzadeh, B., & Oromiehi, A. R. (2008). Biodegradable biocomposite films based on whey protein and zein: barrier, mechanical properties and AFM analysis. International Journal of Biological Macromolecules, 43(2), 209-215. http://dx.doi. org/10.1016/j.ijbiomac.2008.05.006. PMid:18619671. 36. Bergo, P., Sobral, P. J. A., & Prison, J. M. J. (2010). Effect of Glycerol on Physical Properties of Cassava Starch Films. Journal of Food Processing and Preservation, 34(s2), 401-410. http://dx.doi.org/10.1111/j.1745-4549.2008.00282.x. 37. Almeida, A. P., Pinto, E. P., Santos, P. G. P., Filho, H. D. F., & Matos, R. S. (2019). Distribution of microorganisms on surface of kefir biofilms associated with açaí extract. Scientia Amazônia, 8(3), 10-18. 38. Kantorsk, K. Z., & Pagani, C. (2007). Influência da rugosidade superficial dos materiais odontológicos na adesão bacteriana: revisão de literatura. Revista Brasileira de Odontologia, 19(3), 325-330. http://dx.doi.org/10.1016/j.ijbiomac.2008.05.006. 39. Pinto, E. P., Tavares, W. S., Matos, R. S., Ferreira, A. M., Menezes, R. P., Costa, M. E. H. M., Souza, T. M., Ferreira, I. M., Sousa, F. F. O., & Zamora, R. (2018). Influence of low and high glycerol concentrations on wettability and flexibility of chitosan biofilms. Quimica Nova, 41(10), 1109-1116. http:// dx.doi.org/10.21577/0100-4042.20170287. 40. Dias, P. A., Rosa, J. V., Tejada, T. S., & Timm, C. D. (2016). Propriedades antimicrobianas do kefir. Arquivos do Instituto Biológico, 83(e0762013), 1-5. http://dx.doi.org/10.1590/18081657000762013. 41. Bosch, A., Golowczyc, M. A., Abraham, A. G., Garrote, G. L., De Antoni, G. L., & Yantorno, O. (2006). Rapid discrimination of lactobacilli isolated from kefir grains by FT-IR spectroscopy. 9/10

Matos, R. S., Gonçalves, E. C. M., Pinto, E. P., Lopes, G. A. C., Ferreira, N. S., & Resende, C. X. International Journal of Food Microbiology, 111(3), 280287. http://dx.doi.org/10.1016/j.ijfoodmicro.2006.05.010. PMid:16860422. 42. Talaro, K. P. E. (2002). Foundations in microbiology. New York: Mc Graw-Hill. 43. Domingues, F. N., Oliveira, M. D. S., Siqueira, G. R., Roth, A. P. T. P., Santos, J., & Mota, D. A. (2011). Estabilidade aeróbia, pH e dinâmica de desenvolvimento de microrganismos da cana-de-açúcar in natura hidrolisada com cal virgem. Revista Brasileira de Zootecnia, 40(4), 715-719. http://dx.doi. org/10.1590/S1516-35982011000400003.

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44. Ordóñez, J. A. (2005). Tecnologia de alimentos: componentes dos alimentos e processos. São Paulo: Artmed. 45. Ramírez Granillo, A., Canales, M. G., Espíndola, M. E., Martínez Rivera, M. A., de Lucio, V. M., & Tovar, A. V. (2015). Antibiosis interaction of Staphylococccus aureus on Aspergillus fumigatus assessed in vitro by mixed biofilm formation. BMC Microbiology, 15, 33. PMid:25880740. Received: Apr. 14, 2020 Revised: Sept. 07, 2020 Accepted: Oct. 12, 2020

Polímeros, 30(3), e2020033, 2020

ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.04220

Physicochemical characterization, drug release and mechanical analysis of ibuprofen-loaded uhmwpe for orthopedic applications Loise Silveira da Silva1 , Izabelle de Mello Gindri1 , Gean Vitor Salmoria1,2*  and Carlos Rodrigo de Mello Roesler1  1 Laboratório de Engenharia Biomecânica – LEBm, Hospital Universitário, Universidade Federal de Santa Catarina – UFSC, Florianópolis, SC, Brasil 2 Núcleo de Inovação em Moldagem e Manufatura Aditiva – NIMMA, Departamento de Engenharia Mecânica, Universidade Federal de Santa Catarina – UFSC, Florianópolis, SC, Brasil

*gean.salmoria@ufsc.br

Abstract In this study, the preparation of a novel functional material for orthopedic implants using compression molding was investigated. The new functional material is envisioned to avoid inflammatory reactions in vivo after prosthesis implantation. Ibuprofen-loaded UHMWPE samples were prepared in two concentrations (3% and 5%) and samples were characterized in terms of physicochemical and mechanical properties. In addition, the drug-release profile was investigated. The manufacturing process resulted in a homogeneous polymer matrix with homogeneous drug dispersion. The addition of ibuprofen had a minor effect on physicochemical properties but a more significant influence on the mechanical behavior of the specimens was observed. Drug release was demonstrated and overall the results obtained showed a positive outcome with regard to the intended use. The properties analyzed remained within an acceptable range for medical application and the drug-release profile obtained for the material developed shows promise for its use as an anti-inflammatory system. Keywords: UHMWPE, ibuprofen, biomaterial, drug delivery, orthopedic implants, material characterization. How to cite: Silva, L. S., Gindri, I. M., Salmoria, G. V., & Roesler, C. R. M. (2020). Physicochemical characterization, drug release and mechanical analysis of ibuprofen-loaded UHMWPE for orthopedic applications. Polímeros: Ciência e Tecnologia, 30(3), e202034. https://doi.org/10.1590/0104-1428.04220

1. Introduction Total knee arthroplasty (TKA) is a standard technique used to improve mobility and reduce pain in patients with osteoarthritis, and it has been in use for the past 50 years[1,2]. TKA is a procedure that replaces the knee joint with a set of tibial and femoral components composed of metal alloys with a polymer spacer as the articular surface[3]. Ultra-high molecular weight polyethylene (UHMWPE) has been used as the standard material to produce the insert implanted between the femoral and tibial components in TKA procedures, due to its excellent mechanical properties and biocompatibility[4]. Despite the attractive mechanical properties of UHMWPE, which account for its use to manufacture spacers, the constant movement and contact against the tribological pair lead to wear of the polyethylene component[1]. In addition to geometrical changes associated with the wear, which may affect the performance of the implant, wear particles are generated. This wear debris is known to trigger foreign body reactions that may lead to chronic inflammation and ultimately to osteolysis, future aseptic loosening and the

Polímeros, 30(3), e202034, 2020

need for revision surgery[5,6], which represents a negative aspect of UHMWPE and limits the life span of the implant[7]. When these particles are detected by the immunological system, they are phagocytosed by macrophages and multinucleated cells are formed, which lead to the activation of pro-inflammatory cytokines, such as TNF-α, IL-1β and IL-6. This process results in the proliferation and maturation of osteoclasts, multinucleated cells responsible for bone resorption. On the other hand, the proliferation of osteoblasts, bone forming cells, is reduced, generating an imbalance between osteogenesis and bone resorption[8]. One approach to inhibit the development of this inflammatory reaction is the use of anti-inflammatories or bone resorption suppressors as a post-operative treatment[9]. Recent studies have demonstrated the 10-year implant survival rate for TKR was 96.1%, and the 20-year implant survival rate was 89·7%. However, for younger and more active patients the lifetime risk of revision increased up to 35% with great differences between female and male patients (15% lower for women in same age group)[10]. Although TKA has high success rates, with more than a million procedures

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

Silva, L. S., Gindri, I. M., Salmoria, G. V., & Roesler, C. R. M. conducted every year[9], problems related to inflammatory reactions, such as osteolysis, are one of the main factors linked to the high number of revision surgeries[11]. According to AJRR (American Joint Replacement Registry)[12], in 2016, 3.4% of all knee replacement procedures performed in the United States were revision surgeries, and 13% of these were related to inflammatory reactions surrounding the implant area. Various alternatives related to the manufacturing process and material composition have been tested over the past few years to maintain the positive aspects of this material and, at the same time, improve its relationship with the human body and the wear properties. Crosslinked-UHMWPE has emerged as a promising material for total hip arthroplasty with some improved properties when compared with conventional UHMWPE. However, in TKA, this material has not demonstrated the same benefits. This is due to the multidirectional cyclic movements of the knee combined with a reduction in the mechanical properties and fatigue resistance observed for crosslinked-UHMWPE[13]. Thus, conventional UHMWPE has shown better performance in TKA procedures and remains, along with Vitamin E-loaded UHMWPE, the preferred choice of material. UHMWPE with the antioxidant Vitamin E incorporated was developed as an alternative to reduce the oxidative reaction that occurs during in vivo use. In the material matrix this antioxidant reacts with free radicals and retards the oxidative process[14]. UHMWPE with Vitamin E has been used in arthroplasty procedures and has shown positive results in terms of wear resistance and particle interaction with the immunological system[4]. In order to reduce the number of revision procedures by reducing inflammatory reactions and infections, researchers all around the world have been testing the incorporation of different drugs, such as vancomycin[15], gentamicin[5], bupivacaine[16] and bisphosphonates[17], to avoid infection and inhibit bone resorption. Although there have been some good outcomes, this kind of technology is still in the pre‑clinical testing phase as it deals with humans, and therefore has not yet been used in arthroplasty procedures. Considering the constant need to improve the performance of UHMWPE in TKA, this paper proposes the formulation of a novel biomaterial based on the incorporation of ibuprofen (IBU) into the UHMWPE polymeric matrix. IBU has been used as a post-operative treatment in arthroplasty procedures, as a treatment for local pain and inflammations associated with arthritis and musculoskeletal problems[18] and incorporated in different polymers to reach and treat specific areas of the human body[19]. The IBU-loaded UHMWPE prepared in this study is aimed at reducing the inflammatory reaction after implantation, which would help to reduce the number of revision surgeries caused by periprosthetic loosening due to osteolysis. IBU-UHMWPE specimens were prepared by compression molding and then characterized in terms of their physicochemical properties, drug release profile and mechanical properties. The results and data obtained were used to discuss the potential of the material developed, to analyze the effect of IBU on the physicochemical and mechanical properties, and to investigate the drug-release profile. 2/8

2. Experimental 2.1 Materials UHMWPE resin (GUR 1020, Ticona) was the main material used to prepare the specimens, along with IBU (Viafarma, Joinville, Brazil). As a way to compare the material properties and behavior, specimens with drug concentrations of 3wt% and 5wt% were prepared as well as specimens without the addition of the anti-inflammatory drug. The specimens were named UHMWPE, UHMWPE 3% IBU and UHMWPE 5% IBU.

2.2 Preparation of specimens Specimens were prepared according to the methodology proposed by Suhardi et al.[15] with some modifications. The compression molding process was carried out in a hydraulic press with a load capacity of 15 tons (Bonevau, Rio do Sul, Brazil). Samples of UHMWPE 3% IBU (0.3 g IBU: 9.7 g of UHMWOE) and UHMWPE 5% IBU (0.5 g IBU: 9.5 g of UHMWPE) were previously prepared by mixing the polymer and the drug for 10 min. Each formulation was placed in the mold and kept at 150 °C and 5 MPa for 15 min to reach total material consolidation. After this period, the specimens were cooled to room temperature. All test specimens were prepared according to standard specifications as will be detailed in the following sections.

2.3 Specimen characterization 2.3.1 Physicochemical analysis Differential scanning calorimetry (DSC) was conducted on all specimens to determine the thermal parameters, such as melting point temperature (Tm), melting enthalpy and the degree of crystallinity according to ASTM F2625-10. Three specimens, with masses of between 0.003 g and 0.008 g, were extracted from each formulation proposed and sealed in an appropriate aluminum pan for posterior analysis in DSC equipment (PerkinElmer 600, São Paulo, Brazil). Tests were run with a heating/cooling rate of 10 °C/min and the crystallinity degree was obtained by integration of the endotherm peak between 50 °C and 160 °C. Thin films were prepared for Fourier-transform infrared spectroscopy and microscopy experiments. The films were obtained from different regions of specimens to address the material uniformity. Specimens of 1 cm x 1 cm x 1 cm were removed from each specimen and embedded in paraffin. Slices of 200 µm and 90 µm were obtained from paraffin embedded specimens with the aid of a microtome for Fourier infrared (FTIR) spectroscopy and microscopy analysis, respectively. The FTIR spectroscopy was performed in the transmission mode and according to ASTM 2102-13. Optical microscopy (Nikon E-200) was performed to evaluate the quality of the consolidation process. The films were evaluated under 100 X magnification. Lastly, density measurements of three specimens of each formulation were taken following Archimedes’ principle. The specimens were prepared with at least 1 mm of thickness for each 1 g of material with a maximum of 5 g, as recommended by the technical standard ASTM D792-13. Polímeros, 30(3), e202034, 2020

Physicochemical characterization, drug release and mechanical analysis of ibuprofen-loaded uhmwpe for orthopedic applications 2.3.2 Drug release test

3. Results and Discussion

Spectrophotometric analysis was conducted using a UV-Vis spectrophotometer (Model UV-5200, Global Trade Technology, Monte Alto, Brazil) to obtain the release profile and investigate the way in which the drug would be eluted from the UHMWPE when implanted in the human body. IBU, in concentrations of (mg/ml) 0.0652, 0.125, 0.25, 0.5, 1.0, 2.0, 2.5, 3.0 and 4.0, was dissolved in phosphate buffer solution (PBS, pH = 7.4) and then analyzed in the spectrophotometer at a wavelength (λmax) of 264 nm to prepare the IBU calibration curve (Figure S1, Supporting Information, Supplementary Material). The UHMWPE specimens with IBU were then immersed in 3 mL of PBS and placed in a Dubnoff bath at a temperature of 37.0 ± 0.5 °C. The entire PBS volume was collected from each specimen after 6 h and 24 h, then every 24 h for one week and subsequently twice a week until the 30th day of the drug release test (Figure S2). After each PBS collection, the amount taken for the test was replaced with fresh PBS. Due to the nature of the system developed, the models chosen to analyze the data were the zero-order model, Higuchi model and Korsmeyer-Peppas model (also known as the power law model)[20]. The curves resulting from this analysis are provided in Figures S3 – S5.

3.1 Physicochemical analysis

2.3.3 Mechanical testing Mechanical testing was performed according to technical standard ASTM D638 with some modifications. These experiments were carried out on a universal testing machine (DL 3000, EMIC, São José dos Pinhais, Brazil) operating with a load cell of 500 N and at a test speed of 1 mm/min, with the use of an Instron extensometer (Cat. No. 2630-107, Instron, Norwood, USA) with a gauge length of 25 mm and travel length of +25/-2.5 mm. The length between grips during the tests was 33.9 mm. Three specimens of each drug concentration were machined into a dumbbell-shaped format, 2 mm thick and with a cross‑section area of 6 mm2, to obtain the elastic modulus and tensile yield strength of the specimens for posterior analysis of the effect of the IBU on the mechanical properties. The elastic modulus specimens were obtained in the linear strain region of 0.0005 to 0.0025 mm/mm.

The consolidation process was investigated using optical microscopy. Images obtained for UHMWPE, UHMWPE 3% IBU and UHMWPE 5% IBU are shown in Figure 1. Figure 1a shows a UHMWPE specimen, on which only cut marks are observed, indicating a satisfactory consolidation procedure. The presence of the drug in the polymeric matrix is confirmed for both UHMWPE 3% IBU and UHMWPE 5% IBU, as demonstrated in Figures 1b and 1c, respectively, in which IBU crystals can be observed. The DSC and density results are summarized in Table 1. The data demonstrate that the addition of IBU did not affect the melting point temperature, onset temperature, crystallinity percentage and polymer density (p >0.05). The reference density values for medical grade UHMWPE are in the range of 927 - 944 kg/m3[21]. Also, according to the literature, crystallinity values should be between 50 and 55%[4] and the melting point temperature in the range of 125 - 138 °C[22]. The results of our analysis exhibit a positive outcome, showing values in agreement with the reference data provided in the literature and in the technical standard specification ASTM F648 regarding UHMWPE for medical applications[21]. The material composition was confirmed through FTIR analysis. A spectrum for each composition is shown in Figure 2, along with the IBU spectrum to locate and compare the IBU peaks in the UHMWPE 3% IBU and UHMWPE 5% IBU specimens. The IBU, UHMWPE, UHMWPE 3% IBU and UHMWPE 5% IBU spectra are shown in Figures 2a-d. The peak at around 1720 cm-1 in the IBU, UHMWPE 3% IBU and UHMWPE 5% IBU spectra corresponds to the carbonyl group present in the drug composition. The detection of this peak in the UHMWPE 3% IBU and UHMWPE 5% IBU specimens confirms the incorporation of the drug in the polymeric matrix. Characteristics of the UHMWPE bands include: symmetric and asymmetric stretching of C-H in the CH2 groups located at 2950 cm-1 and 2850 cm-1, respectively; and vibration bands in the range of 1350 cm-1 and 1450 cm-1

Figure 1. Microscopic images taken with a magnification of 200X: (a) UHMWPE, (b) UHMWPE 3% IBU and (c) UHMWPE 5% IBU. Table 1. Physicochemical properties. UHMWPE UHMWPE 3% IBU UHMWPE 5% IBU

Tm [°C]

Tonset [°C]

Xc [%]

Density [kg/m3]

131.94 ± 0.98 131.02 ± 0.55 131.06 ± 0.50

122.47 ± 0.94 122.34 ± 0.32 121.72 ± 0.27

53.12 ± 1.70 53.69 ± 3.97 52.58 ± 0.37

927.30 ± 3.70 930.43 ± 7.22 928.28 ± 3.48

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Silva, L. S., Gindri, I. M., Salmoria, G. V., & Roesler, C. R. M.

Figure 2. FTIR spectra for (a) IBU, (b) UHMWPE, (c) UHMWPE 3% IBU and (d) UHMWPE 5% IBU.

associated with the CH2 bending movement and at 720 cm-1 with the CH2 rocking vibration. For the IBU, bands between 3300 cm-1 and 2500 cm-1 correspond to the stretching of OH in the carboxylic acid groups and the vibration of C-H in CH3 and in the phenyl functional groups. As previously mentioned, a characteristic band of IBU, with a high intensity, is located at 1720 cm-1 representing the stretching of the carbonyl bond C=O[23]. 4/8

3.2 Drug release test The two UHMWPE formulations with IBU are classified as a possible drug-release system by diffusion, where the drug is eluted from the polymeric matrix over a prolonged time frame. The data obtained from the drug release tests were analyzed by mathematical methods developed from the First Law of Fick[24,25]. The IBU calibration curve used to analyze Polímeros, 30(3), e202034, 2020

Physicochemical characterization, drug release and mechanical analysis of ibuprofen-loaded uhmwpe for orthopedic applications the data obtained by spectrophotometry analysis is shown in Figure S1 in the Supporting Information. The coefficients used to calculate the amount of IBU released at each interval of analysis were extracted from this curve, leading to the total amount released by the system at the end of the test. The results obtained over the 30-day period of IBU release are shown in Figure 3 (%) and Figure S2 (mg/ml). The results for the percentage of drug released show that of the total drug theoretically incorporated in the specimens, UHMWPE 3% IBU eluted approximately 7.8% and UHMWPE 5% IBU released 9.7% within the first 7 days of the tests. The UHMWPE 3% IBU specimen released 1.59 mg in the first week while UHMWPE 5% IBU released 2.90 mg/ml. In the second week, around 12.00% of the total was eluted from UHMWPE 3% IBU and 14.79% from UHMWPE 5% IBU, representing 2.26 mg and 4.42 mg, respectively. After 30 days, UHMWPE 3% IBU had released 3.81 mg and UHMWPE 5% IBU 7.05 mg, representing 19.62% and 23.91%, respectively. These data demonstrate that the IBU release is more accentuated at the beginning of the eluting period, followed by a more controlled and smooth release over time. The total amounts theoretically incorporated in each specimen were 20.39 mg for UHMWPE 3% IBU and 29.95 mg for UHMWPE 5% IBU. The kinetic release profiles (Figure S3-S5) for the materials developed showed that the data for the UHMWPE 5% IBU specimen best fitted the Higuchi’s Model (with a correlation coefficient of 0.9942), while for the UHMWPE 3% IBU specimen the best fit was obtained with the Korsmeyer‑Peppas model (with a correlation factor of 0.9949). Although the UHMWPE 3% IBU release coefficient indicated a non‑Fickian diffusion release[24,25], with a release exponent of 0.6091, the correlation coefficient obtained with Higuchi’s Model was also higher than 0.99, which suggests that the UHMWPE 3% IBU release may also be governed by Fickian diffusion. An overview of the results obtained from the kinetic analysis is given in Table 2.

The kinetic release coefficients obtained for the Higuchi model confirm that the IBU release is governed by diffusion, with values of less than 0.5, indicating transport governed by Fick’s Law[20,24]. In addition, considering Higuchi’s Model, the kinetic release coefficient of UHMWPE 3% IBU was 0.013, which is smaller than the value for UHMWPE 5% IBU of 0.015. Therefore, the higher concentration of IBU in the polymer seems to facilitate the diffusion of the drug, as a greater portion of it may be in contact with the polymer matrix. The application of mathematical models indicated that the UHMWPE/IBU formulation as a system is governed by Higuchi’s Model, especially the formulation with the higher IBU concentration. This suggests that at the higher concentration a greater portion of the anti-inflammatory is in contact with the matrix, which would facilitate its diffusion process. The results demonstrate a gradual drug release over time, which is of interest for polymeric systems intended for anti-inflammatory applications[23], such as the prevention of osteolysis. Furthermore, the drug release occurred over a longer period compared with other systems, such as bupivacaine[16], vancomycin[15] and gentamicin‑loaded UHMWPE[5], in which antibiotic elution occurred over 5, 12 and 25 days, respectively. While antibiotic systems are designed to achieve release in short periods after implantation to avoid infections, an anti-inflammatory is expected to act for a longer period, to mitigate capsular formation after implantation as well as osteolysis due to particle generation.

3.3 Mechanical testing The tensile curves obtained from the tests are shown in Figures 4a, 4b and 4c for UHMWPE, UHMWPE 3% IBU and UHMWPE 5% IBU, respectively. The three formulations tested showed similar mechanical behavior and the error between specimens within the same group was small. The differences between the material compositions demonstrate a decrease in their elastic modulus (p<0.05), as expected

Figure 3. Drug release in percentage (%) over a 30-day period. Table 2. Kinetic analysis results. SPECIMEN KINETIC MODEL ZERO-ORDER MODEL HIGUCHI MODEL KORSMEYER-PEPPAS MODEL

K 0.002 0.013 K 0.008

UHMWPE 3% IBU r 0.989 0.994 r 0.996

n 0.624

K 0.003 0.015 K 0.013

UHMWPE 5% IBU r 0.992 0.994 r 0.998

n 0.530

K = release kinetic coefficient; r = correlation coefficient; n = release exponent

Polímeros, 30(3), e202034, 2020

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Silva, L. S., Gindri, I. M., Salmoria, G. V., & Roesler, C. R. M.

Figure 4. Engineering stress - strain curves obtained from tensile tests (CP = specimen).

Figure 5. Comparison of elastic modulus and tensile yield strength results for the different formulations tested. Table 3. Elastic modulus and tensile yield strength results. E [MPa] UHMWPE UHMWPE 3% IBU UHMWPE 5% IBU

825.89 ± 3.62 696.12 ± 28.92 661.45 ± 22.61

Tensile yield strength [MPa] 16.02 ± 0.44 15.89 ± 0.21 15.50 ± 0.14

with the addition of a doping agent in the polymeric matrix. Results are summarized in Table 3. The results in Table 3 and Figure 5 indicate that the addition of 3% of IBU to the polymeric matrix decreased the elastic modulus by 16%. However, despite this decrease, the value remained between 500 MPa and 800 MPa, that is, within the range recommended for the use of UHWMPE as a biomaterial[22]. The 5% formulation provided a similar result, with a decrease of up to 20% in relation to the UHMWPE without the drug. These results suggest that the addition of IBU facilitates the molecular chain mobility and increases the ductility of the material. No significant difference was noted for the tensile yield strength when comparing UHMWPE, UHMWPE 3% IBU and UHMWPE 5% IBU (p>0.05), which means that the addition of the drug did not have a notable impact on this parameter. The same behavior has been observed for alendronate and vancomicyn-loaded UHMWPE, with drug contents of less than or equal to 5%.

4. Conclusions The results reported herein demonstrate the incorporation of IBU, an anti-inflammatory drug, into UHMWPE for the first time. The consolidation process selected produced specimens with a homogeneous matrix, total polymer fusion and a good IBU dispersion, as demonstrated by the microscopy analysis. The addition of IBU resulted in minor 6/8

effects on the crystallinity, melting point temperature, onset temperature and density. The drug release experiments demonstrated controlled and sustained drug release. Specimens with higher drug content resulted in greater percentages of release. The mechanical tests showed the influence of IBU on the mechanical properties of UHMWPE. The modulus of elasticity values for the two IBU formulations (0.0005 and 0.0025 mm/mm) were up to 20% lower when compared with the UHMWPE without the drug. Despite the impact on the mechanical properties, the results obtained were satisfactory, demonstrating that the addition of IBU had only a slight impact on the quasi-static mechanical properties of the UHMWPE. The wear behavior and cyclic load effects need to be evaluated in future work. Overall, the specimens of UHMWPE with IBU incorporated showed positive results with regard to their use in arthroplasty procedures. The characteristics and properties are consistent with those recommended for biomedical applications and the release profiles are promising for an anti-inflammatory system. Further tests and improvements must be carried out to achieve a better combination of parameters and obtain a raw material suitable for knee implants, aimed at reducing the number of inflammatory reactions and, consequently, the number of revision surgeries.

5. Acknowledgements The authors would like to thank Coordenação de Aperfeiçoamento de Pessoal de Nível Superior CAPES (I.M.G and L.S.S), and Conselho Nacional de Desenvolvimento Científico e Tecnológico - CNPq (G.V.S and C.R.M.R) for financial support.

6. References 1. Kurtz, S. M., & Kurtz, S. M. (2004). Chapter 8 – The clinical performance of UHMWPE in knee replacements. In S. M. Polímeros, 30(3), e202034, 2020

Physicochemical characterization, drug release and mechanical analysis of ibuprofen-loaded uhmwpe for orthopedic applications Kurtz (Ed.), The UHMWPE handbook: ultra-high molecular weight polyethylene in total joint replacement (pp. 151-188). Cambridge: Academic Press. https://doi.org/10.1016/B978012429851-4/50009-7. 2. Kurtz, S. M., & Kurtz, S. M. (2004). Chapter 7 – The origins and adaptations of UHMWPE for knee replacements. In S. M. Kurtz (Ed.), The UHMWPE handbook: ultra-high molecular weight polyethylene in total joint replacement (pp. 123-150). Cambridge: Academic Press. https://doi.org/10.1016/B978012429851-4/50008-5. 3. Carr, B. C., & Goswami, T. (2009). Knee implants - review of models and biomechanics. Materials & Design, 30(2), 398-413. http://dx.doi.org/10.1016/j.matdes.2008.03.032. 4. Bracco, P., Bellare, A., Bistolfi, A., & Affatato, S. (2017). Ultra-high molecular weight polyethylene: influence of the chemical, physical and mechanical properties on the wear behavior. A Review. Materials (Basel), 10(7), 791. http:// dx.doi.org/10.3390/ma10070791. PMid:28773153. 5. Manoj Kumar, R., Gupta, P., Sharma, S. K., Mittal, A., Shekhar, M., Kumar, V., Manoj Kumar, B. V., Roy, P., & Lahiri, D. (2017). Sustained drug release from surface modified UHMWPE for acetabular cup lining in total hip implant. Materials Science and Engineering C, 77, 649-661. http://dx.doi.org/10.1016/j. msec.2017.03.221. PMid:28532076. 6. Nabeshima, A., Pajarinen, J., Lin, T., Jiang, X., Gibon, E., Córdova, L. A., Loi, F., Lu, L., Jämsen, E., Egashira, K., Yang, F., Yao, Z., & Goodman, S. B. (2017). Mutant CCL2 protein coating mitigates wear particle-induced bone loss in a murine continuous polyethylene infusion model. Biomaterials, 117, 1-9. http://dx.doi.org/10.1016/j.biomaterials.2016.11.039. PMid:27918885. 7. Topolovec, M., Cör, A., & Milošev, I. (2014). Metal-on-metal vs. metal-on-polyethylene total hip arthroplasty tribological evaluation of retrieved components and periprosthetic tissue. Journal of the Mechanical Behavior of Biomedical Materials, 34, 243-252. http://dx.doi.org/10.1016/j.jmbbm.2014.02.018. PMid:24608233. 8. Steinbeck, M. J., & Veruva, S. Y. (2016). Pathophysiologic reactions to UHMWPE wear particles. In In S. M. Kurtz (Ed.), UHMWPE biomaterials handbook ultra high molecular weight polyethylene in total joint replacement and medical devices (pp. 506-530). USA: Elsevier. http://dx.doi.org/10.1016/B978-0323-35401-1.00028-4 9. Purdue, P. E., Koulouvaris, P., Potter, H. G., Nestor, B. J., & Sculco, T. P. (2007). The cellular and molecular biology of periprosthetic osteolysis. Clinical Orthopaedics and Related Research, 454, 251-261. http://dx.doi.org/10.1097/01. blo.0000238813.95035.1b. PMid:16980902. 10. Bayliss, L. E., Culliford, D., Monk, A. P., Glyn-Jones, S., PrietoAlhambra, D., Judge, A., Cooper, C., Carr, A. J., Arden, N. K., Beard, D. J., & Price, A. J. (2017). The effect of patient age at intervention on risk of implant revision after total replacement of the hip or knee: a population-based cohort study. Lancet, 389(10077), 1424-1430. http://dx.doi.org/10.1016/S01406736(17)30059-4. PMid:28209371. 11. Huang, Y.-F., Zhang, Z.-C., Xu, J.-Z., Xu, L., Zhong, G.-J., He, B.-X., & Li, Z.-M. (2016). Simultaneously improving wear resistance and mechanical performance of ultrahigh molecular weight polyethylene via cross-linking and structural manipulation. Polymer, 90, 222-231. http://dx.doi.org/10.1016/j. polymer.2016.03.011. 12. Berry, D. J., Bozic, K. J., & Lewallen, D. G. (2016). AJRR Annual Report 2016 (pp. 46). Rosemont, IL: American Academy of Orthopaedic Surgeons (AAOS). 13. Pruitt, L. A. (2005). Deformation, yielding, fracture and fatigue behavior of conventional and highly cross-linked ultra high Polímeros, 30(3), e202034, 2020

molecular weight polyethylene. Biomaterials, 26(8), 905915. http://dx.doi.org/10.1016/j.biomaterials.2004.03.022. PMid:15353202. 14. Wernle, J. D., Mimnaugh, K. D., Rufner, A. S., Popoola, O. O., Argenson, J.-N., & Kelly, M. (2017). Grafted vitamin-E UHMWPE may increase the durability of posterior stabilized and constrained condylar total knee replacements. Journal of Biomedical Materials Research. Part B, Applied Biomaterials, 105(7), 1789-1798. http://dx.doi.org/10.1002/jbm.b.33710. PMid:27192378. 15. Suhardi, V. J., Bichara, D. A., Kwok, S., Freiberg, A. A., Rubash, H., Malchau, H., Yun, S. H., Muratoglu, O. K., & Oral, E. (2017). A fully functional drug-eluting joint implant. Nature Biomedical Engineering, 1(6), 1-21. http://dx.doi.org/10.1038/ s41551-017-0080. PMid:29354321. 16. Gil, D., Grindy, S., Muratoglu, O., Bedair, H., & Oral, E. (2019). Antimicrobial effect of anesthetic-eluting ultra-high molecular weight polyethylene for post-arthroplasty antibacterial prophylaxis. Journal of Orthopaedic Research, 37(4), 981-990. http://dx.doi.org/10.1002/jor.24243. PMid:30737817. 17. Yang, D., Qu, S., Huang, J., Cai, Z., & Zhou, Z. (2012). Characterization of alendronate sodium-loaded UHMWPE for anti-osteolysis in orthopedic applications. Materials Science and Engineering C, 32(2), 83-91. http://dx.doi.org/10.1016/j. msec.2011.09.012. 18. Celebi, D., Guy, R. H., Edler, K. J., & Scott, J. L. (2016). Ibuprofen delivery into and through the skin from novel oxidized cellulose-based gels and conventional topical formulations. International Journal of Pharmaceutics, 514(1), 238-243. http:// dx.doi.org/10.1016/j.ijpharm.2016.09.028. PMid:27863667. 19. Salmoria, G. V., Paggi, R. A., Castro, F., Roesler, C. R. M., Moterle, D., & Kanis, L. A. (2016). Development of PCL/ ibuprofen tubes for peripheral nerve regeneration. Procedia CIRP, 49, 193-198. http://dx.doi.org/10.1016/j.procir.2015.11.014. 20. Klauss P. (2010). Desenvolvimento de dispositivos poliméricos implantáveis para a liberação de fármaco fabricados por sinterização seletiva a laser [Tese de doutorado]. Programa de Pós-graduação em Ciência e Engenharia dos Materiais, Centro Tecnológico, Universidade Federal de Santa Catarina, Florianópolis. 21. ASTM International. (1980). F648-80: standard specification for ultra-high-molecular-weight polyethylene powder and fabricated form for surgical implants. West Conshohocken: ASTM International. 22. Kurtz, S. M. (2004). The UHMWPE handbook: ultra-high molecular weight polyethylene in total joint replacement. Cambridge: Academic Press. 23. Vieira, E. S., Salmoria, G. V., de Mello Gindri, I., & Kanis, L. A. (2018). Preparation of ibuprofen-loaded HDPE tubular devices for application as urinary catheters. Journal of Applied Polymer Science, 135(2), 1-8. http://dx.doi.org/10.1002/ app.45661. 24. Manadas, R., Pina, M. E., & Veiga, F. (2002). A dissolução in vitro na previsão da absorção oral de fármacos em formas farmacêuticas de liberação modificada. Revista Brasileira de Ciência do Solo, 38(4). http://dx.doi.org/10.1590/S151693322002000400002. 25. Costa, P. J. C. (2002). Avaliação in vitro da lioequivalência de formulações farmacêuticas. Revista Brasileira de Ciências Farmacêuticas, 38(2), 141-153. http://dx.doi.org/10.1590/ S1516-93322002000200003. Received: May 09, 2020 Revised: Sept. 15, 2020 Accepted: Oct. 12, 2020 7/8

Silva, L. S., Gindri, I. M., Salmoria, G. V., & Roesler, C. R. M.

Supplementary Material Supplementary material accompanies this paper. Figure S1. IBU calibration curve. Figure S2.Drug release in mg/day in 30 days. Figure S3. Zero order model fit. Figure S4. Higushi model fit. Figure S5. Korsmeyer-peppas model fit. This material is available as part of the online article from http://www.scielo.br/polimeros

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Polímeros, 30(3), e202034, 2020

ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.05620

Comparison of MA-g-PP effectiveness through mechanical performance of functionalised graphene reinforced polypropylene Saravanan Natarajan1 , Rajasekar Rathanasamy1* , Sathish Kumar Palaniappan2 , Suresh Velayudham3 , Hari Bodipatti Subburamamurthy1  and Kaushik Pal4  Department of Mechanical Engineering, Kongu Engineering College, Erode, Tamil Nadu, India Department of Mining Engineering, Indian Institute of Technology Kharagpur, West Bengal, India 3 Department of Mechanical Engineering, Sona College of Technology, Salem, Tamil Nadu, India 4 Department of Mechanical & Industrial Engineering, Indian Institute of Technology Roorkee, Uttarakhand, India 1

*rajasekar.cr@gmail.com

Abstract This work aims in developing carboxyl functionalised graphene based PP nanocomposites by using melt mixing method to enhance the mechanical and thermal properties. Maleic anhydride grafted polypropylene was used as a compatibilizer to achieve better compatibility between the non-polar polymer and polar nanofiller. FTIR study confirms the presence of functional groups at corresponding absorption levels. TEM and SEM image shows the uniform distribution of COOH-Gr onto the PP matrix with the addition of MA-g-PP onto it. The tensile strength and young’s modulus of PMG5 depicts better improvement of 62% and 20% compared to neat sample. The increase in storage modulus of 19.02% was obtained for PG and 43.48% for PMG samples. The reduction in tan δ peak confirms the minimum heat buildup and as a result, leads to better damping characteristics of the nanofiller incorporated PP matrix. Keywords: functionalised graphene, maleic anhydride grafted polypropylene, polypropylene, scanning electron microscopy, transmission electron microscopy. How to cite: Natarajan, S., Rathanasamy, R., Palaniappan, S. K., Velayudham, S., Subburamamurthy, H. B., & Pal, K. (2020). Comparison of MA-g-PP effectiveness through mechanical performance of functionalised graphene reinforced polypropylene. Polímeros: Ciência e Tecnologia, 30(3), e2020035. https://doi.org/10.1590/0104-1428.05620

1. Introduction Polymer nanocomposites are the material wherein nanofillers are reinforced in a polymer matrix. Polypropylene (PP) is a thermoplastic material used in several applications like reusable containers of various types and stationery plastic parts, automotive components, laboratory equipments and polymer banknote [1]. PP is a breed of polymer which transforms its phase into liquid when it is heated and when freezing, it turns into a glassy state. Recently researches are being carried out with nanofillers such as exfoliated clay, alumina, silica, graphene, carbon nanofibers, carbon nanotubes, nanocrystalline metals and also nanoscale inorganic fillers or modified fibers [2]. Nanofillers enable better dispersal of the fillers into the matrix, outcome as enhancement of properties of base matrix [3,4]. Superior interaction of nanofillers with other particles and improvement in properties is because of high surface area of the nanofillers [1]. Another effect of the nanofillers is the change in glass transition temperature (Tg). Both increase and decrease in Tg have been noticed depends on the interaction between the polymer and nanofiller. Some of the nanofillers used are carbon black, nano clay, carbon nanotubes, graphite oxide, graphene, copper nanoparticles

Polímeros, 30(3), e2020035, 2020

etc. Nanoparticles show improvement in properties at very low loadings. Thus it leads to significant reduction in weight for the same strength, performance and properties. The present report emphasis on the use of graphene, a carbon based nanofiller. A distinct characteristics of carbon is the wide range of forms that it can assume when two or more atoms bond. Thus the carbon has attracted and continued in the direction to attract, a considerable research and development interest from researchers of all parts of the world [4,5]. The carbon usage in nanotechnology is quite promising research area and considerable funding was invested by the government to carbon nanotechnology research. Pure PP without addition of any reinforcements is not suitable for potential applications. Some of the undesirable properties such as stiffness and rigidity may be enhanced by addition of suitable reinforcements. Though PP finds wide range of applications, its strength is not acceptable. The focus lies on enhancing its strength characteristics. Unfortunately, addition of reinforcements may not have superior adhesive

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

Natarajan, S., & Rathanasamy, R., & Palaniappan, S. K., & Velayudham, S., & Subburamamurthy, H. B., & Pal, K. property. This research focuses at increasing the strength and adhesive characteristics of PP in order to extent its usage [5]. PP nanocomposites have been synthesized with a variety of nanocomposite like modified calcium carbonate (CaCO3), nanoclay, layered silicate and multi walled carbon nanotubes (MWCNT). Derivatives of graphite have also been used as nanofillers. An effort has been made to use COOH functionalised graphene (COOH-Gr) as the nanofiller in turn to enhance the thermal property of the PP nanocomposite. Modification of graphene by adding carboxyl or hydroxyl groups on outer periphery of the graphene particles has found to be increases its interaction with polymer matrix [6-8] . Functional groups added may be carboxyl, amine or epoxy groups. COOH-Gr is highly polar and PP is nonpolar, mismatch occurs during nanocomposite production. Maleic anhydride grafted polypropylene (MA-g-PP) is the compatibilizer used to further increase the interaction among the fillers and the matrix [9,10]. This is possible because of the polar to polar interaction between the compatibilizer and COOH-Gr. Polymer nanocomposite with graphite oxide derivative graphene material as fillers have made known dramatic improvements in its tensile strength, elastic modulus, thermal behavior and electrical conductivity of the nanocomposite. Enhancements at low loadings have been observed due to large interfacial area and high aspect ratio. This work aims to prepare PP/functionalised graphene nanocomposite and compare the reinforcing result of COOH-Gr on the mechanical and thermal behavior of PP nanocomposite. This study also explains the morphology of prepared nanocomposite and to review the distribution of nanofillers in polymer matrix.

2. Materials and methods 2.1 Materials Commercial grade of PP was provided by Reliance Industries Limited, Mumbai, India. MA-g-PP with

8-10 wt.-% maleic anhydride content was obtained from SIGMA-ALDRICH and the chemical formula for the same is (C3H6)m (C7H8O3)n. COOH-functionalised graphene with carboxyl content of 20% was supplied by United Nanotech Innovations Pvt. Ltd., Bangalore.

2.2 Methods 2.2.1 Melt mixing The preparation of polymeric nanocomposite was carried out using melt mixing method [9,11]. The mixing operations were done at 160oC for a time period of 15 minutes. The speed is constantly maintained 50 rpm for throughout the mixing process. Initially the PP granules were feed in between the two rolls and allowed to mix for 5 minutes, Compatibilizer (MA-g-PP) is added to semisolid state PP content and mixed together for 3 more minutes. Functionalized graphene was poured in between the rolls once after the polymer gets completely melted. The whole blend is mixed for further 7 minutes to attain homogeneous mixing. 2.2.2 Mixing formulation The mixing formulation for preparing the PP nanocomposite is shown in Table 1. “PP” denotes pure polypropylene, “PM” denotes PP with MA-g-PP, “PG” denotes PP with COOHGr and “PMG” denotes PP with COOH-Gr and MA-g-PP. 2.2.3 Compression molding Compression molding is one of the plastic forming process in which a plastic material is placed directly in metal mold and kept in between two hot plates. Polymers softened by the applied heat and pressure to get the shape of the mold. The temperature of 160 oC and pressure of 0.8 MPa are maintained upto the molding polymer gets cured. Tensile and DMA tests specimens were made from compression molding process. The DMA specimens dimension were 6.4 cm length x 1.27 cm width x 0.32 cm thick.

Table 1. Samples mixing formulation. Sl. No. 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17. 18. 19.

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Sample code PP PM1 PM2 PM3 PM4 PM5 PM6 PG1 PG2 PG3 PG4 PG5 PG6 PMG1 PMG2 PMG3 PMG4 PMG5 PMG6

PP (wt. - %) 100 97.5 95 92.5 90 87.5 85 100 100 100 100 100 100 97.5 95 92.5 90 87.5 85

MA-g-PP (wt.-%) 2.5 5 7.5 10 12.5 15 2.5 5 7.5 10 12.5 15

COOH - Gr (wt.-%) 0.5 1 1.5 2 2.5 3 0.5 1 1.5 2 2.5 3

Polímeros, 30(3), e2020035, 2020

Comparison of MA-g-PP effectiveness through mechanical performance of functionalised graphene reinforced polypropylene 2.2.4 Fourier Transform Infrared Spectroscopy (FTIR) The Fourier transform infrared spectroscopy (FTIR) spectrum of graphene and modified graphene has been documented using a NEXUS 870 FTIR (Thermo Nicolet) equipment. The samples were prepared by making palate with KBr for characterization [12,13]. The signal resolution was 2 cm-1 and a minimum of 64 scans were within the range of 4000-500 cm-1. 2.2.5 Transmission electron microscopy (TEM) Transmission electron microscopy (TEM) imaging was carried out using JEM-2100 (JEOL JEM Co.) electron microscope. Dispersion morphology of composite samples are microtomed as very thin slices of 80 nm approximately by means of ultra microtome with a diamond knife and the temperature of samples were retain at -80 °C using liquid nitrogen. 2.2.6 Mechanical properties Dumb-bell shaped samples were made from compression molding process. The investigations were done under ambient condition with a universal testing machine (Hounsfield H10KS), the specimens were prepared for ASTM D 41206 and ASTM D 624-00(2007) standards. Tensile strength and Modulus of elasticity were measured at ambient temperature. The primary length of the specimens was 40 mm and 50 mm/ min speed was maintained for jaw separation. 2.2.7 Dynamic Mechanical Analysis (DMA) The TA Instruments DMA 2980 model instrument is used to characterize the materials mechanical behavior at different temperatures. The dynamic properties of the samples with dimensions 35 mm length, 12 mm width and 3 mm thick was studied at constant frequency mode with 1 Hz frequency [14]. The samples were heated at a rate of 3oC/min to the temperature range of 30 oC to 130 oC. 2.2.8 Scanning electron microscopy (SEM)

Each functional group has corresponding absorption levels which are identified from the FTIR results.

3.2 TEM TEM image of 2 wt.-% of functionalised graphene dispersed in PP matrix is revealed in Figure 2. The agglomeration of COOH-Gr in the matrix is observed in Figure 2. However, Figure 3 illustrates the TEM image of polymer matrix with 2 wt.-% of functionalised graphene and the addition of 10 wt.-% compatibilizer onto it. It proves the successful exfoliation of single layer graphene which gets uniformly distributed in the polymer matrix [9,18], this is because of the addition of compatibilizer which improved the bonding between matrix and the filler. However, Figure 3 reveals that addition of compatibilizer shows better dispersion of nanoparticles compared to Figure 2.

3.3 Mechanical properties Mechanical properties were studied for pure and functionalised graphene filled PP matrix is shows in the Figure 4. For the uncompatibilized system (PG1 to PG4) the tensile strength gradually enhance with increase in nanofiller content. The tensile strength values of the samples PG1, PG2, PG3 and PG4 increased by 18%, 22%, 36% and 47% respectively when compared to pure PP. The maximum increase in tensile strength was achieved for PG4 (29.4 MPa for 2.0 wt.-% of graphene). Further loading of COOH-Gr in the PP matrix (PG5 and PG6) shows drop in tensile strength compared to PG4. The increase in addition of graphene may lead to agglomeration of nanoparticles in the PP matrix [3,19]. Hence the tensile strength starts to decrease. The blend PM1, PM2, PM3, PM4, PM5 and PM6 shows a slow gradual improvement on tensile strength when compared with pure polymer. The compatibilized system shows considerable enhancement in tensile strength compared to uncompatabilized and pure systems. The percentage enhance in tensile strength for the samples PMG1, PMG2, PMG3, PMG4 and PMG5 compared

The tensile fractured face of the specimens were imaged with scanning electron microscope (Model: JSM-5800 of JEOL Co., Acceleration voltage: 20 kV, Coating type: gold) at 1000 times zooming [15]. The surface morphology of fabricated specimens can also be studied using scanning electron microscopy [1,5,16].

3. Results and discussions 3.1 FTIR The existence of functional groups O-H and C-O are confirmed from FTIR (Figure 1) results. This is arrived by the identification of peaks at 1276, 1710 and 3432 cm-1 which match up to C-O stretching, C=O stretching and O-H stretching correspondingly. Hence the sample was confirmed to be COOH modified graphene [3,13,17]. The covalent bonds of molecules are like springs that are stiff but stretch and bend. At normal temperatures these bond vibrate in accordance with their electronic states. The stretching takes place during the absorption of infrared rays when these bonds experience a transition between the energy levels induced by photons of appropriate energy. Polímeros, 30(3), e2020035, 2020

Figure 1. FTIR spectrum of graphene and COOH modified graphene. 3/7

Natarajan, S., & Rathanasamy, R., & Palaniappan, S. K., & Velayudham, S., & Subburamamurthy, H. B., & Pal, K.

Figure 2. TEM image of PP+2 wt. % of COOH-Gr.

Figure 3. TEM image of compatabilized PP + 2 wt.-% of modified graphene.

Figure 4. Tensile strength of PP, PG and PMG. 4/7

to pure PP are 18%, 31%, 40%, 51% and 62% respectively. The maximum tensile strength of 32.4 MPa was achieved for the sample PMG5. The same shows 62% enhancement in tensile strength match up to pure PP. The presence of compatibilizer enhances the dispersal of graphene in the base PP matrix [9,10,20]. However with increase in filler content beyond 5 wt.-% for PMG5, the tensile strength starts to decrease. For both uncompatibilized and compatibilized system the elastic modulus gradually increases with increase in addition of graphene in the PP matrix. The Young’s modulus for blend PM1, PM2, PM3, PM4 and PM5 increases upto 6% compared to pure polymer, later on decreases for PM6. For the uncompatibilized system 5%, 7%, 8%, and 10% increase of modulus was obtained for the samples PG1, PG2, PG3 and PG4 respectively. From Figure 5, on further adding of COOH-Gr (2.5 wt.-% and 3 wt.-%) decreases the elastic modulus of PG5 and PG6. Similarly for the compatibilized system shows 10%, 12%, 13%, 18% and Polímeros, 30(3), e2020035, 2020

Comparison of MA-g-PP effectiveness through mechanical performance of functionalised graphene reinforced polypropylene

Figure 5. Young’s modulus of PP, PG and PMG.

Figure 8. SEM image of fracture surface with 2.5 wt.-% COOHGr + MA-g-PP + PP.

3.4 DMA studies

Figure 6. Storage modulus for PP, PG and PMG.

Figure 6 characterizes the temperature dependence of storage modulus (E’) of nanocomposite. The nanocomposite containing modified graphene showed a development in storage modulus on compared with PP and PG. An increase of 19.02% was observed for PP and graphene nanocomposite compared to pure PP. Several studies [21,22] have confirmed that inclusion of graphene (particles, sheets) can boost the thermal behavior of the polymer to different degree. Some studies [23,24] exposed that incorporation of functionalized graphene (particles, sheets) accelerates the thermal degradation of polymer in inert condition. The enhancement in storage modulus is additional evidence for the compatibilized nanocomposites. The compatibilized COOH-Gr system showed an increase of 43.48% in storage modulus compared to PP. With increase in filler loading the peak intensity of tan δ consecutively decreases. This may be due to better reinforcing effectiveness of the filler with the PP matrix. Figure 7 signifies the tan δ of PP, PG and PMG. The reduction in tan δ peak confirms the minimum heat buildup and as a result, leads to better damping characteristics of the nanofiller incorporated PP matrix.

3.5 Morphological study

Figure 7. Tan δ peak for PP, PG and PMG.

20% increase of modulus for the samples PMG1, PMG2, PMG3, PMG4 and PMG5 respectively. On further addition of COOH-Gr (3 wt.-%) decreases the modulus of PMG6. At last, this overall increase of mechanical properties is due to homogeneity in mixing and better dispersion of functionalised fillers in matrix. Polímeros, 30(3), e2020035, 2020

The polar nature of COOH-Gr will not be compatible upon mixing in bulk non-polar PP matrix. In order to avoid the compatibility mismatch, a polar compatibilizer MA-g-PP derived from base PP is used as a medium to attain effective reinforcement of COOH-Gr in base matrix. However, PMG5 sample shows better tensile strength and Young’s modulus. So, the SEM study of PMG5 sample has also been done to add more proof of the obtained results. SEM image (Figure 8) shows a fibrillated morphological texture, which proves the enhancement in compatibility between polar COOH-Gr and non-polar PP matrix by introducing MA-g-PP as compatibilizer [18,20]. 5/7

Natarajan, S., & Rathanasamy, R., & Palaniappan, S. K., & Velayudham, S., & Subburamamurthy, H. B., & Pal, K.

4. Conclusion In summary, PP nanocomposite was prepared by incorporating COOH functionalised graphene to investigate the enrichment of mechanical and thermal stability of PP. FTIR results confirm the presence of functional groups in modified graphene by corresponding absorption levels. TEM image justifies the successful exfoliation of single layer graphene, uniformly distributed in PP matrix. Significant enhancement in the mechanical properties and dynamic mechanical properties was achieved for compatibilized PP nanocomposites compared to uncompatibilized and pure system. The surface morphology shows a fibrillated morphological texture, which proves the enhancement in compatibility between polar COOH-Gr and non-polar PP matrix by introducing MA-g-PP as compatibilizer. As of the results, it confirms that addition of nanofiller reinforcement into the polymer matrix, compatibility match between the fillers and base matrix plays a role in varying the properties of the polymer.

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mechanical properties of compatibilized polypropylene nanocomposites. Polymer, 46(23), 10237-10245. http:// dx.doi.org/10.1016/j.polymer.2005.08.035. 10. Mohaiyiddin, M. S., Lin, O. H., Akil, H. M., Yee, T. G., Adik, N. N. A. N., & Villagracia, A. R. (2016). Effects of polypropylene methyl polyhedral oligomeric silsesquioxanes and polypropylene-grafted maleic anhydride compatibilizers on the properties of palm kernel shell reinforced polypropylene biocomposites. Polímeros: Ciência e Tecnologia, 26(3), 228-235. http://dx.doi.org/10.1590/0104-1428.2038. 11. Kalaitzidou, K., Fukushima, H., & Drzal, L. T. (2007). Mechanical properties and morphological characterization of exfoliated graphite–polypropylene nanocomposites. Composites. Part A, Applied Science and Manufacturing, 38(7), 16751682. http://dx.doi.org/10.1016/j.compositesa.2007.02.003. 12. Sahoo, S., Karthikeyan, G., Nayak, G., & Das, C. K. (2012). Modified graphene/polyaniline nanocomposites for supercapacitor application. Macromolecular Research, 20(4), 415-421. http://dx.doi.org/10.1007/s13233-012-0042-1. 13. Rajasekar, R., Kim, N. H., Jung, D., Kuila, T., Lim, J. K., Park, M. J., & Lee, J. H. (2013). Electrostatically assembled layer-by-layer composites containing graphene oxide for enhanced hydrogen gas barrier application. Composites Science and Technology, 89, 167-174. http://dx.doi. org/10.1016/j.compscitech.2013.10.004. 14. Lei, S., Hoa, S. V., & Ton-That, M.-T. (2006). Effect of clay types on the processing and properties of polypropylene nanocomposites. Composites Science and Technology, 66(10), 1274-1279. http://dx.doi.org/10.1016/j. compscitech.2005.09.012. 15. Pal, K., Rajasekar, R., Kang, D. J., Zhang, Z. X., Pal, S. K., Das, C. K., & Kim, J. K. (2010). Effect of filler and urethane rubber on NR/BR with nanosilica: morphology and wear. Journal of Thermoplastic Composite Materials, 23(5), 717-739. http://dx.doi.org/10.1177/0892705709355234. 16. Nikje, M. M. A., Moghaddam, S. T., & Noruzian, M. (2016). Preparation of novel magnetic polyurethane foam nanocomposites by using core-shell nanoparticles. Polímeros: Ciência e Tecnologia, 26(4), 297-303. http:// dx.doi.org/10.1590/0104-1428.2193. 17. Choi, E.-Y., Han, T. H., Hong, J., Kim, J. E., Lee, S. H., Kim, H. W., & Kim, S. O. (2010). Noncovalent functionalization of graphene with end-functional polymers. Journal of Materials Chemistry, 20(10), 1907-1912. http://dx.doi. org/10.1039/b919074k. 18. Song, P., Cao, Z., Cai, Y., Zhao, L., Fang, Z., & Fu, S. (2011). Fabrication of exfoliated graphene-based polypropylene nanocomposites with enhanced mechanical and thermal properties. Polymer, 52(18), 4001-4010. http://dx.doi. org/10.1016/j.polymer.2011.06.045. 19. Simanke, A. G., Azeredo, A. P., Lemos, C., & Mauler, R. S. (2016). Influence of nucleating agent on the crystallization kinetics and morphology of polypropylene. Polímeros: Ciência e Tecnologia, 26(2), 152-160. http://dx.doi. org/10.1590/0104-1428.2053. 20. Kumar, K. V. M., Krishnamurthy, K., Rajasekar, R., Kumar, P. S., Pal, K., & Nayak, G. C. (2019). Influence of graphene oxide on the static and dynamic mechanical behavior of compatibilized polypropylene nanocomposites. Materials Testing, 61(10), 986-990. http://dx.doi.org/10.3139/120.111411. 21. Liang, J., Huang, Y., Zhang, L., Wang, Y., Ma, Y., Guo, T., & Chen, Y. (2009). Molecular‐level dispersion of graphene into poly (vinyl alcohol) and effective reinforcement of their nanocomposites. Advanced Functional Materials, 19(14), 2297-2302. http://dx.doi.org/10.1002/adfm.200801776. Polímeros, 30(3), e2020035, 2020

Comparison of MA-g-PP effectiveness through mechanical performance of functionalised graphene reinforced polypropylene 22. Xu, Y., Hong, W., Bai, H., Li, C., & Shi, G. (2009). Strong and ductile poly (vinyl alcohol)/graphene oxide composite films with a layered structure. Carbon, 47(15), 3538-3543. http://dx.doi.org/10.1016/j.carbon.2009.08.022. 23. Lee, Y. R., Raghu, A. V., Jeong, H. M., & Kim, B. K. (2009). Properties of waterborne polyurethane/functionalized graphene sheet nanocomposites prepared by an in situ method. Macromolecular Chemistry and Physics, 210(15), 1247-1254. http://dx.doi.org/10.1002/macp.200900157.

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24. Raghu, A. V., Lee, Y. R., Jeong, H. M., & Shin, C. M. (2008). Preparation and physical properties of waterborne polyurethane/functionalized graphene sheet nanocomposites. Macromolecular Chemistry and Physics, 209(24), 24872493. http://dx.doi.org/10.1002/macp.200800395. Received: June 17, 2020 Revised: Aug. 07, 2020 Accepted: Oct. 13, 2020

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ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.09119

Isolation and characterization of micro cellulose obtained from waste mango Miguel Angel Lorenzo-Santiago1 and Rodolfo Rendón-Villalobos2*  Centro de Ciencias de Desarrollo Regional, Universidad Autónoma de Guerrero, Acapulco, Guerrero, México Centro de Desarrollo de Productos Bióticos, Instituto Politécnico Nacional, Yautepec, Morelos, México

*rrendon@ipn.mx

Abstract Cellulose fibers are used in polymeric matrices due to their thermal and mechanical capabilities. These biopolymers can be isolated from different natural sources. In this study, micro cellulose was obtained from mango fibrous endocarp (Mangifera caesia Jack ex Wall) waste. Isolation was performed using sulfuric acid and sodium hydroxide as removal agents of lignin and hemicellulose. A comparative analysis between native fiber (NF) and micro cellulose (MC) was performed, using FTIR, DSC and SEM techniques to assess their composition, thermal, structural and crystallinity capacities. The structures of the fibers were not damaged due to the chemical treatments received and their sizes ranged between 40 and 400 μm in length. The extraction of cellulose from mango waste represents an important start in obtaining biopolymers that can be used in the food, pharmaceutical, and other industries. Keywords: biopolymers, mango waste, micro cellulose. How to cite: Lorenzo-Santiago, M. A., & Rendón-Villalobos, R. (2020). Isolation and characterization of micro cellulose obtained from waste mango. Polímeros: Ciência e Tecnologia, 30(3), e2020036. https://doi.org/10.1590/0104-1428.09119

1. Introduction In recent years, different ways have been sought to control and reduce pollution caused by petroleum-derived materials[1-3], and one of the most used alternatives in the world is recycling[4]. However, and even though this conventional method reduces the ecological impact of plastic has not been possible to reduce the large volumes of those kinds of materials that are produced each year[5]. This is mainly due to the fact that a great many products are manufactured in way that make the plastic content difficult to separate and, therefore very little plastic is currently recycled. An estimated 8.3 billion metric tons of plastic are produced worldwide, and only 9% goes through recycling process[4]. Another alternative for reducing environmental impact of plastics that has been proposed is the use of polymeric materials combined with pro-oxidizing substances[2,6]. However, the wastes the materials generate after they are degraded consist of heavy metals and petroleum products[7-9], which tend to accumulate in the environment. To reduce the impact caused by petroleum products, there is need to look for materials that have a polymer matrix formed by biopolymers (natural polymers)[10]. Since the polymer matrix is biodegradable, the composite which is its integral part is also expected to be biodegradable and environmentally friendly[11]. It has an advantage over synthetic polymers because it can be totally biodegraded in products such as CO2, water and also organic fertilizer[12]. Bioplastics and biomaterials are composed of biopolymers, biomass derivatives or blends of natural and synthetic

Polímeros, 30(3), e2020036, 2020

polymers, which will help mitigate environmental issues in the long term. Agriculture in Mexico is an important productive sector. The role played by the agricultural sector in the other sectors of Mexico such as economic, social and environmental determines its greatest impact on the country’s development. However, post-harvest residues (bagasse, husks and seeds) represent an environmental problem due to the followings: lack of distribution channels, pests, environmental policies and inadequate management of the large volumes of residues that are generated. Therefore, agricultural residues must be considered as a source sustainable and renewable of products such as cellulose fiber used as a reinforcement for bioplastics and biomaterials, due to their high availability and potential[13,14]. Waste mango (seeds, shells and cotyledons) currently have little impact on industries. However, these wastes have been studied to obtain active and functional compounds from mango peel[15], antioxidants from cotyledon[16], to use fibrous endocarp in the paper industry[17], and pulp to extract reducing sugars, organic acids and obtain alcohols[18]. Unlike other agricultural residues, mango (Mangifera caesia Jack ex Wall) is composed of more than one biopolymer, with starch and cellulose being the highest proportion[19]. In addition to being the most abundant natural polymer in the world, cellulose is renewable, biodegradable and is a classic example of a matrix reinforcement biopolymer[20,21] and is expected to be able to reduce and replace many non-renewable materials[22]. The fibrils have two regions

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

Lorenzo-Santiago, M. A., & Rendón-Villalobos, R. known as crystalline (ordered regions) and amorphous (disorder regions). These regions can be hydrolysed with acid, which makes the crystalline region less reactive. This results in obtaining microfibrils, nanofibrils and cellulose nanocrystals[23]. Cellulose fibers have an inherent structural hierarchy that originates from their different biological sources. Obtaining cellulose from different lignocellulosic fibers has attracted the interest of the scientific and industrial community in recent years[24,25]. The native fibers are longitudinal, rigid and difficult to modify due to the components that form them. These compounds work as a natural glue that gives firmness and shape. The fibers are hydrophobic. This is why, it is necessary to perform a chemical and thermal treatment, that can help to remove glue from the fibers without damaging the structure of the micro-cellulose fibers. The particle size reduction in the cellulose fibers through thermal treatment, generates major areas of contact with chemical substances such as acids and alkalis. Likewise, with the ultrasonic treatments crystallinity increase and morphology of cellulose fibers can be improved; in the same way individual fibrils could be obtained from waste mango and which serve in developing a bio-based material. In the present study, cellulose fibers were obtained from waste mango fibrous endocarp. The shape, size and structure of the fibers were characterized as well as the decrease in lignin and hemicellulose subjected to a modified acidalkaline treatment was analysed by wet chemistry methods. Taking advantage of waste residues facilitates the use of cellulose as a biopolymer and reduces the impact generated by inappropriate disposal of agricultural waste and are able to generate profits.

2. Materials and Methods 2.1 Cellulose insulation Fibrous endocarp of Ataulfo mango fruits (Mangifera caesia Jack ex Wall) was obtained in the local market in El Arenal, Guerrero, Mexico, and used as the raw material for this study. Reagents used to obtain and extract the native cellulose fibre (NF) and chemically modified cellulose fibre (designated micro-cellulose, MC) were: Sodium hydroxide (CAS 1310-73-2), sulfuric acid (CAS 7664-939), hydrogen peroxide (CAS 7722-84-1) and distilled water (CAS 7732-18-5).

2.2 Pre-treatment of mango fibre The mango endocarp was cut into squares of approximately 1 cm2. 50 g was weighed and hydrated using distilled water. It was stirred at room temperature for two hours and subsequently dried in a continuous flow oven at 40 ° C for 24 hours[26].

2.3 Alkaline treatment The dried fibrous endocarp was treated in a solution of 2% NaOH (w/v) in a 1:20 ratio (fibre: solution). The solution was continuously stirred at 80 ° C for two hours. After the alkaline treatment was done, the fibres were washed with distilled water and dried at 40 ° C for 24 h. The dried fibres 2/8

were grinded using a E3303.00 mini cutting mill (Eberbach Corp.), equipped with a stainless steel blades to separate the bonded fibres formed after the drying process. After the milling was completed, the fibres were stored in an airtight bag[27].

2.4 Whitening The fibres were bleached in a 1:20 ratio (fibre: solution) in a solution composed of H2O2 (v/v) and 4% NaOH (w/v), and stirred for two hours at 50 ° C. After the treatment, the fibres were washed with distilled water and dried at 40 ° C for 24 h[28].

2.5 Acid Hydrolysis The fibres were subjected to 52% sulfuric acid (w/w) in a 1:20 ratio (fibre: solution), and stirred for two hours at 45 ° C. The reaction was stopped by thermal shock using a vessel with distilled water at a temperature of 10 ° C[27].

2.6 Sonication Sonication process was applied after acid hydrolysis in order to disintegrate the fibres into microfibres in water dispersion to reduce the viscosity and allow efficient propagation of the vibration through the dispersion. Szymańska-Chargot et al.[29] and Kasuga et al.[30] techniques were modified and the dispersion was then subjected to sonication for 5 min at a power of 99 W using a Branson 5510MT ultrasonic cleaner. Subsequently, the samples were centrifuged at 6300 x g at 4 °C for 15 min and dried at 40 °C for 24 h[27].

2.7 Chemical analyses The lignin content was determined according to Klason method[31]. The dried fibres was extracted by using Soxhlet method[32] with ethanol-benzene solvent for 5 h. The residue extract was oven dried at 103 o C for 1 h, and later treated with 72% H2SO4 for 2 h while stirring at 37 o C. The material was then diluted to 3% H2SO4 and then refluxed at 80 o C for 4 h; the sample was oven dried at 105 o C for 1 h and cooled in a desiccator to obtain constant weight. The lignin content was calculate using the following equation[31]: Lignin % =

A100 W

(1)

Where A is weight of lignin and W is oven-dry weight of test specimen. Holocellulose was measured by means of the Wise chlorite method[33] through dried fibers treatment with an acidified sodium chlorite water solution in acidic medium at 70-80 ºC for 1 h. The process was repeated until the product became white. Determination of cellulose to dry fiber samples was using the Seifert method[34] wherein the cellulose is separated from the plant material in the process of its processing with aqueous solutions of acetylacetone, dioxane, and hydrochloric acid. Hemicellulose content was calculated theoretically from the difference in the contents of holocellulose and cellulose. Polímeros, 30(3), e2020036, 2020

Isolation and characterization of micro cellulose obtained from waste mango 2.8 Scanning Electron Microscopy (SEM) Cellulose samples were sprinkled on a double adhesion carbon conductive tape, which was previously fixed on aluminium support of the scanning electron microscope Carl Zeiss EVO LS 10. Subsequently, they were observed at a voltage of 2.5 kV, with a resolution of 3-10 nm; the micrographs were taken at 100, 500 and 1500 magnifications[35].

3.2 Structural characterization

2.9 Fourier Transform Infrared Spectroscopy (FTIR) Fourier transform infrared spectroscopy (FTIR) studies were carried out using a Perkin-Elmer-Spectrum 100/100 N FT-IR infrared spectrometer. The infrared region was within a range of 4000-650 cm-1 in the transmittance mode, with a resolution of 16 cm-1 and 8 scans[36].

2.10 Differential Scanning Calorimetry (DSC) A differential scanning calorimeter model TA 2010 was used. The equipment was calibrated with indium (melting point of 156.4 ° C and enthalpy of 6.8 cal/g). Between 5 and 10 mg of the sample was weighed, using a micro balance with an accuracy of ± 0.01 mg, a heating program ranging from 20 to 400 ° C and a heating rate of 10 ° C/min[30]. This generated an inert atmosphere in the cell by circulating 50 mL/min of nitrogen gas (99.9% purity).

2.11 Crystallinity The degree of crystallinity was determined based on the DSC enthalpy of fusion data using the following equation[37]: Xc =

∆H f ∆H o

(2)

Where XC is the crystalline fraction, ΔHF represents the enthalpy of fusion measured by DSC and ΔHO is the enthalpy of fusion for 100% crystalline polymer.

3. Results and Discussions 3.1 Chemical Analyses Despite the economic importance of agricultural residues there are no studies on the chemical composition of the mango endocarp. In general, the studies of mango chemistry are limited to chemical characteristics based mainly on mango agroindustrial waste, fruit by-products and silage[38-41]. The results of lignin and calculated hemicellulose of native (NF) and alkaline modified (MC) fibers decreased, from 10.32 wt % to 0.64 wt % for NF and 29.75 wt % to 1.11 wt %, respectively. Regarding the lignin content, the data from this investigation (10.32 wt %) is close to those found for Guzmán et al.[38], different values are obtained when the food-processing industry method is employed (10.32 wt % vs. 9.0 wt %). This method is similar to Klason lignin extraction process, except a final step of 1 h of calcination at 550 ° C. Nevertheless, a large difference is observed between our lignin results: 5.76 wt % and 6.97 wt % of lignin in mango peel and mango by-products, respectively[39,40]. Concerning the hemicellulose content from this study (29.75 wt %) is lower than the values (31.1 wt % and 31.75 wt %) found Polímeros, 30(3), e2020036, 2020

for mango peel and mango by-products[38,41]. This can be explained by the fact that it seems that in the foodprocessing industry method, the neutral detergent solution used for determining the Neutral Detergent Fibre (NDF = Hemicellulose + Cellulose + Lignin) content does not remove completely hemicelluloses and is, therefore, this method is not indicated for this work.

The native fibres presented very dense outer layers covering the cellulose sacs. The layers are considered to be composed of lignin (Figure 1a). Lignin wraps the cellulose in the inner part covering it with its stiffness and hardness property. Small globular particles and pores along the fibres and impurities can be seen on the surface. Figure 1b shows an accumulation of long fibres, rolled and small traces of waste. This is attributed to the modification of the alkaline treatment using NaOH. Cellulose fibres can be seen more clearly, but there is still a need to remove part of the lignin and hemicellulose that holds them together. Cellulose fibres are more available due to the removal of lignin, and their characteristics (e.g. shape, surface) are similar to those of the commercially available cellulose (Civeq™); the SEM images (Figure 1b and c) evidence that separate individual fibres have also similar morphologies with smooth surfaces. The fibres went through a whitening process with the help of hydrogen peroxide. The peroxide helped in eliminating some waxes, tannins, lignin and alkaline treatment residues. In Figure 2a, completely separated fibres can be seen; there are no traces of lignin, as well as porous components. In this procedure, the decrease in thickness, shape, size and discoloration of the fibres is more evident. After bleaching, the fibres had sizes ranging from 40 to 400 μm in length with a uniform surface and rectangular shape (Figure 2b). Owi et al.[42] characterized micro celluloses chemically isolated from fruit bunches and sugarcane bagasse as well as commercial cellulose (Sigma™) with characteristics similar to those obtained in this study after extraction, particularly in fibre shape and surface. In a similar manner, Nagalakshmaiah et al.[43] report the shape and diameter of chili fibres by an acid hydrolysis treatment. These results are consistent with the ones presented in this study. This is consistent with the report of Atiqah et al.[44] that amorphous portions like hemicellulose and lignin were removed after applying alkali and bleaching treatment on raw kenaf fibre.

3.3 FTIR analysis Figure 3 shows the native and micro cellulose fibre spectra obtained by chemical treatment. A peak at 894 cm-1 can be observed in the fibres with alkaline treatment. This signal is attributed to the anomeric carbon present in cellulose. This signal is usually in the polysaccharide absorbing region of 950-700 cm-1[45], which also reveals the component structures of ß-glucans. In the native fibre, this signal is less intense because the lignin coating prevents specific cellulose signals that are present. A narrow absorption band in the region between 1170 and 920 cm-1 ascribed to the polysaccharide bands, shows the stretching vibrations of C – O, C – C and C – OH bonds which 3/8

Lorenzo-Santiago, M. A., & Rendón-Villalobos, R.

Figure 1. Micrograph of (a) native mango fiber, (b) alkaline modified fiber (c) commercial cellulose (Civeq™).

Figure 2. Micrographs of (a) fiber after whitening (b) and thermal sonication.

are in cellulose, lignin and hemicellulose[46,47]. A strong signal was present in the treated fibre (MC). The increase is based on the breaking of bonds during the acid-base treatment. The native fibre showed a peak at 1230 cm-1. This is attributed to the stretching of C – O – C bonds of the alkyl-aryl ether, a compound present in the structure of lignin, and which disappears after chemical treatments. This means that the lignin was partially removed in MC as further confirmed by lignin results obtained in this work, 0.64 wt %, determined by Klason method.

Figure 3. Infrared spectra for native fibre (NF) and micro cellulose (MC). 4/8

In the native fibre, a characteristic peak of the C–C unsaturated bond of lignin was observed at 1623 cm-1, as well as the water band at 1636 cm-1 decreasing its intensity slightly in MC. Likewise, a peak of 1740 cm-1 is observed, Polímeros, 30(3), e2020036, 2020

Isolation and characterization of micro cellulose obtained from waste mango which is associated with the stretching of C = O groups linked to aliphatic carboxylic acid and ketone. It could also be related to the presence of hemicellulose[48]. Both peaks decrease when the sample is subjected to acid treatment, which allows the removal of lignin and hemicellulose, from 10.32 to 0.64 wt % and from 29.75 to 1.11 wt %, respectively. Between the range of 2940 and 2800 cm-1, there are two peaks in the native cellulose, which are assigned to the asymmetrical and symmetrical stretching vibrations of alkyl, aliphatic and aromatic rings (H – C – H) groups present in hemicellulose and lignin. These bands were ascribed to the aliphatic materials such as cutin, waxes and cutan[49]. The signal decreases when the sample is subjected to an acid-base treatment, due to the removal of these components[38]. 1740 cm-1 band associated with the stretching of C=O in ester groups disappeared after basic hydrolysis, confirming the partial removal of hemicellulose, which presented variation in hemicellulose content (e.g. 1.11 wt % compared to 29.75 wt % for NF sample).

3.4 Thermal characterization As for the DSC results, both samples studied in nitrogen atmosphere show a similar trend and, all the reactions during the decomposition of the cellulose were endothermic process in inert atmosphere (Figure 4). The native endocarp fibre presented three steps of thermal degradation in relation to cellulose, lignin and hemicellulose. This is in line with a research done by Stevulova et al.[50]. NF shows an endotherm at 113.64 °C and a plateau at 145.3 °C; while that of MC is at 135.11 °C. This is mainly due to the elimination of moisture by the heat range applied in the DSC, as well as reducing sugars and other volatile components joined to the fibres, that are thermally decomposed. Espinosa-Andrews and Urias-Silvas[51] mention that there is a humidity loss, specifically in fructans, in temperatures under 130 °C. In addition to this, a small endotherm at 232.12 °C can be observed in NF probably due to the degradation of hemicellulose, reducing sugars and those hemicelluloses that are still joined to the cellulose-lignin complex[48]. NF and MC presented endotherm peaks at 325.52 °C and 335.25 °C, respectively. The endothermic peak observed in a 220-370 °C temperature range corresponds to that used for the simultaneous decomposition of hemicelluloses, cellulose and lignin. Based on the report of Miyahara et al. [52] , this endothermic peak represents the degradation of hemicellulose. As reported by Xiao et al.[53], cellulose is degraded within the range of 250 and 345 °C. The enthalpies in the initial endotherm peak of the NF sample (121.92 J/g) are lower compared to those of the MC sample (183.27 J/g). This could be due to a greater presence of traces of waste and volatile components that were impregnated through acid-alkali hydrolysis. Therefore, a greater amount of heat energy is needed for degradation to occur. Conversely, with respect to final endotherm peak for both samples, the enthalpy of MC (172.92 J/g) is lower compared to that of the native fibre (178.32 J/g), which could be due to the elimination of larger sugar chains through acid-alkali hydrolysis. Besides, under nitrogen atmosphere, cellulose is more resistant to thermal treatment, probably due to its crystalline structure[50]. The range at which lignin is Polímeros, 30(3), e2020036, 2020

Figure 4. DCS curves of native (NF) and modified (MC) fibre. Table 1. Percentage of crystallinity of native fibre (NF) and microcellulose (MC). Sample NF MC

Crystallinity % 75.18 77.53

decomposed is between 250 and 600 ° C. The analysis was performed up to 400 ° C, which indicates the degradation of 40% of the total lignin in the native sample. In the composition of native cellulose, it can be seen that the compound that degrades first is hemicellulose, followed by cellulose and lignin[53].

3.5 Crystallinity Applying the Equation 1 and with the use of the DSC enthalpy of fusion and taking 230J/g as the value for the fully crystalline Polycaprolactam (PA6)[54], a lower value of crystallinity for native fibre (NF) was found compared to microcellulose (MC) (Table 1). As reported by Li et al.[55], the percentage of crystallinity in native fibre will always be lower, due to the presence of hemicellulose and lignin, amorphous compounds that encapsulate the fibrils. It is believed that the crystallinity index in micro cellulose increases due to the elimination of the amorphous regions of cellulose, in which hydrolysis drives the breakdown of glycosidic bonds, releasing individual crystals[56]. The crystallinity of mango micro cellulose reported in this study is similar to that reported by Bolio-López et al.[57] in banana peel micro cellulose (72% crystallinity) but greater than that obtained by Naranjo et al. [47] in agave fibre (50.07%). The greater the crystallinity of a fibre, the better the thermal stability and properties it can provide in a polyblend to obtain biodegradable materials[58].

4. Conclusions Cellulose microfibers were isolated from waste mango fibrous endocarp (Mangifera caesia Jack ex Wall), after which they were subjected to an acid-alkaline treatment and an ultrasonic treatment that helped homogenize their sizes. The acid-alkaline treatment helped decrease the size of the fibres and was also involved in the removal of lignin and cellulose. Ultrasonication process is a very effective way to depolymerize cellulose and that led to reduce the formation 5/8

Lorenzo-Santiago, M. A., & Rendón-Villalobos, R. of agglomeration and increase the dispersion of cellulose disintegrating into individual microfibers. We hope that with this simple chemical-ultrasonic treatment applied to cellulose will further stimulate interest in isolating microcellulose fibers from natural sources. The use of agricultural waste as a natural source of biopolymers has become an excellent option for innovation and the development of new products with potentially biodegradable properties by having a polymeric matrix formed by biopolymers and, their utilization can minimize negative environmental effects.

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Polímeros, 30(3), e2020036, 2020

ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.05720

Evaluation of fracture toughness of epoxy polymer composite incorporating micro/nano silica, rubber and CNTs Ronaldo Câmara Cozza1*  and Vikas Verma2  Departamento de Engenharia Mecânica, Centro Universitário da Fundação Educacional Inaciana “Padre Sabóia de Medeiros” – FEI, São Bernardo do Campo, SP, Brasil 2 Thermochemistry of Materials SRC, National University of Science and Technology – NUST/MISiS, Moscow, Russia

*rcamara@fei.edu.br

Abstract In engineering applications, fracture toughness is an essential requirement that determines the life of a material. Epoxy polymers are widely used in fibre-reinforced composite materials. Due to their structural efficiency and durability, the use of adhesive and composite materials based on epoxy polymers is widespread in aerospace and automobile industries. In this paper fracture toughness of hybrid epoxy polymer composite with addition of nano/micro figures of silica, rubber and carbon nano tubes (CNTs) is evaluated. It is observed that silica addition promoted nano toughening effect with plastically deformation capability in epoxies. Rubber and multi walled CNTs increased the toughness with negligible reduction in stiffness in epoxies. Future research emphasis can be laid on crucial understanding of stress transfer mechanisms and interfacial bond strength between nano particles – epoxy system and on nanofillers modified epoxies as matrices or interleafs for carbon or glass fiber composites to increase the interlaminar delamination toughness. Keywords: CNT, epoxy, rubber, silica. How to cite: Cozza, R. C., & Verma, V. (2020). Evaluation of fracture toughness of epoxy polymer composite incorporating micro/nano silica, rubber and CNTs. Polímeros: Ciência e Tecnologia, 30(3), e2020030. https://doi. org/10.1590/0104-1428.05720

1. Introduction Epoxy polymers are widely used in fibre-reinforced composite materials[1]. The use of adhesive and composite materials based on epoxy polymers is widespread in the aerospace, automobile and wind-energy industries due to their structural efficiency[2]. Their outstanding temperature resistance and durability to weathering, fuel, de-icing fluids, etc. leads to them invariably being the preferred materials, compared to acrylics and polyurethanes, for external aerospace applications. Furthermore, their insulating properties, good temperature resistance and ease of processing also allow epoxy polymers to be used extensively in the electronics and electrical components[2]. However the limitations of epoxies as structural materials are due to their poor resistance to the initiation and growth of cracks. Thus, improvements in their fracture performance are highly sought after by industry. However there have been several researches on thermoplastic and thermoset polymer to overcome these drawbacks by modifying epoxy resins with integration of various nanofillers as a second microphase, for advanced composite applications[2,3]. An epoxy polymers based composite material has numerous advantages as it offers excellent mechanical properties and thermal stability[3]. Hybrid polymer composites (HPC) are one of the recent developments to reduce the cost of expensive composites containing reinforcements like carbon fiber by incorporating a proportion of cheaper, low-quality fibers such as glass, textile, natural fibers, and nano figures like (silica, rubber,

Polímeros, 30(3), e2020030, 2020

CNT, clay, graphene). A lot of work has been carried in past to improve the fracture toughness of a polymer composite. Most of the reviewed data in this paper are based on the classification of incorporating different fibers, CNTs and micro/nano size silica, rubber particles in the epoxy resin as shown in schematic diagram Figure 1. Following the above brief introduction, Section 2 presents an overview of fracture toughness evaluation of epoxy resin modified with silica nano and micro particles; Section 3 is concerned with fracture toughness behaviors of epoxy resin modified by rubber particles, Section 4 deals with the effect of fracture toughness of an epoxy resin modified by different CNTs with varying wt.%. Section 5 describes the recent development achieved towards fracture toughness evaluation of different hybrid polymer composite. Section 6 concludes with research gap to address for some unattended problem in hybrid polymer composite.

2. Overview of Fracture Toughness Evaluation of Epoxy Resin Modified with Silica Nano and Micro Particles The use of silicate-based filler in polymer composite has been a great interest for research as it improves the mechanical properties, especially fracture toughness of epoxies[1,2].

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

Cozza, R. C., & Verma, V.

Figure 1. Materials used for toughening of epoxies.

Hsieh et al.[1] used four different epoxy polymers to modify with silica (10-20 wt.%) nano particles. They are anhydride-cured DGEBA (diglycidyl ether of bis-phenol A), polyether-amine cured DGEBA/F (diglycidyl ether of bisphenol F), and polyether-amine cured DGEBA and amine-cured TGMDA (tetraglycidylmethylenedianiIine). Single-edge notch-bend (SENB) and double-notched four point specimen have been conducted to observe the toughness differences. Toughest material identified was polyether-amine cured DGEBA/F epoxy posses low transition temperatures (TG) and high molecular weight (MC). AFM phase images of unmodified and epoxy polymers containing the silica nanoparticles, are shown in Figure 2a, b. Unmodified epoxy polymers showed as homogeneous thermoset polymers whereas epoxy polymers having silica nanoparticles, exhibited a very well-dispersed phase of silica nanoparticles, with no indications of agglomeration. Predominant toughening mechanisms in epoxy polymers containing silica nanoparticles were identified as localized shear bands initiated by the stress concentrations around the periphery of the silica nanoparticles with debonding features of the silica nanoparticles followed by subsequent plastic void growth of the epoxy polymer. Bray et al.[2] modified the piperidine-cured epoxy polymer by adding silica nano particles of varying diameter. Toughness increased steadily as the concentration of silica-nano particles was increased with no significant effect of particle size being observed. The toughening mechanisms were identified as the formation of localized shear-band yielding in the epoxy matrix polymer initiated by the silica nanoparticles, and debonding of the silica nanoparticles followed by plastic void growth of the epoxy matrix polymer as shown in Figure 3a, b. It was observed that incorporation of nanomaterials in the polymer matrix maintaining a homogeneous dispersion and good adhesion is considered to be a highly effective in improving the mechanical properties of resins[3]. Finite element analysis can also be used to estimate the fracture behavior of hybrid polymer composite (HPC)[4]. Epoxyfunctionalised toughening particles results in superior tensile, compressive and impact properties which can be tested fracture toughness testing methods like indentation strength 2/14

method (IS), the single edge notched beam (SENB) and the Chevron notched beam method (CN)[5,6]. The “nano” effects of silica (< 25 vol.%) or rubber (> 10 wt.%) nanoparticles in toughening epoxy resin are confirmed by Marouf et al.[7]. Dittanet et al.[8] studied the effect of silica nanoparticle content and particle size on glass transition temperature (TG) coefficient of thermal expansion (CTE), Young’s modulus (E), yield stress (s), fracture energy (GIC) and toughness (KIC). The addition of silica nanoparticles does not have significant effect on TG or yield stress of epoxy resin. The CTE (coefficient of thermal expansion) values of nano silica filled epoxies were found to decrease with increasing silica nanoparticle content. The young’s modulus was found to be significantly improved with the addition of silica nanoparticles & increased with increasing fiber content. Particle size did not exhibit any effect on the young’s modulus. The fracture toughness & fracture energy showed significant improvement with the addition of silica nanoparticles and increasing with increasing filler material. Observation of the fracture surfaces showed evidence of debonding of silica nanoparticles, matrix void growth, and matrix shear banding, which are credited for the increases in toughness for nano silica-filled epoxy systems. Dittanet et al.[9] formulated the epoxy mixtures of two different size distributions of silica-based particles (micron-size and nanometer-size) for possible synergistic toughening effects. The influence of bimodal particle size mixture and silica particle content on the glass transition temperature, coefficient of thermal expansion, Young’s modulus, yield stress, and fracture toughness, was investigated. Fracture toughness improved by approximately 30% when mixtures of microparticles and nanoparticles were used. The increased fracture energy found in bimodal systems is mainly due to the contribution from the plastic void growth mechanism, by the silica nanoparticles, together with matrix shear banding. Johnsen et al.[10] examined the toughening mechanisms of epoxy/nanosilica composites using single-edge notch bend (SENB) specimens. Fracture energy increased from 103 J/ m2 to 291 J/m2 at 4.0 wt.% loading, 182% increment. Further increase in nano filler loading by 7.8 wt.% increased the fracture toughness performance. SEM studies showed the presence of debonding and plastic void growth as contributing factors for the enhancement of fracture toughness. Blackman et al.[11] investigated the fracture and fatigue performance of nanosilica modified epoxy polymers. Interesting observations were made by the authors when the loading of nanosilica was increased progressively. An increment of 51.7% in fracture toughness was observed in comparison with the neat epoxy. Microscopic studies were performed to analyse the nature of such enhancement. Debonding of nanosilica and plastic void growth were the factors that essentially contributed to the toughening mechanisms of the nanocomposite. The effects of particle size and volume fraction on the fracture toughness of epoxy composites filled with spherical silica particles were investigated by Adachi et al.[12] and found that the fracture toughness increased drastically as the volume fraction increased and the particle diameters decreased. Nano and micron-silica particles bidispersed epoxy composites showed that the composites had a lower fragility with higher strength and fracture toughness with increase in nanoparticles[13]. Glass transition temperature and fragility are found useful parameters for estimating Polímeros, 30(3), e2020030, 2020

Evaluation of fracture toughness of epoxy polymer composite incorporating micro/nano silica, rubber and CNTs

Figure 2. AFM phase images of the polyether-amine cured DGEBA epoxy polymer: (a) Unmodified and (b) 0.138 vF of silica nanoparticles[1].

Figure 3. (a) and (b) High resolution SEM images of the fracture surfaces for the epoxy polymers containing silica nanoparticles[1,2].

the fracture toughness of the silica particulate-filled epoxy composites[13,14]. Apart from mechanical properties epoxy resins in the uncured state have relatively high viscosities[15]. The viscosity of the uncured mixture must be sufficiently low to permit the evacuation of air bubbles. Fracture mechanics of the epoxy composites revealed that crack growth depends upon the intensity of stress at the tip of the crack[16]. Effect of particle size on the fracture behaviour of cured epoxy resin filled with spherical silica particles showed that the main crack propagation was hampered by large particles and a damage zone was formed at the main crack tip region in the large particle filled resin due to crack diversion and debonding of particle/matrix interfaces[17,18]. It was found that fracture surface morphology in unfilled and particulatefilled epoxy resins under static loading was governed by slow, sub-critical crack growth manifested by resin-particle debonding[19]. Unfilled resins exhibited unstable (or stickslip) crack propagation at low and high cross-head speeds whereas filled resins showed stable crack propagation at all speeds in double torsion test technique[20]. The research conducted on improvement of fracture toughness of epoxy composite incorporated with nano or micro-sized silica particles resulted in improvement of mechanical properties of epoxy resins such as elastic modulus, hardness, impact resistance and fracture toughness. Epoxy-nanocomposite resins filled with 12 nm spherical silica particles decreased Polímeros, 30(3), e2020030, 2020

in cure and glass transition temperature (for loadings of 10 wt.% and above) with increased silica loading[21]. Liang and Pearson[22] in their study found that no significant differences in fracture behavior were observed between the epoxies filled with different nano silica (20 nm or 80 nm NS particles). Effect of silica nanoparticles on toughness of two epoxy systems cured by Jeffamine D230 (J230) and 4,40-diaminodiphenyl sulfone (DDS), improved the toughness of J230-cured epoxy from 0.73 to 1.68 MPa.m1/2, and for the other system improved from 0.51 to 0.82 MPa. m1/2[23]. Fracture behaviours of nanosilica filled bisphenol-F epoxy resin investigated at ambient and higher temperatures showed improved elastic modulus, microhardness, impact resistance and fracture toughness of epoxy matrix with increasing nanoparticle volume content as the nanoparticles were almost homogeneously dispersed in epoxy matrix[24]. It was also shown that addition of 5 vol %. silica-nanoparticles could improve the stiffness and the toughness of an epoxy resin at the same time and the nano-reinforced material behaved more ductile and showed a bigger yielding than the pure epoxy[25]. Zhang et al.[26] found that the static/ dynamic modulus, microhardness, and fracture toughness of the nanocomposites enhanced with increasing silica content up to 14 vol.% (23 wt.%) due to homogeneously distributed nanoparticles which improved both the stiffness and toughness of the epoxy. Epoxy-silica nanocomposites 3/14

Cozza, R. C., & Verma, V. produced by dispersing silica-organosol particles in TGDDM/ DDS resin mixtures resulted an increased interfacial adhesion which improved fracture properties[27]. Organophilic layered silicates (OLS), modified by means of cation exchange, were added in amounts of 5-15 wt.% to vinylester (VE)/ epoxy (EP) hybrid resins which formed interpenetrating networks (IPN) after curing at T = 150 °C, doubled the fracture energy of the resins with 5 wt.% OLS[28]. The type of the organic cation modification of the OLS had no effect on morphology or toughness which was attributed to the coarse dispersion of the OLS and its encapsulation by the EP phase in the VE/EP hybrid resin[28]. Fu et al.[29] provided a review on the effects of particle size, interfacial strength, and filler loading on the elastic modulus, strength, fracture toughness, and impact behaviors of polymer composites containing nano/micro-size particulates. Rosso et al.[25] employed the well-dispersed silica nanocomposites for tensile and fracture tests, indicating that the addition of 5 vol.% silica nanoparticles could improve the stiffness and fracture energy to 20% and 140%, respectively. Guo and Li[30] performed compressive loading on the SiO2/epoxy nano composites under different loading rates, revealing that the compressive strength of the composites with silica nanoparticles is higher than pure epoxy at higher strain rates; nevertheless, there is no clear connection between the compressive strength and the nanoparticle contents at lower strain rates. Summarising, the use of silica nanoparticles as an effective toughening agent for epoxies which can plastically deform and the nano toughening effect is confirmed in silica/epoxy nano -composites. Crack growth originates from debonding of silica nanoparticles which promote matrix plasticity via shear yielding and void growth as the main energy dissipation mechanisms. The incorporation of silica particles uniformly dispersed in the matrix as seen in

the micrographs shown and stated by different researchers discussed above of the processed epoxy composite, either nano-size or micro-size, into neat epoxies increases fracture toughness, KIC and normalized fracture energy, GIC. Nanosize fillers are more effective tougheners than micro-size fillers over the full range of filler loading up to 60 wt.%. Table 1 summarises the published reports on silica particlefilled epoxies to evaluate the influence of these fillers on increasing fracture toughness.

3. Overview of Fracture Toughness Evaluation of Epoxy Resin Modified with Rubber Several methods have been proposed to improve the fracture toughness of epoxy resins and addition of suitable rubber before epoxy resin curing is claimed to be a successful routine[31]. Chikhi et al.[31] worked on liquid amine-terminated butadiene acrylonitrile (ATBN) copolymers containing 16% acrylonitrile at different contents to improve the toughness of DGEBA epoxy resin using polyaminoimidazoline as a curing agent and post cured at 120 °C. In modified epoxy resin, all reactivity characteristics (gel time and temperature, cure time and exotherm peak) were decreased. Where as a 3-fold increase in Izod IS was obtained by just adding 12.5 phr ATBN compared to the unfilled resin. Addition of ATBN, the Izod IS increased drastically from 0.85 to 2.86 kJ/m2 and from 4.19 to 14.26 kJ/m2 for notched and unnotched specimens respectively while KIC varies from 0.91 to 1.49 MPa.m1/2. Concerning the adhesive properties, the tensile shear strength (TSS) increased from 9.14 to 15.96 MPa just by adding 5 phr ATBN. Finally SEM analysis suggested that rubber particles cavitation and localized plastic shear yielding induced by the presence of the dispersed rubber particles within the epoxy matrix as the prevailing toughening mechanism. Zhao et al.[32] worked on enhancement in tensile strength and

Table 1. Summary on Silica/Epoxy composite. Type of Epoxy Polymers DGEBA DGEBA/F Polyether DGEBA TGMDA DGEBA

SiO2 Addition 0-20 wt.%

30 vol.%

DGEBA

Unmodified 23 74 170 Neat resin 42 μm + 23 nm 42 μm + 74 nm 42 μm + 170 nm 20 nm

DGEBF

90 nm

DGEBA

4/14

Modifier SiO2 size (nm) –

10 vol.%

Fracture Fracture energy toughness (J/m2) (MPa.m1/2) 1.45 ± 0.12, 616 ± 109, highest highest for for polyether polyether DGEBA DGEBA at 0.138 at 0.138 volume volume fraction fraction silica silica 1.11 ± 0.06 303 ± 14 2.52 ± 0.11 973 ± 17 2.89 ± 0.11 1264 ± 14 2.65 ± 0.06 1030 ± 08 1.07 ± 0.1 280 ± 20 1.96 ± 0.15 760 ±10 1.93 ± 0.12

740 ± 80

1.94 ± 0.15

750 ± 10

13 vol.%

0.59-1.42

100-460

0-7%

Increase in strain rate

–

Curing Agent

Ref.

Accelerated methylhexahy drophthalic acid anhydride

Hsieh et al.[1]

Piperidine

Dittanet and Pearson[8]

5phr piperidine

Dittanet and Pearson[9]

Methylhexahydrophthalic acid anhydride (HE600) Poly(propyleneglycol)diglycidyl ether (PPGDE)

Johnsen et al.[10] Guo and Li[30]

Polímeros, 30(3), e2020030, 2020

Evaluation of fracture toughness of epoxy polymer composite incorporating micro/nano silica, rubber and CNTs fracture toughness at 77 K of diglycidyl ether of bisphenol-F (DGEBF) epoxy using diethyl toluene diamine (DETD) as curing agent and carboxylic nitrile-butadiene nano rubber (NR) particles. Dispersion of NR in the cured epoxy samples is shown in Figure 4. It was found NR dispersion throughout the epoxy matrix with no agglomeration[32]. The tensile strength is increased compared with that of the pure epoxy. The SENB test is used to determine the KIC. Fracture toughness (KIC) is enhanced by 48.3% at 15 phr NR compared with that of the pure epoxy. Moreover, the comparison of mechanical properties between 77 K and room temperature indicated that the tensile strength, young’s modulus and fracture toughness at 77 K are higher than those at room temperature with reduced young’s modulus of epoxy resins. Chen and Taylor[33] investigated the fracture toughness and mechanical properties of an anhydride-cured diglycidylether of bisphenol-A (DGEBA)

epoxy polymer modified with poly(methyl methacrylate)-bpoly(butylacrylate)-b-poly(methyl methacrylate) (MAM). These triblock copolymers toughen the epoxy polymer significantly, with only slight reductions in the mechanical and thermal properties of the epoxy polymer. The maximum values of fracture toughness and fracture energy (1.22 MPa. m1/2 and 450 J/m2, respectively) were measured which is an increase of 100 and 350%, respectively, compared with the unmodified epoxy. Carboxyl terminated butadiene acrylonitrile rubber (CTBN), amine terminated butadiene acrylonitrile rubber (ATBN), epoxy terminated butadiene acrylonitrile (ETBN) and hydroxyl terminated poly butadiene liquid rubber (HTPB) have been used to enhance the fracture toughness of epoxy resins at room temperature (RT) where micro- sized rubber-modified epoxy resins showed enhanced fracture toughness with lowered strength[34-51]. SEM photograph of a fractured surfaces of 10 phr rubbermodified epoxy (Figure 5) showed stress-whitened zone near the crack tip occurred due to micro-cavitations of rubber particles because of high hydrostatic stress beneath the blend crack tip[51]. Rubber modifiers are also used in the forms of reactive oligomers (e.g., carboxyl terminated butadiene acrylonitrile CTBN, amine terminated butadiene acrylonitrile ATBN, etc)[52] enhance the fracture toughness of epoxies. the rubber nanoparticles in the form of spherical and wormlike micelles increase KIC monotonically. The microsize rubber particles (CTBN and ATBN) may impart higher, similar or lower KIC values in epoxies depending on the nature of epoxy and the rubber modifier, concentration of rubbery phase, and interface properties[52]. Summarising, the toughening effects of rubber particles on epoxy resins, it is concluded that the micro-size and nano rubber particles (CTBN and ATBN) impart higher toughness in epoxies. Table 2 Summaries the published reports on rubber particle-filled epoxies to evaluate the influence of these fillers on increasing fracture toughness.

4. Overview of Fracture Toughness Evaluation of Epoxy Resin Modified with Carbon Nano Tubes (CNTs) Figure 4. Dispersion of NR particles in the NR-epoxy blends[13].

Figure 5. SEM micrograph of rubber-modified epoxy showing stress-whitened zone[51]. Polímeros, 30(3), e2020030, 2020

Carbon nano tubes (CNTs) are generally considered as one of the potential fillers to improve the mechanical properties of polymer matrices[53-55]. Ayatollahi et al.[53] had investigated the effects of MWCNT as nanofillers on epoxy matrix under bending and shear loading conditions. Shear loading was found more effective in comparison to normal loading. Single-edge notch bend specimen (SENB) was used for this study. Fracture toughness increased with 0.1 wt.% to 0.5 wt.% MWCNT addition resulting in 30% higher than that of the neat epoxy. Hsieh et al.[54] used Multi-walled carbon nanotubes, with a typical length of 140 μm and a diameter of 120 nm, to modify anhydride-cured epoxy polymer. TOM images revealed that the MWCNTs were agglomerated in the epoxy polymer at all the concentrations employed (Figure 6). The modulus, fracture energy and the fatigue performance of the modified polymers have been investigated. The addition of nanotubes increased the modulus of the epoxy, fracture energy from 133 to 223 J/m2 and threshold strain energy 5/14

Cozza, R. C., & Verma, V. Table 2. Summary on Rubber/Epoxy composite. Type of Epoxy

Modifier rubber

DGEBA

Amine terminated butadiene acrylonitrile rubber (ATBN) Carboxylic nitrilebutadiene nano (M22N, M52N, M52)* Epoxy terminated butadiene acrylonitrile (ETBN) & ATBN Hydroxyl terminated polybutadiene liquid rubber (HTPB)

DGEBF DGEBA DGEBA

DGEBA

Rubber Addition (phr) 5-20

Fracture toughness (MPa.m1/2) 0.91 to 1.49

15 12 wt.%

48.3% increment at 77K 1.22

0-15

1-1.65

0-20

1.5

Curing Agent

Ref.

Polyaminoimidazoline

Chikhi et al.[31]

Diethyl toluene diamine (DETD) Methylhexahydrophthalic acid anhydride Trimethylene glycol di-paminobenzoate (TMAB)

Zhao et al.[32]

Cyclic anhydride (Hy906)

Chen and Taylor[33] Hwang et al.[47]

Thomas et al.[51]

*Poly(methyl methacrylate)-bloc-poly(butylacrylate)-bloc-poly(methyl methacrylate) (PMMA–b-PbuA–b-PMMA) BCPs (M22N, M52 and M52N).

to debonding between MWCNTs and the matrix followed by plastic deformation can also be observed – Figure 7.

Figure 6. TOM images of the nanocomposites containing 0.5 wt.% MWCNT[15].

Figure 7. FEG-SEM image of the nanocomposite containing 0.5 wt.% MWCNT showing nanotube bridging[15].

release rate Gth, increased from 24 to 73 J/m2 with 0.5 wt.% nanotubes addition. Electron microscopy of the fracture surfaces showed clear evidence of nanotube debonding and pull-out with void growth around the nanotubes, in both the fracture and fatigue tests. Examination of the fracture surfaces showed nano tube bridging and voids due 6/14

Gkikas et al.[55] studied the effect of dispersion on MWCNT toughened epoxy. Single edge notch 3-point bending (SENB) was used to determine the toughness of the nanocomposite. At 0.5 wt.% CNT reinforcement and the sonication power at full amplitude (100%) for 1 h increased the toughness by 95%. Further sonication for 2 h revealed a reduction in the fracture properties of the nanocomposite. The importance of the duration and amplitude of the sonication process for good dispersibility was highlighted. For nanocomposite with the same CNT loading, reducing the sonication power to half of the maximum amplitude (50%) and increasing the time to 2 h, increased the fracture toughness by 550% as compared to the neat epoxy. The investigated results on the mechanical and fracture toughness properties reported by various researchers, the properties of polymer composites are enhanced at very low CNTs loading. Most of the loading values gathered from the reviewed papers are approximately in the same range. Amongst all the CNTs, DWCNT had produced the highest stiffness nanocomposites when surface functionalization was carried out. The values of young’s modulus presented for most of the MWCNTs nanocomposites were much lower in comparison with SWCNTs and DWCNTs ones. This effect was explained due to the difference in specific surface area (SSA). Optimizing the sonication power during mixing resulted in the highest value of Mode I fracture toughness (1300 J/m2) as reported by Gkikas et al.[55]. Results concerning appropriate type as single-walled CNT (SWCNT), double-walled CNT (DWCNT), or multi-walled CNT (MWCNT) in epoxies showed that CNTs addition increased the tensile properties; hardness, impact resistance, and fracture toughness of epoxy resins as discussed below[37,56-85]. Compressive behavior of epoxy resin system with addition of different kind of nanoparticles revealed that carbon nanotubes increased compressive yield strength modestly at low strain rates, but there is a significant increase at high strain rate[57]. Single walled carbon nanotube (SWCNT) in a polymer matrix increased fracture behavior with increasing weight percentage of SWCNT[58]. Finite-element modelling had demonstrated that once one Polímeros, 30(3), e2020030, 2020

Evaluation of fracture toughness of epoxy polymer composite incorporating micro/nano silica, rubber and CNTs silica nanoparticle debonds then its nearest neighbours are shielded from the applied stress field, and hence may not debond[58]. Covalently functionalized nanotubes have been observed as particularly good additives for polymer reinforcement, allowing for an excellent stress–strain transfer between nanotubes and polymer matrix[60]. Generally, the structural toughness arising from conventional energy dissipation mechanisms has the potential to be much more significant in nanocomposites than in composites containing conventional micron-size fibers[62] whereas uniform distribution of MWCNTs in the matrix and the formation of voids significantly affect the fracture and fatigue behavior of MWCNT-reinforced composites[56,63]. Gojny et al.[66] found that the mechanical properties of potential matrices of fibre-reinforced polymers (FRP), such as epoxy resins, were significantly increased by low contents of carbon nanotubes (CNT) (tensile strength, young’s modulus and fracture toughness). The glass-fibre-reinforced polymers (GFRP) containing 0.3 wt.% amino-functionalised double-wall carbon nanotubes (DWCNT-NH2) exhibited significantly improved matrix-dominated properties (e.g., interlaminar shear strength)[66]. Enhanced dispersion and alignment of CNTs in polymer matrices improved by optimum physical blending, in situ polymerization and chemical functionalization led to improve mechanical, electric, thermal, electrochemical, optical and super-hydrophobic properties[67]. The alignment of CNTs could also be increased by ex situ alignment due to force, electrical and magnetic field-induced methods[67]. It was found that the tensile strength improved with the increase of MWCNTs addition, and as MWCNTs loading reached 8 wt.%, the tensile strength reached the highest value of 69.7 MPa[68]. Park et al.[69] oxyfluorinated multi-walled carbon nanotubes (MWNTs) at different temperatures and it was found that surface fluorine and oxygen contents increased with increasing oxyfluorination temperature which led to an increase of surface polarity of the MWNTs, resulting in increasing KIC and impact strength due to the improvement of interfacial adhesion force. Better dispersibility and stronger interfacial bonding between MWCNTs and epoxy matrix by acid treatment and triethylene-tetramine (TETA) modification of multi-walled carbon nanotubes (MWCNTs). MWCNTs were treated with concentrated H2SO4/HNO3, and then triethylenetetramine (TETA) grafting showed formation of TETA thin layer on the MWCNT surface, which contributes to the homogenous dispersion of MWCNTs in epoxy matrix and the improvement of the MWCNT-epoxy interfacial interaction[70,71]. Amino-functionalized MWNTs/epoxy nanocomposites, in which MWNTs with amino groups acted as a curing agent resulting in improvement of the tensile strength (+51%) and impact strength (+93%) is obtained with amine-treated MWNTs at an 1.5 wt.% content[72] and fracture toughness KIC turned out to be significantly increased (45%) adding only 0.3% of amino-functionalised doublewalled carbon nanotubes (DWCNT-NH2)[73]. It was found that tetra-functional epoxy resin, MY-720 mixed with 1% by weight MWNT increased the fracture toughness of the neat resin by more than three folds in the three-point bend mode[74]. Catalytically grown double-wall carbon nanotubes (DWCNT) were used in epoxy resin and a very good dispersion of both DWCNTs and carbon black (CB) in an epoxy resin could be observed with increased KIC-value was Polímeros, 30(3), e2020030, 2020

observed[75]. Rahmat and Hubert[78] in their review presented carbon nanotube-polymer interactions in nanocomposites and concluded that an optimum carbon nanotube-polymer interaction is a key factor towards reaching the full potential of carbon nanotubes in nanocomposites while presence of moisture absorption may cause degradation resulting in weak interfacial bonding due to epoxy swelling[79] and it is known that nano-composite toughness increased with enhanced interfacial adhesion[80]. Multi-walled carbon nanotube reinforced epoxy nanocomposites, with multiwalled carbon nanotubes (MWCNTs) to diglycidyl ether of bisphenol-F epoxy via the ultrasonic technique resulted in a higher reinforcing efficiency and consequently, the cryogenic tensile strength, Young’s modulus, failure strain and impact strength at 77 K are all enhanced[37]. Nanotube morphology and impurity content can significantly affect the effective properties of the resulting composite as found by Hernández-Pérez et al.[81]. Nanotubes with different aspect ratios showed very limited improvement in tensile properties, where the impact resistance and fracture toughness of the nanocomposites were significantly improved, for the composites employing the nanotubes with higher aspect ratio[81]. Controlling curing process[82], preventing agglomerate of nanotubes as large clumps of black powder[83], purification of single-wall carbon nanotubes[84] also help in proper dispersion resulting in enhanced properties of epoxy composites. The nanocomposites produced exhibited an enhanced strength and stiffness and significant increase in fracture toughness (43%) at 0.5 wt.% amino-functionalised DWCNT addition[85]. The outcome of multiphase study revealed great potential route that could be explored as established with the presence of liquid rubber. The addition of rubber increased Mode I fracture toughness by 409% in comparison with pristine epoxy. MWCNT is still the most interesting candidate with a promising outlook, by choosing an appropriate dispersion and functionalization techniques. Table 3 Summaries the published reports on CNTs-filled epoxies to evaluate the influence of these fillers on increasing fracture toughness.

5. Hybrid Polymers Composites Hybrid polymer composites (HPC) are one of the recent developments to reduce the cost of expensive composites containing reinforcements like carbon fiber by incorporating a proportion of cheaper, low-quality fibers such as glass, textile, natural fibers, and nano figures like silica, rubber, CNT, clay, graphene. The concept of hybrid toughening refers to the use of two or more toughening agents to achieve some synergistic effect in the toughness of the overall nanocomposites[4]. Gouda et al.[4] investigated the fracture behavior of a high silica glass-satin textile fiber reinforced hybrid polymer composite (HPC) under the full range of in-plane loading conditions experimentally and numerically. The observation of fracture surface shows that the Mode I fracture surface is of brittle matrix cracking with relatively smooth flat matrix fracture and showed little debonding between fiber and matrix. Crack specimens are tough in tensile loading condition and weak in shear loading conditions. Tsai et al.[88] investigated the interlaminar fracture toughness of glass fiber/epoxy composites having silica nanoparticles 7/14

8/14

– – – ~ 15

~ 15

– *f-SWCNT – *p-MWCNT

*f-MWCNT

Epoxy 862

1-10

– – – 1-10

140 > 500 nm – –

– 10 & 30

Size length (µm)

*p and f-MWCNTs were used to signify the pristine and ozone-functionalized MWCNTs.

DGEBA

120 10 to 15 – –

– 30 & 40

MWCNT MWCNT – (SW/DW/MW)CNT

DGEBF

DGEBA LY 564 DGEBA

Size diameter (nm)

Name of the modifier element – MWCNT

Type of Epoxy

Table 3. Summary on CNTs/Epoxy composite.

0.5-1

– 1 – 0.5-1

Addition content (wt.%) – 0.1 0.5 1 0.5 0.5-1 – 0.3-0.5 0.79 ± 0.048 0.56 ± 0.03 0.66 ± 0.04 0.46 ± 0.04 0.59 ± 0.05 0.60 ± 0.08 0.64 ± 0.07 0.70 ± 0.06

Fracture Toughness (MPa.m1/2) 1.619 ± 0.14 1.865 ± 0.07 2.045 ± 0.1 1.93 ± 0.19 0.69 to 0.98 2.5 0.65 ± 0.062 0.72 ± 0.014 –

135, 145

– – 64 110, 108

133 to 233 1200 – –

Fracture Energy (J/m2) –

4methylhexahydrophtahlic anhydride (MHHPA) HE600

Epicure W

HE 600 anhydride – Amine hardener H137i

Polyamine HA-11

Curing agent

Tang et al.[87]

Sun et al.[86]

Hsieh et al.[54] Gkikas et al.[55] Gojny et al.[85]

Ayatollahi et al.[53]

Ref.

Cozza, R. C., & Verma, V.

Polímeros, 30(3), e2020030, 2020

Evaluation of fracture toughness of epoxy polymer composite incorporating micro/nano silica, rubber and CNTs and the rubber particles. Reactive liquid rubber (CTBN) and core-shell rubber (CSR), were employed to modify the fracture toughness of epoxy resin. Silica nano particle increased the fracture toughness of the fiber composite with pure or CTBN-modified epoxy. The silica nanoparticle effect would become detrimental when the fiber composites are already modified by CSR. By considering the overall mechanical performances, the fiber composite with silica nanocomposites together with CTBN rubber particles demonstrate superior properties. Hsieh et al.[89] presented the structure/property relationship of an anhydride-cured epoxy modified with silica nanoparticles and/or a rubber microparticle. Transmission electron micrographs of ‘hybrid’ epoxy polymers containing 2.3 wt.% silica nanoparticles and 9 wt.% CTBN is shown in Figure 8. The fracture energy of the bulk epoxy increased from 77 to 212 J/m2 with 20 wt.% silica nanoparticles addition. The observed toughening mechanisms were debonding of the epoxy polymer from the silica nanoparticles, followed by plastic void growth of the epoxy. Localised plastic shear-banding in the polymer was also observed. Maximum fracture energy of 965 J/m2 was measured for a ‘hybrid’ epoxy polymer containing 9 and 15 wt.% of the rubber microparticles and silica nanoparticles respectively. Both the theoretical and experimental studies clearly revealed the benefits of using silica nanoparticles, as opposed to much larger micrometre-sized silica particles, in terms of observing a relatively high toughness for the modified epoxy polymer. Although the toughness increase is dependent on the toughener used, its concentration and particle size, the second-phase materials have been shown of increasing the toughness of epoxy polymers. It should be noted that the toughening effect is also dependent on the properties of the epoxy itself, as some epoxies are tougher in nature than others[89]. Figure 9 shows evidence of cavitation of the rubber particles, though it is difficult to identify the mechanisms associated with the silica nanoparticles due to the roughness of the surfaces. However, SEM studies of the ‘hybrid’ fracture surfaces showed that the silica nanoparticles are present as both individual particles and agglomerates[89]. Gouda et al.[90] studied Mode I fracture behavior of glass-carbon fiber reinforced hybrid polymer composite experimentally and by finite element analysis at room temperature with curing hardener (HY 951) and woven glass and carbon bi-direction mesh. The difference in the magnitude of elastic modulus along and across the fiber orientation is 417 MPa. Mode I fracture surface is indicative of brittle cleavage failure with relatively smooth and flat matrix fracture with a little debonding between fiber and matrix. The fracture energy of hybrid carbon fiber reinforced polymers was investigated by Karapappas et al.[91]. The composites were modified by multi-walled carbon nanotubes addition in the matrix material. The interlaminar fracture properties under Mode I and Mode II remote loading were studied as a function of the carbon nanotube content in the matrix. Carbon nanotubes in the epoxy matrix led a significant increase in the load bearing ability as well as in the fracture energy, for both Mode I and Mode II tests. It is speculated that carbon nanotubes due to their large aspect ratio have a significant toughening effect since extra energy is needed in order to pull them out from the matrix and start the crack propagation following a Polímeros, 30(3), e2020030, 2020

Figure 8. TEM image of ‘hybrid’ epoxy polymers containing 2.3 wt.% silica nanoparticles and 9 wt.% CTBN[18].

Figure 9. SEM image of fracture surface of epoxy polymers with 4.5 wt.% silica nanoparticles and 9 wt.% CTBN[18].

kinking out pattern at nano scale. Lee et al.[92] investigated Mode I and Mode II interlaminar fracture toughness of the hybrid laminates with nonwoven carbon tissue (NWCT) under severe temperature conditions, the double cantilever beam (DCB) and end notched flexure (ENF) tests for the carbon fiber reinforced polymer (CFRP) and the hybrid specimens were conducted. The GIC and GIIC values of the hybrid laminates were compared with those of the CFRP laminates in the range of −60 °C to +80 °C. The mean GIC values of the CFRP and hybrid specimens at −60 °C and −30 °C were not changed significantly when compared with those the CFRP and hybrid specimens at room temperature. The mean GIIC values of the hybrid specimens at −60 °C to +80 °C were about 171% to 189% higher than those of the CFRP specimens at −60 °C to +80 °C, respectively. When compared with the decreasing rate of the mean GIIC values of the CFRP specimens at +80 °C, the decreasing rate of the mean GIIC values of the hybrid specimens at +80 °C slowed down significantly due to the carbon shortfibers bridging, carbon short-fibers breakage and hackles. Borowski et al.[93] investigated the fracture toughness of carbon fiber reinforced polymer (CFRP) laminates produced by using an epoxy nanocomposite matrix reinforced with 9/14

Cozza, R. C., & Verma, V. Table 4. Summary on Hybrid Polymers Composites. Name of the modifier element High silica glass fiber Woven Satin textile fiber Type of Fiber

*GFRP

% of content addition (wt.%) t = 0.26 mm

GFRP GFRP

– SiO2-CTBN SiO2-CSR Rubber (µm) – SiO2 (nm) Carbon

– 10 wt.% each 10 wt.% each 9 wt.%, 15 wt.% 0/90

*CFRP

MWCNT

CFRP

MWCNT

0.1 0.5 1 – 0.5 1 1.5

% of KIC, MPa.m1/2 increment 7.431 [0°] 4.199 [90°] – – – – 34.07 (Across) 32.84 (Along) – – – – – – –

% of GIC, J/m2 increment – 190 930 1030 (390%) 965 – 200-300 400-500 400-600 943 (442%) 1175 (25%) 1132 (20%) 1102 (17%)

Ref. Gouda et al.[4] Tsai et al.[88]

Hsieh et al.[89] Gouda et al.[90] Karapappas et al.[91]

Borowski et al.[93]

*Glass fibre reinforced-polymer (GFRP) and carbon-fibre reinforced-polymer (CFRP).

carboxyl functionalized multi-walled carbon nanotubes (COOH–MWCNTs). Four MWCNTs contents of 0.0 – 1.5% per weight of the epoxy resin/hardener mixture were examined. Double cantilever beam (DCB) tests performed to determine the Mode I interlaminar fracture toughness of the unidirectional CFRP composites showed 25%, 20%, and 17% increase in the maximum interlaminar fracture toughness of the CFRP composites. The involvement of silica and rubber particles in glass fiber epoxy has shown a major improvement in Mode I fracture toughness with lowered stiffness due to low modulus rubber particles as highlighted by Tsai et al.[88]. Study conducted by Hsieh et al. [89] reported maximum fracture energy of 965 J/m2 for a hybrid polymer composite containing 9 wt. % and 15 wt. % of the rubber microparticles and silica nanoparticles, respectively. Karapappas et al.[91] and Borowski et al.[93] has reported the increase in fracture toughness of hybrid carbon fiber reinforced polymers for low loading of Multi-walled carbon nanotubes. It was found that toughness of epoxy can be significantly increased by incorporating either rubber or silica nano-particles, however, hybrid nanocomposites do not display any significant effect on toughness[94]. Carolan et al.[95] observed increase in toughness in the bulk epoxy polymer by the addition of a combination of silica nanoparticles and/or CSR nanoparticles. Rubbery and glassy epoxy resins reinforced with carbon nanotubes created a bridge between the nanotubes and matrix, and differences in viscosity resulted better nanotubes dispersion in the rubbery epoxy resin than in glassy epoxy with 28% increase in tensile Young’s modulus in the rubbery system using 1 wt.% functionalized nanotubes[76]. The values of toughness of the CFRP laminates, compared to the bulk epoxy polymer, were further enhanced by additional fibre-based toughening mechanisms, i.e., fibre bridging, fibre debonding and fibre pull-out. It was also observed that epoxy/acryl triblock copolymer alloys, applied as the toughening modifiers for the epoxy resins, rubbery epoxy particles successfully acted as toughening agent for glassy epoxy matrices[96-98], effects of both rubbery phase and nanosilica on mechanical properties of epoxy showed that fracture surfaces were accompanied 10/14

with multiple voids, providing evidence of debonding of the nanoparticles[99]. It has been well established that the incorporation of second soft microphase such as rubber particles into epoxy polymers increase their toughness, without significant impairing other properties whereas addition of nanosilica particles to rubber-toughened epoxy may lead to very significant increase in the toughness of the matrix[100]. Summarising, the effect of having multiphase in polymer composite is highly beneficial in order to attain good mechanical and fracture properties of polymer composite. Table 4 presents the summary of published reports on hybrid polymers composites. It can be concluded from the present study that addition of a second dispersed phase (silica, rubber, CNTs) in epoxies results in enhancing the material fracture toughness and inducing a remarkable increase of damage tolerance performance.

6. Conclusions In the research conducted so far, the fracture toughness of the epoxy based polymer matrix composite incorporating silica, rubber and CNTs in micro/nano sizes is and attempt is made to understand the controlling toughening mechanisms of these materials. Following are the major conclusions with proposed future direction: (1) Silica addition promoted nano toughening effect with plastically deformation capability in epoxies. Addition of rubber increased the toughness with negligible reduction in stiffness in epoxies. Double and multi walled CNTs addition significantly increased the tensile properties; hardness, impact resistance and fracture toughness of epoxy resins; (2) Summarising, the effect of having multiphase in polymer composite is highly beneficial in order to attain good mechanical and fracture properties of polymer composite; (3) The advantages of using nanofillers modified epoxies as matrices or interleafs for carbon fiber and glass fiber Polímeros, 30(3), e2020030, 2020

Evaluation of fracture toughness of epoxy polymer composite incorporating micro/nano silica, rubber and CNTs composites to increase the interlaminar delamination toughness are not fully explored for better efficiency than some conventional methods; (4) More emphasis can be laid on the crucial understanding of stress transfer mechanisms and interfacial bond strength between nanoparticles and the epoxy system and the fracture toughness of hybrid polymer composite under different temperature conditions needs to be elaborated; (5) The experiments conducted so far to measure different modes of fracture toughness of hybrid nanocomposites may be numerical validated (modeling); (6) In CNT modified epoxies, further studies are required to clarify the predominant toughening mechanisms with respect to the different CNT morphology although crack bridging is important.

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