Polímeros: Ciência e Tecnologia 3rd. issue, vol. 31, 2021

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Polímeros VOLUME XXXI - Issue III - July/Sep., 2021

30 years of Polímeros Three Decades Sharing Polymer Science

São Paulo 994 St. São Carlos, SP, Brazil, 13560-340 Phone: +55 16 3374-3949 Email: abpol@abpol.org.br 2021

DESCUBRA o conjunto de instrumentos que conduzem a percepções mais profundas sobre as PROPRIEDADES e ESTRUTURA DO POLÍMERO em cada etapa



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Three Decades Sharing Polymer Science Sebastião V. Canevarolo  Editor-in-Chief, Departamento de Engenharia de Materiais, Universidade Federal de São Carlos – UFSCar, São Carlos, SP, Brasil How to cite: Canevarolo, S. V. (2021). Three Decades Sharing Polymer Science. Polímeros: Ciência e Tecnologia, 31(3), e2021037. https://doi.org/10.1590/0104-1428.ED3103

30 years In Nov/2021, Polímeros journal completes three decades disseminating scientific knowledge in the polymer area! During this period, it passed through different moments, in the beginning it was the learning how to

manage a scientific publication, then being accepted by the community interested in the “Plastic World”, following the insecurity of opening up to the world by publishing only in English, and quite recently finding itself forced to apply the Publication Article Charge PAC, in order to survive. These have been just some of the enormous challenges the journal had to face during its first 30 years. But all in all, the result has been positive, justified by the gradual and constant increase in its impact factor, the last one an impressive ~50% increase to 1.492 (2020). This was only possible thanks to the generous contribution of all, authors, reviewers, members of the Editorial Council and Committee, collaborators, and the entire “Polymer Community”, to which we sincerely thank. But the world paces in continuous evolution, new challenges present themselves, now with the inevitable implementation of the “Open Science”[1,2], when the journal again finds itself face to face with the new rules of scientific dissemination, generating doubts on how well to disseminate them to authors and reviewers, and gradually implement them. Another huge challenge is to find scientifically and economically viable ways to recycle and reuse discarded plastic products, especially films and bottles, which ends up in the ocean as microplastics[3], consequence of the eager seek of the world’s population to increase their well-being, direct result of their enrichment. Let 2022 be not just the beginning of a new year, but the beginning of a new decade for Polímeros, full of goals to be achieved, challenges to be overcome, and above all, to remain a reliable and scientifically sound source of open access polymer information to everyone.

References 1. PLOS. Retrieved in 2021, December 22, from https://plos.org/ 2. SciELO. Retrieved in 2021, December 22, from https://scielo.org/en/ 3. Isobe, A., Iwasaki, S., Uchida, K., & Tokai, T. (2019). Abundance of non-conservative microplastics in the upper ocean from 1957 to 2066. Nature Communications, 10(1), 417. http://dx.doi.org/10.1038/s41467-019-08316-9. PMid:30679437.

Polímeros, 31(3), e2021037, 2021



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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 I - 2 0 2 1 Indexed


“ 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 ­T e c h n ic a l I n f o r m a t i o n ” — “ L a t i n d e x ” — “ W e b o f S ci e n c e ”


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 Ailton S. Gomes (UFRJ/IMA), Rio de Janeiro, RJ (in memoriam) Alain Dufresne (Grenoble INP/Pagora) Bluma G. Soares (UFRJ/IMA) César Liberato Petzhold (UFRGS/IQ) Cristina T. Andrade (UFRJ/IQ) 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 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


“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 2021

Financial support:

Available online at: www.scielo.br

Polímeros / Associação Brasileira de Polímeros. vol. 1, nº 1 (1991) -.- São Carlos: ABPol, 1991Quarterly v. 31, nº 3 (July/Sept. 2021) ISSN 0104-1428 ISSN 1678-5169 (electronic version)

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

1. Polímeros. l. Associação Brasileira de Polímeros. E2

Polímeros, 31(3), 2021

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

O r i g in a l A r t ic l e Assessment of derived sunflower oil as environmentally friendly plasticizers in Poly Vinyl Chloride Boussaha Bouchoul and Mohamed Tahar Benaniba....................................................................................................................................... 1-9

Epoxy hybrid composites reinforced with nanodiamond-silica for abrasive applications Marcos Antônio Barcelos, Mariana Valinhos Barcelos, Gabriel Rodrigues de Almeida Neto, Antônio Cesar Bozzi and Rubén Jesus Sánchez Rodriguez .................................................................................................................................................................... 1-10

Esterification of oleic acid employing sulfonated polystyrene and polysulfone membranes as catalysts Ana Paula de Lima, Andressa Tirone Vieira, Bárbara Nascimento Aud, Antonio Carlos Ferreira Batista, Luís Carlos de Morais, Anízio Márcio de Faria, Rosana Maria Nascimento de Assunção and Daniel Pasquini ............................................................................... 1-8

ABS/recycled PCTG blend compatibilized with ionomer: effect on impact resistance and morphology Juliana Augusto Molari and Deborah Dibbern Brunelli ................................................................................................................................ 1-8

Structural changes of polyethylene in blown films with different pro-oxidants João Augusto Osório Brandão, Fernando Dal Pont Morisso, Edson Luiz Francisquetti and Ruth Marlene Campomanes Santana ............ 1-8

Antimicrobial activity of silver composites obtained from crosslinked polystyrene with polyHIPE structures Roberta Trovão Santos, Nathália Smith Santos, Mirian Araújo de Oliveira, Fernanda de Andrade Buás Campeão, Maria Aparecida Larrubia Granado Moreira Rodrigues Mandu, Mônica Regina Costa Marques and Luciana da Cunha Costa ............ 1-10

Effects of gamma radiation on nanocomposite films of polycaprolactone with modified MCM-48 Marcos Vinícius Paula, Leandro Araújo de Azevedo, Ivo Diego de Lima Silva, Glória Maria Vinhas, Severino Alves Junior .................... 1-10

Effect of drying different inclusion plasters on the mechanical properties of thermoactivated acrylic resins Tarcisio José de Arruda Paes-Junior, Natália Rivoli Rossi, Tayná Mendes Inácio de Carvalho, Vanessa Cruz Macedo, Michelle de Sá dos Santos Gomes, Leonardo Jiro Nomura Nakano and Cristiane Mayumi Inagati.............................................................. 1-6

ABS/Recycled PCTG blend compatibilized with SBS: effect on mechanical properties and morphology Juliana Augusto Molari, Deborah Dibbern Brunelli ....................................................................................................................................... 1-8

Role of cellulose nanocrystals in epoxy-based nanocomposites: mechanical properties, morphology and thermal behavior Nayra Reis do Nascimento, Ivanei Ferreira Pinheiro, Guilherme Fioravanti Alves, Lucia Helena Innocentini Mei, José Costa de Macedo Neto and Ana Rita Morales ...................................................................................................................................... 1-13

Structural and optical properties o plasma-deposited a-C:H:Si:O:N filmsa Juliana Feletto Silveira Costa Lopes, Jean Tardelli, Elidiane Cipriano Rangel and Steven Frederick Durrant ............................................ 1-8

Incorporation of astrocaryum vulgare (tucuma) oil into PCL electrospun fibers Nathan Rampelotto Bressa, Vinícius Rodrigues Oviedo, Aline Machado Bessow Machado, Willians Lopes de Almeida, Tiago Moreno Volkmer, Luis Alberto Loureiro dos Santos, Michele Rorato Sagrillo and Luiz Fernando Rodrigues Junior ......................... 1-8 Cover: Transmission electron micrographs of cellulose nanocrystals, CNCs Arts by Editora Cubo.

Polímeros, 31(3), 2021




SABIC CREATES THE FIRST CERTIFIED CIRCULAR POLYMERS FROM ADVANCED RECYCLING OF RECOVERED OCEAN-BOUND PLASTIC SABIC, a global leader in the chemical industry, and Malaysia-based plastic recycling company HHI, have announced a pioneering new collaboration to create the first certified circular polymers produced through the advanced recycling of recovered mixed and used ocean-bound plastic. The certified circular polyolefins from ocean-bound plastic, from SABIC’s TRUCIRCLE™ portfolio of circular solutions, will be used by SABIC’s customers to announce new products over the coming months. As well as helping to protect our oceans and waterways, the ocean-bound plastic collection helps to create value for local communities by increasing demand for recycled plastic across the industry. The material is recovered from ocean-feeding waterways and inland areas within a 50 kilometres radius of the ocean by HHI partners predominantly in Malaysia. The recovered material is then sent to HHI, where they convert the used plastic into pyrolysis oil through an advanced recycling. The pyrolysis oil is then used by SABIC in their production process as an alternative to traditional fossil materials to make new certified circular polymers. The material has been certified under the Zero Plastic Oceans accreditation, and HHI is the first organisation to have received certification confirming the materials it recycles qualify as ocean-bound. HHI created its own model to outline the steps required to facilitate the transition to a circular economy. The model has five stages which are to collect ocean-bound plastic through its extensive network; convert them into high-quality, manufacturable materials; collaborate with partners to create new products; provide customers with the platform to champion their use of more sustainable materials; and catalyse a generation of conscientious consumers who will opt for sustainable materials. Abdullah Al-Otaibi, General Manager, ETP & Market Solutions at SABIC, said: “We are acutely aware of the challenges we face globally to stop plastic from becoming waste. Developing an entirely circular recycling system is a huge but necessary step we need to take together and will require all players across the value chain to collaborate. That’s why we’re committed to developing long-term solutions and working with new partners like HHI to significantly upscale the production of more sustainable materials, including those produced using recycled ocean-bound materials, for the benefit of our customers, society and the environment.” Kian Seah, CEO at HHI, explains: “At HHI, our circular economy model helps to guide us in all of our endeavours, from business planning and collaborations with partners such as SABIC, to eco-initiatives, as we strive to protect our ocean and communities. We believe that we have the ability to work towards a cleaner future that views plastic as a valuable resource to keep within the value chain. “We are incredibly proud of what we have achieved so far with SABIC, but also realise that we are early into our journey towards enabling a circular economy, and it is by no means a straightforward one. Our common spirit and passion has helped us overcome significant challenges to make this innovative process a reality and to ensure the reliability of technologies, quality of the end material and viability of the circular consumption model. We share a commitment to reshape the way we produce and recycle plastics and address environmental and societal challenges in a lasting, meaningful way.” The circular polymers produced from ocean-bound plastic form part of SABIC’s TRUCIRCLE™ portfolio and services for circular innovations. Launched in 2019, SABIC’s TRUCIRCLE portfolio spans mechanically recycled products, certified circular products from advanced recycling of used plastic and certified renewables products from bio-based feedstock, as well as design for recyclability and closed loop recycling initiatives. Source: SABIC – www.sabic.com/en


BSP: First batch of polymer ₱1,000 bills to arrive April 2022; Australia to print banknotes The country will get to test its first-ever polymer banknotes when they arrive next year, the Bangko Sentral ng Pilipinas (BSP) announced on Thursday. BSP Governor Benjamin Diokno said the first batch of the more durable bills is expected to arrive in April 2022. The BSP partnered with the Reserve Bank of Australia and its subsidiary Note Printing Australia for the production of the banknotes. “Australia’s the first country to issue full-series polymer banknotes and has produced and supplied polymer banknotes to several countries. As such, their advanced technology and expertise in the printing of polymer banknotes will be the best benchmark for our first circulation,” Diokno said. The central bank has been authorized to produce 500 million pieces of ₱1,000 polymer bills for its pilot - running between 2022 and 2025 - according to BSP Deputy Governor Mamerto Tangonan. Polymer banknotes are said to have a lifespan that is 2.5 to 4 times longer than paper ones, and can be sanitized without getting damaged — a handy feature especially as the COVID-19 pandemic lingers on. Given their longer lifespan, it will take longer to replace them — effectively costing the government less in their production. Source: CNN Philippines - www.cnnphilippines.com

Polímeros, 31(3), 2021

February International Conference on Material Science and Engineering Date: February 24-25, 2022 Location: Prague, Czech Republic Website: materialsscience.conferenceseries.com

March World Congress on Carbon and Advanced Energy Materials Date: March 08-09, 2022 Location: Auckland, New Zealand Website: global.materialsconferences.com 23rd World Congress on Materials Science and Engineering Date: March 09-10, 2022 Location: Barcelona, Spain Website: materialsscience.insightconferences.com

April 37th International Conference of the Polymer Processing Society (PPS-37) Date: April 11-15, 2022, (hybrid) Location: Fukuoka, Japan Website: www.pps-37.org 2nd Annual congress on SMARTMATERIALS 2022 Date: April 21-22, 2022 Location: Florida, United States Website: smartmaterials.euroscicon.com Fire Retardants in Plastics - 2022 Date: April 26-27, 2022 Location: Houston, USA Website: www.ami.international/events/ event?Code=C1189#15909

May Polymers for Fuel Cells, Energy Storage and Conversion Date: May 15-18, 2022 Location: Napa, United States Website: www.polyacs.net/2022fuelcells 7th Annual Conference and Expo on Biomaterials Date: May 18-19, 2022 Location: London, United Kingdom Website: biomaterials.insightconferences.com MOMPS-X — 10th International Symposium on Molecular Order and Mobility in Polymer Systems Date: May 23-27, 2022 Location: Saint Petersburg, Russia Website: momps2020.macro.ru Polymers 2022 — New Trends in Polymer Science: Health of the Planet, Health of the People Date: May 25-27, 2022 Location: Turin, Italy Website: polymers2022.sciforum.net 3rd Annual Congress on Biofuels and Biopolymers Date: May 28 - 29, 2022 Location: Vancouver, Canada Website: biofuels.enggconferences.com

June Fire and Polymers – 2022 Date: June 5-8, 2022 Location: Napa, United States Website: www.polyacs.net/22fipo Chemical Recycling - 2022 Date: June 15-16, 2022 Location: Cologne, Germany Website: www.ami.international/events/ event?Code=C1185#15657 Polymer Sourcing & Distribution – 2022 Date: June 28-30, 2022 Location: Hamburg, Germany

Website: www.ami.international/events/event?Code=C1186 2nd Global Conference on Advances in Polymer Science and Nanotechnology Date: June 27-28, 2022 Location: Berlin, Germany Website: polymerscience.peersalleyconferences.com EPF – European Polymer Congress Date: June 26 - July 1, 2022 Location: Prague, Czech Republic Website: www.epf2022.org

July 49th World Polymer Congress – MACRO2022 Date: July 17-21, 2022 Location: Winnipeg, Canada Website: www.macro2022.org PVC Formulation Asia - 2022 Date: July 19-20, 2022 Location: Bangkok, Thailand Website: www.ami.international/events/event?Code=C1178 84th Prague Meeting on Macromolecules – Frontiers of Polymer Colloids Date: July 24-28, 2022 Location: Prague, Czech Republic Website: www.imc.cas.cz/sympo/84pmm/ 8th International Conference on Chemical and Polymer Engineering (ICCPE’22) Date: July 31 - August 2, 2022, (hybrid) Location: Prague, Czech Republic Website: cpeconference.com

August The Global Meet on Bio-Polymers and Polymer Science (GMBPPS2022) Date: August 25-27, 2022 Location: Paris, France Website: primemeetings.org/2022/polymer-science

September Polymer Physics Meeting — Retirement Conference for Dame Athene Donald Date: September 12-13, 2022 Location: Cambridge - United Kingdom Website: events.iop.org/polymer-physics-meeting-retirementconference-dame-athene-donald Advances in Polyolefins (APO-22) Date: September 18-21, 2022 Location: Rohnert Park (Northern California), United States Website: www.polyacs.net/22apo

October Polymers and Nanotechnology Date: October 16 - 19, 2022 Location: Napa, United States Website: www.polyacs.net/21polynano, Polymers in Medicine and Biology Date: October 23 - 26, 2022 Location: Napa, United States Website: www.polyacs.net/21polynano

2023 March 18th International Plastics and Petrochemicals Trade Exhibitions Date: March 21-24, 2023 Location: Riyadh, Saudi Arabia Website: saudi-pppp.com/saudi-plastics-petrochem

Polímeros, 31(3), 2021 E5


ABPol Associates Sponsoring Partners

Polímeros, 31(3), 2021


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Assessment of derived sunflower oil as environmentally friendly plasticizers in Poly Vinyl Chloride Boussaha Bouchoul1*  and Mohamed Tahar Benaniba2 Research Center in Industrial Technologies CRTI, P.O. Box 64 Cheraga 16014, Algiers, Algeria Laboratory of Multiphase Polymeric Materials (LMPMP), Faculty of Technology, Setif-1 University, 19000 Setif, Algeria 1


*b.bouchoul@crti.dz; b_bouchoul@yahoo.com

Abstract Epoxidized sunflower oil (ESO) and epoxidized sunflower oil methyl ester (ESOME) were synthetized and employed as secondary bio based plasticizers and combined with primary natural plasticizers. As such, di-esters isosorbide (DEI) and acetyl tributyl citrate (ATBC) used in Poly Vinyl Chloride (PVC) as compared to di-(2-ethylhexyl) phthalate (DEHP). PVC Sheets were obtained by processed the polymer and its additives on two-roll mill. The effect of the plasticizers combination on flexible sheets properties, including stabilization, migration, fusion rheological properties and light transmission have been discussed, Results have shown that ESO and ESOME have improved PVC formulations thermal stability, the weight loss by volatility has decreased in the formulations which contained ESO or ESOME. Mixtures of plasticizers have lower migration rates. Plastograph test has indicated a better compatibility of PVC with these plasticizer mixtures. The light transmission rate through PVC sheets plasticized by bio-based plasticizers mixtures is more than those containing DEHP. Keywords: biobased plasticizers, migration, PVC, plasticizers combinations, sunflower oil. How to cite: Bouchoul, B., & Benaniba, M. T. (2021). Assessment of derived sunflower oil as environmentally friendly plasticizers in Poly Vinyl Chloride. Polímeros: Ciência e Tecnologia, 31(3), e2021025. https://doi.org/10.1590/01041428.20210015

1. Introduction Polyvinyl chloride (PVC) has a wide range of applications in various fields that have several advantages such as product performance, material processing capability, thermal stability, relatively low cost and versatility. Generally, it has been used with various additives such as thermal stabilizers, plasticizers, lubricants, fillers …etc, according to the application or depending on desired properties[1,2]. Plasticizers are the most consumable additives in PVC[3]. These materials are mainly used to modify the polymer chain flexibility, the glass transition temperature (Tg), degree of crystallinity, optical transparency and Young modulus[3-5]. The most widely used PVC plasticizer in the world is di(2-ethylhexyl) phthalate (DEHP) as a primary plasticizer. Plasticizers based on phthalates account for more than 90% of the total plasticizers used in PVC which DEHP represents more than 50% of the total phthalates[6-9]. Unfortunately, this phthalates migrate easily from PVC formulations and reduce the performance of PVC products with time[10]. Yet, petroleum-based plasticizers reveal negative environmental and health effects[11-13]. They would result in possible toxicity and high danger to contaminate environment and human food, as they are also suspected of having carcinogenic effects[1,5,14]. These have led researchers[15-19] to be interested in substituting these polluting substances with bio-based products that have no negative impact on the environment or

Polímeros, 31(3), e2021025, 2021

on human health[11] by using some derivatives of vegetable oils, such as, soybean oil, linseed oil and sunflower oil, and to make these vegetable oils compatible with PVC, they have subjected to an epoxidation reaction where an oxygen atom introduced into the doubles bonds of carbon chains of fatty acids forming an oxirane ring, where the resulting properties are dependent on the epoxidation level which is linked with the iodine number[20]. Epoxidized soybean oil has been widely used in PVC as a secondary plasticizer, but it shows partial compatibility with PVC compared to petroleum based plasticizers[21]. But with its esterification by an alcohol improves the plasticizing effect on PVC and the resistance to migration against DEHP[22]. For this, we use sunflower oil, which has a high content of non-conjugated double bonds that are simply epoxidized by oxygen peroxide to produce epoxidized sunflower oil (ESO), and it shows a plasticizing effect higher than that of epoxidized soybean oil (ESBO)[23]. Then ESO is esterified by methanol to obtain epoxidized sunflower oil methyl ester (ESOME)[17]. These are used as secondary plasticizers to improve the thermal stability of PVC and increase the compatibility[24]. The use of a plasticizer alone does not satisfy all the desired performances; therefore, it is possible to use a combination of plasticizers in order to meet more properties of the PVC according to a particular application.



Bouchoul, B., & Benaniba, M. T. The present study evaluates the performance of biobased plasticizers combinations in polyvinyl chloride (PVC). The secondary natural plasticizers were synthesized in our laboratory by epoxidation reaction of sunflower oil and esterification of epoxidized sunflower oil; these products were combined with primary bio-based plasticizers which are di-ester isosorbide (DEI) and acetyl tributyl citrate (ATBC).The formulations are prepared by mixing plasticizers combination and other additives with PVC resin which one contained DEHP alone used as reference. Then, blends aremixed and heated; sheets of plasticized PVC are obtained. Several techniques such as thermogravimetric analysis (TGA) in isothermal mode, volatility and migration tests, internal mixer test and light transmission have been used to study the thermal, physicochemical and rheological properties in order to search for a synergism of properties between these combinations of plasticizers.

• 60 phr for the plasticizer system containing DEI, ATBC, ESO and/or ESOME;

The various plasticized formulations are obtained either in the presence of plasticizer alone, DEI or ATBC or with their binary combinations with ESO and ESOME, in addition to a reference formulation which contains DEHP alone as plasticizer.

2.3 Thermal gravimetric analysis (TGA) TGA was performed using a TA Instruments (TGA Q500, USA). Isothermal weight loss studies were conducted at 160 °C, 180 °C and 200 °C for 90 minutes under 40 mL/min of nitrogen flow. Approximately 10 ± 2 mg of each sample is heated from room temperature to the selected temperature with a constant heating rate of 100 °C/min. The mass evolutions are recorded as a function of time.

2. Materials and Methods

2.4 Volatility test

Polyvinyl chloride (PVC), suspension grade resin (SE 950, K 65.7-67.1), was kindly supplied by SHINTECH (Houston, USA). Plasticizers used were as follows: Acetyl tributyl citrate (ATBC) from Sigma Aldrich, USA, Diesters isosorbide(DEI), ID47 from Roquette Frères, France, Epoxidized sunflower oil (ESO) and Epoxidized sunflower oil methyl ester (ESOME) with 6.1% of oxirane oxygen index, were prepared in our laboratory and have been reported in our previous study[17], Di-(2-ethylhexyl) phthalate (DEHP) as plasticizer reference was supply from Plastimed, Tunisia.

The weight loss of plasticizer from PVC sheets by volatility was determined according to ISO 176-2005. The samples (25x25 mm2 in triplicate) are weighed and then placed in an oven at 100 °C. After 24, 48 and 96 hours, sheets were allowed to cool down to room temperature and are cleaned well to remove any traces of volatile plasticizers, which may condense on the sheet surfaces. The masses were measured by an analytical balance having an accuracy of 0.1 mg and the average percentage of losses by volatility was determined.

2.1 PVC sheets preparation

2.5 Migration test

The preparation of the PVC with plasticizers mixtures was carried out in a two-roll mill of the Rodolfo Comero type. It is a method of mixing the melt material by crushing between two heated and mechanically driven rolls. Blends were obtained by manually mixing the PVC resin and additives (plasticizers, stabilizers and lubricants), and then they were placed in the mixer at 160 °C for 10 minutes until the resin was completely homogenized.

Plasticizer migration from PVC sheets was based on ASTM D1239-14[25]. When was carried out in four different solvents (acetic acid, ethanol, petroleum ether and n-heptane) at 25 °C[26]. Samples of 25×25 mm2 were used in triplicate runs to confirm repeatability, were weighed and then immersed in 150 mL of each solvent for 48 and 96 hours, and then dried. Sheets were weighed and the average mass losses were determined.

2.2 Formulations

2.6 Plastograph test

Table 1 presents the various formulations in which the additive contents were kept constant:

The plastograph is an internal mixer that tracks rotational torque as a function of time. This equipment (Brabender, Germany) allows continuously to monitoring the torque variation as a function of time. Accurately 63 g of the PVC and its plasticizers combinations were added into the mixer chamber at 40 rpm and heated to 180 °C for 40 min, a pressure is applied on the melted material by descent of a piston; the fusion rheological property was measured and recorded by Brabender® Data Correlation software.

• 100 phr of the PVC resin; • 1 phr for the lubricant which is stearic acid (SA); • 2 phr for the thermal stabilizer which is Calcium / Zinc stearates; Table 1. Plasticizer systems compositions. Plasticizers (%) DEI ATBC ESO ESOME DEHP


Formulation Number 0 1 2 3 4 5 - 100 50 00 50 00 50 100 50 50 50 100 -

2.7 Light transmission 6 50 50 -

7 50 50 -

Light transmission in PVC sheets was measured using a UV spectrophotometer (SHIMADZU UV-1800, Japan). The samples of a square shape of sizes 2.5×2.5 cm2 and 0.5 mm thick had substantially flat-to-parallel surfaces free of dust and internal voids[26]. The results were recorded in the wavelength range 200 to 1100 nm at a resolution of 1 nm. Polímeros, 31(3), e2021025, 2021

Assessment of derived sunflower oil as environmentally friendly plasticizers in Poly Vinyl Chloride Transmission percentages were obtained directly from the UV Probe software.

3. Results and Discussions 3.1 Thermal gravimetric analysis (TGA) Figure 1 shows the thermogravimetric profiles of PVC plasticized with DEI or ATBC and its mixtures with ESO

or ESOME at 160 °C, 180 °C and 200 °C. PVC samples containing ESO or ESOME in their plasticizer systems show a lower weight loss after 90 minutes as compared to DEI or ATBC alone. Thus, ESO and ESOME improve the thermal stability of PVC formulations. This effect is attributed to the epoxide functions presented in the structure of secondary plasticizers[11]. It is clear that oxirane rings react with heat-generated HCl at a rate greater than the rate of volatilization of HCl from the PVC sample. The point

Figure 1. Weight loss by TGA in isothermal mode of PVC plasticized with DEI, ATBC and their mixtures with ESO, ESOME and with DEHP as reference. Polímeros, 31(3), e2021025, 2021


Bouchoul, B., & Benaniba, M. T. at which the volatilization of HCl dominates the reactions which cannot be determined during the 90 minutes period[11]. Therefore, maintaining the integrity of the ESO or ESOME as a secondary plasticizer is primordial because of epoxide rings and ester functions that limit the decomposition of PVC samples[10,27]. Activation energies for the PVC degradation process were determined by TGA for PVC plasticized by DEHP and PVC with ATBC or DEI and its mixtures with ESO or ESOME according to the Arrhenius equation: k = Ae− Ea / RT (1)

where k is the degradation rate at a particular temperature, A is the frequency factor, Ea is the activation energy for the degradation process, R is the gas constant, and T is the absolute temperature. The isothermal weight loss data of PVC plasticized with different plasticizers systems at various temperatures a long with the Arrhenius activation energy are presented in Table 2. The plots for each sample are shown in Figure 2. The activation energy for the initial degradation for PVC plasticized with DEHP alone is 81.79 ±9.63 kJ mol-1 while that for PVC with ATBC and DEI alone are 65.55±0.96 kJ.mol-1 and 70.25±1.44 kJ.mol-1, respectively. Therefore, PVC with DEHP has lost more weight than with ATBC or DEI, the activation energy has increased in the case of ATBC with ESO (75.09±2.97 kJ.mol-1) and almost remains stable with ESOME (62.90±3.44 kJ.mol-1) as compared to the plasticizer alone, but in the case of DEI, the activation energy has decreased with the incorporation of ESO (44.97±9.2 kJ.mol-1) or with the incorporation of ESOME (53.79±10.89 kJ.mol-1) in comparison to DEI alone. these energy reductions suggest that when PVC is plasticized with plasticizer system containing ESO and especialy ESOME that capture the HCl that is released by PVC and as such preventing the autocatalytic role of HCl, so the weight loss is reduced, therefore the activation energy has decreased[11]. Plasticizers combination reduces migration as compared to a single plasticizer[28].

and are important parameters for assessing the migration of plasticizers[5]. Figure 3 shows the mass loss by volatility. At elevated temperatures, plasticizers migrate from the PVC resin to the surface where the weight loss by volatility of the plasticizer will negatively affect the mechanical properties of the PVC films[26]. PVC formulations contain ATBC have the highest mass loss values (2.5%, ​​ 4.61% and 8.22%) over the entire test period (24 h, 48 h, and 96 h) at 100 °C. Meanwhile, the weight loss decreases very significantly with the integration of ESO (1.53%, 2.66% and 4.50%) or ESOME (2.07%, 3.65% and 6.19%), conversely, it is noted that the mass losses in the case of DEI alone (0.95%, 1.62% and 2.4%) are comparable to those of DEHP (0.83%, 1.15% and 1.93%), and they decrease with the incorporation of the ESO into the plasticizer system with the DEI at (0.11%, 0.27% and at 0.35%) and with the insertion of ESOME at 0.29%, 0.32% and at 0.32%. This can be explained by the fact that the decrease of the mass loss by volatility in the formulations which contained ESO or ESOME in their plasticizing system, by the reaction of the epoxide function and the HCl produced by the dehydrochlorination of PVC and the higher number of ester bonds which could reduce the volatility[31]. As for the case of the DEI or the ATBC and even the DEHP, herein, there is no reaction between the plasticizer and HCl, assuming that

3.2 Volatility test The volatility of the plasticizers in the polymer is highly dependent on the molecular weight, solubility, compatibility, and chemical structure of the plasticizers[29-31],

Figure 2. Arrhenius activation energy plots for PVC plasticized by ATBC, DEI and their mixtures with ESO and ESOME and DEHP as reference.

Table 2. Loss weight and maximum rate of degradation by isothermal TGA mode at 160 °C, 180 °C and at 200 °C and Arrhenius activation energy. Plasticizer Systems composition (%) DEHP ATBC ATBC/ESO ATBC/ESOME DEI DEI/ESO DEI/ESOME

100 100 50/50 50/50 100 50/50 50/50

160 °C Vmax (%/ Δm/m0 (%) min) 8.70 0.13 21.73 0.43 11.50 0.21 09.96 0.22 4.79 0.10 3.02 0.09 3.87 0 .09

180 °C Vmax (%/ Δm/m0 (%) min) 23.56 0.43 30.31 0.98 17.45 0.56 15.05 0.51 12.90 0.23 06.64 0.13 7.37 0.14

200 °C Vmax(%/ Δm/m0 (%) min) 33.19 0.88 36.05 2.00 20.79 1.22 18.75 0.96 25.18 0.52 12.92 0.26 14.14 0.32

Activation energy kJ.mol-1 81.79±9.63 65.55±0.96 75.09±2.97 62.90±3.44 70.25±1.44 44.97±9.2 53.79±10.89

m: the final mass of the sample; m0: the initial mass of the sample; Δm = (m0-m); V max: the maximum degradation rate in%/min.


Polímeros, 31(3), e2021025, 2021

Assessment of derived sunflower oil as environmentally friendly plasticizers in Poly Vinyl Chloride the mass lost by volatilization at high temperatures contains the volatilized plasticizer and HCl[32,33]. The greatest loss of mass in the case of ATBC can be explained by its boiling point (173 °C) which is much lower than that of the DEI (375 °C) and the DEHP (385 °C). Therefore, the use of ESO and ESOME as secondary plasticizers is recommended to reduce the loss of mass by volatility and maintain the mechanical properties of sheets for longer life time.

3.3 Migration test

Figure 3. Volatility weight loss of mixtures composed by different combinations of plasticizers at 100 °C for 24, 48 and 96 hours.

It is also found that ESOME has a low migration rate as compared to ESO, the increase in the polarity of ESOME

The probability of extraction of the plasticizer increases when a polymer material comes into contact with liquids (solvents)[34]. The resistance to migration of different plasticizer systems in PVC blends was investigated by the migration test using four different solvents as food simulants. Figure 4 shows the mass loss by migration in ethanol, petroleum ether, acetic acid and n-heptane from PVC sheets after 48 and 96 hours where each value was the average of three specimens. It is worth noting that the migration rate of all plasticizer systems in all solvents is lower than that of DEHP. Mixtures of plasticizers have lower migration rates as compared to a single plasticizer such as the cases of DEI and ATBC. Hence, the combination of different plasticizers in PVC plasticization reduces the migration. Indeed, the interactions created between the plasticizers prevent their easy migration from the material to its surrounding environment[35] that is to say that the plasticizer that migrates with a slower speed limits the speed of the plasticizer which diffuses faster. The use of plasticizer blends in PVC samples has been shown to reduce migration in comparison to those containing a single plasticizer[28].

Figure 4. Weight loss by migration in: (a) ethanol; (b) petroleum ether; (c) acetic acid; and (d) n-heptane of the various compositions of the plasticizer systems. Polímeros, 31(3), e2021025, 2021


Bouchoul, B., & Benaniba, M. T. due to the ester functions which make it possible to create interactions with the PVC (interactions between the function C=O and the C-Cl bond)[36] which is responsible for the reduction of the migration rate. In general, the migration of the plasticizers is influenced by the polarity of the plasticizer and the solvent with which the PVC is in contact[37,38], which affects the compatibility of the plasticizer with PVC, that is to say the plasticizer-PVC interaction forces are greater as compared to those of the plasticizer-solvent, then the migration rate is lower, thus, the excellent solvent resistance of plasticizers retain the PVC properties possibly longer life time[19,39], it is suggested that the migration of plasticizers such as ESBO also depends on the manufacturing process of the PVC films[40].

3.4 Plastograph test Figure 5 shows the evolution of the torque as a function of time for the various compositions of the plasticizer system. To be noted is that we find that the curves of formulations containing DEHP, DEI and ATBC have all a similar shape. After the torque goes stable, its value starts to rise again. This is called the onset of degradation[41]. Which is mainly caused by the formation of double bonds, which occurs from the dehydrochlorination of PVC[15], the further increase in the torque curve whose viscosity increases is a result of the crosslinking of chains containing double bonds[42]. All samples degrade after 1000 seconds, but with the incorporation

of ESO or ESOME in the composition of the plasticizer system, the dynamic thermal stability of PVC increases very significantly. Herein, we notice that the degradation time exceeds 2500 seconds. Table 3 presents the fusion time evolution, the minimum torque and the stability time as a function of the composition of the plasticizer mixture. It is observed that the fusion time of the formulations is minimal in the cases of ATBC and DEI alone, but it increases with the integration of ESO or ESOME in the plasticizer mixtures, where it reaches a maximum point at 436 seconds with DEI/ESO, this fusion time increase in the case of ESO is explained due to the large number of carbon in triglyceride which reduces somewhat its compatibility with PVC as reported by Garcia and Marcilla[43]. All formulations have a lower fusion time than DEHP (482 sec.) indicating a better compatibility of PVC with these plasticizer combinations as compared to DEHP. The same thing happened for the minimum torque which is related to the material melt viscosity[44], which increases with the addition ESO or ESOME in the plasticizer systems. It is explained by the high viscosity of these secondary plasticizers as compared to DEI and ATBC. Regarding the dynamic thermal stability plateau in torque/ time curves, where the time at which the torque is constant is considered as a stability time[41], we observe that ESO or ESOME increases the stability time once incorporated in the plasticizer systems. This confirms the results of TGA, where the epoxide function reacts with the HCl released by the PVC dehydrochlorination reaction and limits the autocatalytic role of HCl in PVC degradation.

3.5 Light transmission The transmission of light is a study of the diffusion of electromagnetic radiation through plasticized PVC films in a particular region of the spectrum: in the ultraviolet wavelengths (from 200 to 400 nm), the visible light (from 400 to 750 nm) and the near infrared (from 750 to 1050nm). To do this, we have chosen the transmission almost at half of each range to note the difference between the different formulations, for instance, at 300 nm for UV, 550 nm for the visible range and 900 for IR as presented in Table 4. PVC is optically a very transparent material, but its transparency could be affected by the addition of additives such as thermal stabilizer[45] or lubricant, which explains the lowering of transparency to than 82% for all sheets.

Figure 5. Torque/Time curves for the various compositions of the plasticizer systems (a) in presence of ATBC; (b) in presence of DEI. 6/9

In Figure 6, the light transmission of PVC blends was plotted as a function of the composition of the plasticizer system and as a function of wavelength. Note that the transmission curves look alike for all PVC blends/plasticizers. The results show that transmission in the case of DEI (84.63%, 80.65%, and 11.38%) or ATBC (84.94%, 81.37% and 17.64%) is superior to DEHP (83.92%, 79.62%, and 2.78%) for the infrared, visible and ultraviolet ranges, while the incorporation of ESO or ESOME is observed to decrease the light transmission. This reduction in light transmission is probably due to the type of plasticizer system composition used in the sheet formulation, probably because of the epoxy group, or the length of the secondary plasticizer chain[46]. Polímeros, 31(3), e2021025, 2021

Assessment of derived sunflower oil as environmentally friendly plasticizers in Poly Vinyl Chloride Table 3. Fusion rheological characteristics of PVC formulations plasticized with various compositions of plasticizer systems. Plasticizer systems composition (%) DEHP 100 ATBC 100 ATBC/ESO 50/50 ATBC/ESOME 50/50 DEI 100 DEI/ESO 50/50 DEI/ESOME 50/50

Fusion time (s) 482 354 414 368 372 436 404

Minimum torque (Nm) 6.8 6.6 7.1 7.7 6.3 8.1 8.2

Stability time (s) 1078 738 >2586 >2632 712 >2564 2694

Table 4. Light transmission in PVC sheets. Plasticizer systems composition (%) UV (300nm) DEHP ATBC ATBC/ESO ATBC/ESOME DEI DEI/ESO DEI/ESOME

Transmission (%) Visible (550 nm)

IR (900 nm)

2.78 17.64 4.08 1.08 11.38 3.57 1.28

79.62 81.37 68.96 69.19 80.65 73.91 68.77

100 100 50/50 50/50 100 50/50 50/50

83.92 84.94 79.72 79.74 84.63 82.92 83.28

4. Conclusions

Figure 6. Light transmission variation of different compositions of PVC films (a) in presence of ATBC; (b) in presence of DEI, as a function of wavelength.

All samples show a sharp decrease in transmission at 380 nm but a transmission band of about 280 nm is observed for all formulations. This significant increase in transmission below 300 nm is associated with the C-Cl bond of PVC[47]. Polímeros, 31(3), e2021025, 2021

According to thermogravimetric analysis, formulations which contain DEHP, DEI and ATBC have maximum weight losses, because those are primary plasticizers and do not have a good thermal satability. Therefore, the incorporation of secondary plasticizers either ESO or ESOME into the plasticizer system improves significatively the thermal stability of PVC blends. The decrease in mass loss by volatility in formulations containing ESO or ESOME in their plasticizer system can be explained by the reaction of the epoxide function and the HCl produced by the dehydrochlorination of PVC. The combination of different plasticizers in PVC plasticization reduces migration. Indeed, the interactions created between plasticizers prevent migration. It is also found that the ESOME has a low migration rate compared to the ESO, the increase in the polarity of the ESOME due to the ester functions which makes it possible to create interactions with PVC, in general, the migration of the plasticizers is influenced by the polarity of the plasticizer and the solvent with which the PVC is in contact, which influences the compatibility of the plasticizer with the PVC. The results of the internal mixer show that the thermal stability improves with the increase of the rate of ESO or ESOME. This confirms the results of TGA, where the epoxide function reacts with the HCl released by the PVC dehydrochlorination reaction and limits the autocatalytic role of HCl in PVC degradation. The light transmission test shows that the incorporation of the ESO or the ESOME leads to a difference in transparency, where the transmission of light decreased with the increase of the rate of the latter in the plasticizer system. As purpose of this study was to provide alternatives to phthalates by the use of epoxidized sunflower oil and the methyl ester of epoxidized sunflower oil as secondary plasticizers, mixed with two primary bio sourced plasticizers such as di-esters Isosorbide (DEI) and acetyl tributyl citrate 7/9

Bouchoul, B., & Benaniba, M. T. (ATBC), to plasticize PVC and seek to obtain properties similar or even better than those of di-(2-ethylhexyl) phthalate, then the mixtures of ESO and ESOME with DEI and with ATBC derived from renewable resources, show plasticization efficiency and can be applied as alternative plasticizers for conventional plasticizers based on a fossil source.

5. Acknowledgements The authors would gratefully like to acknowledge the financial support of General Directorate for Scientific Research and Technological Development (DGRSDT), Algeria, for their support in this work.

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ISSN 1678-5169 (Online)


Epoxy hybrid composites reinforced with nanodiamondsilica for abrasive applications Marcos Antônio Barcelos1,2∗ , Mariana Valinhos Barcelos2 , Gabriel Rodrigues de Almeida Neto3 , Antônio Cesar Bozzi4  and Rubén Jesus Sánchez Rodriguez1  Laboratório de Materiais Avançados – LAMAV, Universidade Estadual do Norte Fluminense – UENF, Campos dos Goytacazes, RJ, Brasil 2 Instituto Federal do Espírito Santo – IFES, Vitória, ES, Brasil 3 Programa de Pós-graduação em Engenharia e Ciência dos Materiais, Universidade Federal de São Carlos – UFSCar, São Carlos, SP, Brasil 4 Laboratório de Caracterização de Superfícies e Materiais, Departamento de Engenharia Mecânica, Universidade Federal do Espírito Santo – UFES, Vitória, ES, Brasil 1


Abstract In this work, a ternary composite of epoxy filled with ND and MS was produced for abrasive applications. Surfactants (oleic acid (OA), sodium dodecyl sulfate (SDS) and Triton TX-100 (TX-100)) were used to improve the particle dispersion and, consequently, the composite properties. The elastic modulus increased up to 76% for the sample with 1 wt% ND and 5 wt% ND using TX-100 (1ND5MS-TX100). Regardless of the filler concentration, the particles did not modify the thermal degradation behavior of the epoxy. Thermogravimetric (TGA) and dynamic mechanical (DMA) analyses suggest a strong particle-matrix interface, also evidenced in scanning electron microscope (SEM) micrographs. The composites presented superior tribological performance. 1ND5MS-TX100 presented a wear rate of 2.19 x 10-3 mm3.Nm-1, 61.3% lower than the epoxy. Also, all composites significantly reduced the roughness of the marble, being proportional to the abrasives concentration. Overall, composites with TX-100 presented improved wear behavior. Keywords: epoxy, nanocomposites, nanodiamond, silica, surfactant, thermoset. How to cite: Barcelos, M. A., Barcelos, M. V., Almeida Neto, G. R., Bozzi, A. C., & Sánchez Rodriguez, R. J. (2021). Epoxy hybrid composites reinforced with nanodiamond-silica for abrasive applications. Polímeros: Ciência e Tecnologia, 31(3), e2021026. https://doi.org/10.1590/0104-1428.20210036

1. Introduction The extraction and processing of ornamental stones is an important economic activity in many countries across the globe[1]. Currently, Brazil is the fourth-largest producer of ornamental stones in the world, with year production of 9.5 million tons. According to the Brazilian Association of the Ornamental Stones Industry (ABIROCHAS), only in 2019, Brazil exported 2.1 million tons of ornamental stones, representing a total of nearly US$ 1 billion[2]. Polishing is one of the most critical steps for processing ornamental stones. This step aims to improve some characteristics of the stone, such as color, texture and beauty, in addition to provide unique properties for specific applications. Currently, a combination of resin and diamond, known as resin-based diamond abrasives, is the most common material used in the polishing process of marble and other ornamental stones. Composites filled with particles with high hardness, such as diamond, silicon carbide, alumina, and other ceramic particles have been studied to improve the wear resistance of polishing tools[3]. In addition to the tribological properties, the particles are expected to enhance the stiffness and strength of the matrix[4]. The size of the particle plays an important

Polímeros, 31(3), e2021026, 2021

role in the overall polishing performance. Larger particles tend to be more easily detached from the matrix, leading to premature failure. Another important factor is controlling the distribution and dispersion of the particles, which may impact the material stiffness, impact, and wear resistance[5]. Nanodiamond (ND) and silica (MS) have been reported to improve the wear behavior of composites[6,7]. The addition of only 0.5 wt% of ND enhanced by 95% the wear rate of neat epoxy[6]. Friction coefficients of epoxy resin decreased by a factor of 4 with 7.5 vol% of ND[8]. It was also observed an increase of scratch resistance by the incorporation of 25 vol% of ND[9]. Similarly, MS was reported to improve the frictional coefficient and specific wear rate of epoxy composites[7,10] and to produce wear-resistant coatings[11,12]. The studies reported so far explored the individual addition of ND and MS in an epoxy resin. It has not been studied the combination of both fillers for the preparation of a hybrid composite of epoxy resin matrix. Also, the fillers were added to the epoxy matrix without surface treatment[9,10,12] or with surface chemical modification[6-8,11]. To the best of the authors knowledge, no work has been reported on the addition of



Barcelos, M. A., Barcelos, M. V., Almeida Neto, G. R., Bozzi, A. C., & Sánchez Rodriguez, R. J. different surfactants to improve these particles dispersion in epoxy matrix, and consequently the tribological behavior. In the present study, a ternary hybrid composite of epoxy resin reinforced by ND and MS was produced. The goal was to evaluate the effect of different surfactants and particles concentration on the mechanical, dynamical mechanical, morphological, and tribological properties of the composites. It is expected that the results of this work will contribute to the development of a novel composite material for the final step of polishing of ornamental stones.

2. Materials and Methods 2.1 Materials The epoxy resin used in this study (DER 331 resin) is a transparent liquid of bisphenol-A diglycidyl ether (DGEBA) acquired from Dow Quimica S/A. The epoxy hardening agent was tetraethylene pentaamine (TEPA) obtained from Fluka. The ND (0 – 250 nm) was supplied by Diambra, and the MS (0.5 – 10 µm, with approximately 80% of the particles between 1 – 5 µm) was obtained from Sigma-Aldrich. Acetone was the solvent used in this study, and it was supplied by Neon. The surfactants used were oleic acid (OA, Proquimios), sodium dodecyl sulfate (SDS, Labsynth) and Triton TM X-100 (TX-100, Sigma-Aldrich).

2.2 Preparation of the composites Before the preparation of the composites, the ND and MS particles were dried at 100 °C for 2 h. Then, the epoxy resin was dehydrated in a vacuum oven at 70 mbar and 90 °C to remove humidity and reduce the presence of bubbles. The particle size distributions were obtained in a Sympatec HELOS laser diffraction equipment with 3 bar of pressure, and their x-ray diffractogram was obtained by x-ray diffraction analysis (XRD) in a Bruker D8 Advanced. The ND and MS particles were used to prepare a ternary composite of epoxy / ND / MS with different surfactants. The surfactant concentration was pre-optimized to yield superior dispersion. First, the particles were dispersed with the surfactants (1 g, 0.2 g, and 5 mL for OA, SDS and TX100, respectively) in 20 mL of acetone by bath sonication (SolidSteel, 135 W, frequency of 40 kHz) and mechanical stirring for 20 min. Then, the suspension was added to the epoxy resin, mechanically stirred for 30 min at room temperature, and bath-sonicated under vacuum for 1 h. TEPA was added to yield a DGEBA: TEPA ratio of 1:

0.17 to achieve the stoichiometric ratio of epoxy / equivalent amine, according to the methodology presented in a previous study of our group[13]. It was mixed for at least 5 min, and degassed under vacuum for 30 s. The mixture was poured in a metallic mold and cured in a multistep process: 24 h at room temperature, 4 h at 60 °C, 4 h at 120 °C, and 2 h at 190 °C. The mold dimension was 65 x 12 x 4 mm for flexural and impact testing and 7.94 mm diameter x 67.9 mm height for the samples for wear analyses. A silicon demolding agent was used to aid the removal of the specimens from the mold. The samples will be designated as xNDyMSSurfactant, where x and y are the weight fraction of ND and MS, respectively, and Surfactant is the surfactant used in this formulation when applicable (Table 1).

2.3 Characterization of the composites The thermal stability, the interaction between the fillers and the polymer matrix, and the actual content of the fillers in the samples were evaluated by TGA in a TA Instruments TGA Q5000. It was used a sample size of 10 mg, in a temperature range of 35-950 °C at a heating rate of 10 °C.min-1, under an air flow rate of 25 ml.min-1. The flexural properties of the samples were obtained in an Instron 5582s universal testing machine at 18 °C with a humidity of 50%. The analyses were performed according to ASTM D790, with a specimen size of 65 × 12 × 4 mm, a crosshead speed of 1.4 mm.min-1, and a span length of 50 mm. Four specimens were used for each formulation. The elastic modulus (E), maximum strain (Ɛ), and flexural strength (σm) were calculated and used for the discussion. The impact properties of the composites were examined in a CEAST impact tester. Four specimens (65 x 12 x 4 mm) of each formulation with a v-shaped notch were tested according to ASTM D256, using a 1 J pendulum and impact velocity of 3.46 m/s. The equipment was calibrated to determine the impact energy (J.m-1) of each test. The samples average impact energy was determined. The morphology of the fracture surface of the flexuraltested samples was analyzed in a Zeiss EVO MA 10 SEM. The samples were mounted on stubs and gold-sputtered. Energy dispersive x-ray spectrometry (EDS) was used to map the distribution of C and Si on the samples fracture surface. DMA analysis was performed in a TA Instruments Q800 dynamic mechanical analyzer to obtain the storage modulus (E’), loss modulus (E”), and tanδ (E”/E’) of the samples. The samples were analyzed using a dual cantilever clamp, frequency of 1 Hz, amplitude of 20 µm, a static

Table 1. The designation and the composition of the produced samples. Sample Matrix 1ND3MS 1ND3MS-OA 1ND3MS-SDS 1ND3MS-TX100 1ND5MS 1ND5MS-OA 1ND5MS-SDS 1ND5MS-TX100


Epoxy (wt%) 100 96 96 96 96 94 94 94 94

ND (wt%) 0 1 1 1 1 1 1 1 1

MS (wt%) 0 3 3 3 3 5 5 5 5

Surfactant OA SDS TX100 OA SDS TX100

Polímeros, 31(3), e2021026, 2021

Epoxy hybrid composites reinforced with nanodiamond-silica for abrasive applications force of 0.1 N, a heating rate of 3 °C.min-1 from -90 to 190 °C. The glass transition temperature (Tg) is obtained by the peak of tanδ. The wear behavior of the composites was analyzed in a Phoenix PLINT TE67 pin-on-disk tribometer coupled with the COMPEND 2000 control and data acquisition software (Figure 1). The analysis was performed for 10 min, with a normal load of 70 N, wear path radius of 27.42 mm, a rotational speed of 150 rpm, a tangential speed of 0.42 m.s-1, and a sliding distance of approximately 257 m. Temperature and humidity were controlled by a Minipa MT-240 thermo hygrometer; the temperature ranged from 18.9 °C to 21.4 °C, while the humidity ranged from 35% to 52%. The composite cylinders were 56 mm height × 7.93 mm diameter. The ornamental stone selected for the test was white marble with green veins with a 76 mm diameter and 4 mm height, a central hole of 10 mm diameter, a lateral hole of 5 mm diameter and 4 mm height, which is only present on the side where the stone is fixed to the equipment. Extra care was taken during manipulation and cleaning to assure minimal contamination and, thus, interference in the test results[14]. The marble surface’s mean roughness (Ra) before and after the wear tests was evaluated by laser confocal microscopy in a LEXT Olympus microscope. Roughness provides valuable information regarding the efficiency of the

polishing process. Reducing the roughness is fundamental to minimizing contact deformation, heat generation, electrical current conduction, and other problems[15]. All data obtained in this study were presented as mean ± standard deviation unless stated otherwise. The samples were compared using one-way ANOVA followed by Fisher’s post-hoc test in Minitab 17.3.1 to determine any statistically significant difference for a significance level of 95%.

3. Results and Discussions 3.1 Thermogravimetric analysis (TGA) The TGA curves provide the weight variation of the sample with the temperature (Figure 2). The fillers, ND and MS, presented high thermal stability. MS did not undergo any weight loss event at up to 800 °C, while ND presented a single weight-loss event, starting at 545.6 °C, related to the oxidation process[16]. Regarding the polymer composite, there is not a notable difference in the thermal behavior between the samples with different surfactants. However, comparing to the matrix, the addition of the fillers slightly increased the thermal stability of the matrix. This may result from the interaction between the polymer chains and the fillers, requiring more energy to overcome it[17,18].

Figure 1. (a) Phoenix PLINT TE67 apparatus; (b) the contact between the composite pin and the disk.

Figure 2. TGA curves of the particles, the epoxy matrix, and the produced composites. Polímeros, 31(3), e2021026, 2021


Barcelos, M. A., Barcelos, M. V., Almeida Neto, G. R., Bozzi, A. C., & Sánchez Rodriguez, R. J. The weight fraction of MS present in the samples can be estimated by the residual weight at 800 °C, because at this temperature the matrix is fully decomposed, presenting a negligible residual weight, and the ND is fully oxidized. The mean residual weight of the composites was 2.83 and 4.51%, agreeing with the weight fraction of MS added to each sample; the small difference is likely caused by the particle loss during the preparation process.

3.2 Mechanical characterization The stress-strain curves obtained by flexural testing at room temperature are presented in Figure 3. The curve selected to represent the sample behavior was the one that presented the flexural properties closer to the mean values. Overall, the introduction of fillers reduced the flexural strength and maximum strain of the matrix; this reduction was proportional to the filler concentration. Depending on the polymer-filler interface properties and the dispersion degree of the fillers, high filler loading may reduce the macromolecular mobility, thus hindering deformation[19]. The elastic modulus, on the other hand, presented a statistically significant increase (p < 0.05) by the introduction of the fillers regardless of the composition (Figure 4a). The matrix presented an elastic modulus of 1.54 GPa, which increased to up to 2.71 GPa for 1ND5MS-TX100, representing a 76% increase. This percentage increase is higher than that observed for the

Figure 3. Stress × strain curves obtained by flexural testing for the samples produced. 4/10

incorporation of up to 40% of micrometric diamond particles in the same type of epoxy matrix (DGEBA-TEPA) [20] . The mildest improvement of elastic modulus was for 1ND5MS-SDS, but it still increased by 22%. The most significant improvement was found for 5 wt% fraction of MS, also suggesting a direct correlation with the filler concentration[21]. Diamond and silica are known for their high elastic modulus, but the overall performance of the composite does not depend exclusively on the filler elastic modulus; it also requires a strong interfacial interaction between the components. TX-100 was the most effective surfactant on increasing the elastic modulus. The other surfactants yielded similar results. The impact resistance of an epoxy-based composite may reduce with a poor dispersion of the filler – acting as a stress concentrator – and with the presence of defects produced during the matrix curing step[22]. The mean energy absorbed by the samples during the fracture induced by the impact test is presented in Figure 4b. Apart from 1ND5MS-SDS, there was no significant difference in the energy absorbed by the samples compared to the matrix (p < 0.05). The highest increase in the mean value was observed for 1ND5MS-OA (11.4%). Even though the fillers did not improve the impact resistance of the epoxy matrix, they did not decrease it either. Therefore, it can be concluded that for high strain rates analysis, the effect of stress concentration was mild, and the ND and MS interfacial bonds were strong enough to maintain the energy absorption capacity of the system.

Figure 4. (a) Elastic modulus obtained in the flexural testing and (b) impact energy for the studied samples. N.S. stands for no surfactant. *p < 0.05 compared to the matrix, **p < 0.05 compared to the N.S. system with the same concentration of particles, ***p < 0.05 compared to the system with the same surfactant but a different concentration of particles (One-way ANOVA, followed by Fisher’s post-hoc test). Polímeros, 31(3), e2021026, 2021

Epoxy hybrid composites reinforced with nanodiamond-silica for abrasive applications 3.3 Morphological characterization The morphological aspects of the fracture surface of the samples after flexural testing were contrasting (Figure 5). Pure epoxy presented a smooth surface, with river patterns, a typical characteristic of brittle fracture[23,24]. On the other hand, the composites presented a rougher surface (Figure 5b), possibly due to crack deflection by the anchored particles[22,25]. The fillers adhered to the surface after the fracture, possibly due to the strong bonding between the polymer matrix and the particles. From the micrographs, the quality of the processing conditions of the specimens was confirmed, as no voids from bubbles trapped into their interior were observed. EDS element mapping is a useful tool to evaluate the particle dispersion qualitatively. The samples’ fracture surface was mapped for identifying carbon and silicon (Figure 6). The region that contains silicon is represented in pink color, while green regions stand for the presence of carbon. Silicon, which is only present in the MS, is evenly distributed across the fracture surface of both samples (Figures 6c and 6d). The pink regions are more frequent and brighter for 1ND5MS, which is exclusively attributed to the higher loading of MS particles in this formulation. The dispersion of ND particles could not be evaluated by this analysis, because its main element (carbon) is the same as the epoxy matrix. Despite the particles having nanostructured grains, confirmed by XRD analysis, they formed micrometric agglomerates (Figure 7). Nanodiamonds produced by the detonation route have a surface rich in oxygenated groups. The secondary forces between these groups, combined with the high surface

area per volume of the nanoparticles, are the driving force for the agglomeration. We have observed in another study the tendency of ND agglomeration[26]. The individual particles were analyzed by laser diffraction of dry dispersion before the addition to the matrix. A mean agglomerate size (x50,3) of 1.8 µm and 1.5 µm was obtained for ND and MS, respectively. Agglomerates as large as 24.9 µm and 17.9 µm were found on the fracture surface of 1ND3MS and 1ND5MS. The use of surfactants effectively reduced the agglomerate size. OA was the least effective surfactant, but still significantly hindered the formation of agglomerates, reaching 4.3 µm and 9.1 µm, for 1ND3MS-OA and 1ND5MS-OA, respectively. The agglomeration of the particles decreased even more with SDS and TX-100. The effect of these surfactants was similar for higher loadings: 2.1 µm and 2.3 µm for 1ND5MS-SDS and 1ND5MS-TX100, respectively. For lower loadings, TX100 was slightly more effective, producing agglomerates of up to 2.9 µm, as opposed to 3.1 µm of SDS.

3.4 Dynamical mechanical characterization The DMA curves presented in Figure 8 show the influence of the fillers on the viscoelastic properties of epoxy resin. The summary of the properties obtained is presented in Table 2. The damping behavior obtained by this analysis is sensitive to the material’s interface; therefore, to elucidate the contribution of each filler interface, samples containing only ND and surfactants were also tested[27] (Figure 8a). The values of the storage modulus of the samples were similar. All samples displayed the same behavior with increasing temperature. Initially, the storage modulus smoothly decreased, and then, it dropped two orders of magnitude. The sharp reduction is associated with the onset of collaborative motion of the chain segments, which can facilitate mechanical failure. The addition of ND shifted Tg to higher temperature, increasing from 127.2 °C to 137.5 °C. The surfactants were effective in increasing it even more, reaching 145.5 °C for OA and SDS. Usually, such behavior is a consequence of the restricted motion of the polymer chains because of the interaction with the filler. Therefore, the improvement after the addition of surfactants might result from the improved interfacial interaction between the matrix and the particles and the enhanced particle dispersion[27,28]. Kubát et al.[27] proposed a parameter A (Equation 1), which estimates the contribution of the interface to the overall damping of the composite; in other words, it can estimate the interaction strength between the filler and the matrix[28]. If A equals zero, there is no dissipation of energy on the interface and good adhesion. Higher values of A suggest a larger contribution of damping for the interface. The A values calculated for the system with 1 wt% ND (Table 2) show the positive effect of the surfactants on the interaction between the filler and the matrix. The composites with surfactant presented significantly smaller A than that of the composite without surfactant.

A = Figure 5. SEM micrographs of the fracture surface of (a) neat epoxy resin (b) and 1ND5MS. Polímeros, 31(3), e2021026, 2021

1 tanδ c − 1 (1) 1 − v f tanδ m

tanδ= c

(1 − Bv f ) tanδ m (2) 5/10

Barcelos, M. A., Barcelos, M. V., Almeida Neto, G. R., Bozzi, A. C., & Sánchez Rodriguez, R. J.

Figure 6. EDS element mapping of C (green) and Si (pink) on the fracture surface of 1ND3MS (a and c) 1ND5MS (b and d). Table 2. Summary of the samples properties obtained by DMA. Sample Matrix 1ND 1ND-AO 1ND-SDS 1ND-TX100 1ND3MS 1ND3MS-AO 1ND3MS-SDS 1ND3MS-TX100 1ND5MS 1ND5MS-AO 1ND5MS-SDS 1ND5MS-TX100

Tg (°C) 127.2 137.5 145.5 145.5 138.1 122.2 130.1 138.6 116.6 147.3 137.5 145.3 112.2

Tanδ at Tg 0.6685 0.6784 0.4957 0.4935 0.5235 0.5609 0.5635 0.6063 0.5409 0.5106 0.5566 0.4906 0.5114

where v is the volume fraction, and the subscripts f, c and m refer to filler, composite and matrix, respectively. Ziegel and Romanov[29] proposed a volume term B to account for the formation of immobilized interphase because of strong adhesion between particle and matrix (Equation 2). Again, the composites treated with surfactants presented higher B than the neat composite (Table 2). This suggests the formation of a thicker interphase layer[28] for the composites with surfactant. Regarding the hybrid composites, overall, Tg increased from a 3 wt% to 5 wt% MS concentration. This may be related to the restriction of the macromolecular motion of the epoxy matrix by the incorporation of the particles. 6/10

A 0.02 -0.26 -0.26 -0.21 -

B -4.00 69.86 70.75 58.62 -

For the 1ND3MS samples, a clear response regarding the type of surfactants used was not observed. However, for the 1ND5MS composition, the Tg of all samples prepared with surfactants were smaller than that of 1ND5MS sample. Amongst the surfactants, TX-100 presented the largest decrease in the Tg. This effect can be related to a different distribution of free volume between the systems. The addition of fillers with a high ratio of surface area per volume creates a large volume of polymer-filler interface[30]. The same particles with improved dispersion degree have larger interfacial free volume. This additional free volume facilitates the chain mobility in this region[31]. Thus, a possible explanation for the reduction of Tg for the samples using TX-100 is the reduction of the agglomerate Polímeros, 31(3), e2021026, 2021

Epoxy hybrid composites reinforced with nanodiamond-silica for abrasive applications

Figure 7. SEM micrographs showing agglomerates on the fracture surface of (a) 1ND5MS; (b) 1ND5MS-OA; (c) 1ND5MS-SDS; (d) 1ND5MS-TX100.

Figure 8. Evolution of storage modulus and tan delta with the temperature for all studied samples. Polímeros, 31(3), e2021026, 2021


Barcelos, M. A., Barcelos, M. V., Almeida Neto, G. R., Bozzi, A. C., & Sánchez Rodriguez, R. J.

Figure 9. Influence of the surfactant used and abrasive particles concentration on (a) the wear rate (k) and (b) the reduction of the mean roughness (Ra) of the marble after polishing with the samples. N.S. stands for no surfactant; (c) Photo of the polished marble after the wear test, and 3D image of the transition between the (d – A) unpolished and (d – B) polished region of the marble. In (a), the data is presented as mean (standard deviation). *p < 0.05 compared to the matrix (one-way ANOVA, followed by Fisher’s post-hoc test).

size, which was previously discussed in the morphological characterization, and reflected in the sample’s flexural elastic modulus and tribological properties (to be discussed). However, a possible plasticizing effect of the surfactants cannot be discarded[30]. Kubát et al.[27] and Ziegel and Romanov[29] parameters could not be used to assess the interface of the hybrid composites, as they were designed for single fillers, and it cannot account for the individual contributions and interactions of each filler. However, there was no significant difference between the height of tan delta peaks of the composites, suggesting that there was no effect of the surfactants between these formulations.

3.5 Tribological analysis The influence of the hard and stiff ND and MS particles on the wear behavior of the epoxy matrix was evaluated in terms of the wear rate (k) and the average surface roughness (Ra) (Figure 9). The k values were considerably reduced for the composites (Figure 9a) (p < 0.05). 1ND5MSTX100 presented a k of 2.19 x 10-3 mm3.Nm-1, 61.3% lower than that of the neat epoxy matrix. Ayatollahi et al. [6] observed a reduction of 95% of k for 0.5 wt% of aminofunctionalized ND with a diameter of 9-15 nm, indicating the important role of well-dispersed nanoparticles. Similarly, Zhang et al.[7] reported a greater reduction of k for surfacefunctionalized nanosilica. Oppositely, in another study[11], 5.9 wt% of nanosilica increased k of epoxy resin, the authors suggest the role of nanoparticle agglomeration on the wear behavior. A strong interface between the matrix and particles might be the reason for the improved wear resistance in our study[8]. Overall, k of the composites was similar, and no difference was observed among the surfactants used. 8/10

The reduction of the marble roughness after polishing with the samples did not follow the same trend as that of k (Figure 9b). The difference of Ra was higher for the samples with a surfactant, being the least effective to the most effective: OA, SDS, and TX-100. Also, it was observed that the Ra increased with the concentration of MS for the system with and without surfactant, showing proportionality to the abrasive particle concentration. 1ND5MS-TX100 presented the highest Ra, 42.4%, nearly 105% higher than that of 1ND5MS. The wear track was seen by visual observation (Figure 9c). Laser confocal microscopy (Figure 9d) confirms the efficiency of the composites produced herein for the polishing of ornamental stones, showing a higher difference of the height of peaks and valleys in the unpolished region. The wear behavior of the composites results from the physical properties of the fillers, such as hardness, which has a primary role in the polishing procedure. The improved elastic modulus obtained in this study is also beneficial for the desired application[7,10,12].

4. Conclusions Hybrid composites of epoxy filled with nanodiamond and silica were successfully produced. The surfactants effectively reduced the tight agglomerates formed by the nanoparticles; however, micrometric clusters were still observed. The impact energy of the composites was not smaller than that of the neat matrix, which is expected when there are agglomerates that may act as a stress concentrators. The synergetic combination of the particles improved the flexural elastic modulus and the wear behavior of the composites. The sample 1ND5MS-TX100 was the one that presented the best results. In addition to presenting the smallest agglomerates, the elastic modulus increased by 76%, Polímeros, 31(3), e2021026, 2021

Epoxy hybrid composites reinforced with nanodiamond-silica for abrasive applications the mean impact energy raised 11.4%; it also displayed the lowest wear rate and the most significant reduction in the roughness of the polished marble. Therefore, this epoxybased composited containing nanodiamond and silica with the addition of TX-100 is promising for abrasive applications, such as the polishing of marble stones.

5. Acknowledgements This study was financed in part by the Coordenação de Aperfeiçoamento de Pessoal de Nível Superior - Brasil (CAPES) - Finance Code 001. The authors would like to thank to Conselho Nacional de Desenvolvimento Científico e Tecnológico (CNPq) (process 310108/2017-9), FAPERJ (process E-26/203.016/2018) and Ifes for the financial support.

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Barcelos, M. A., Barcelos, M. V., Almeida Neto, G. R., Bozzi, A. C., & Sánchez Rodriguez, R. J. Materials Science and Engineering A, 528(24), 7295-7300. http://dx.doi.org/10.1016/j.msea.2011.06.053. 26. Almeida, G. R., No., Barcelos, M. V., Rodríguez, R. J. S., & Gomez, J. G. C. (2017). Influence of encapsulated nanodiamond dispersion on P(3HB) biocomposites properties. Materials Research, 20(3), 768-774. http://dx.doi.org/10.1590/19805373-mr-2016-0715. 27. Kubát, J., Rigdahl, M., & Welander, M. (1990). Characterization of interfacial interactions in high density polyethylene filled with glass spheres using dynamic-mechanical analysis. Journal of Applied Polymer Science, 39(7), 1527-1539. http://dx.doi. org/10.1002/app.1990.070390711. 28. Dong, S., & Gauvin, R. (1993). Application of dynamic mechanical analysis for the study of the interfacial region in carbon fiber/epoxy composite materials. Polymer Composites, 14(5), 414-420. http://dx.doi.org/10.1002/pc.750140508.


29. Ziegel, K. D., & Romanov, A. (1973). Modulus reinforcement in elastomer composites. I. Inorganic fillers. Journal of Applied Polymer Science, 17(4), 1119-1131. http://dx.doi.org/10.1002/ app.1973.070170410. 30. Biswal, M., Mohanty, S., Nayak, S. K., & Kumar, P. S. (2013). Effect of functionalized nanosilica on the mechanical, dynamicmechanical, and morphological performance of polycarbonate/ nanosilica nanocomposites. Polymer Engineering and Science, 53(6), 1287-1296. http://dx.doi.org/10.1002/pen.23388. 31. Rakha, S. A., Raza, R., & Munir, A. (2013). Reinforcement effect of nanodiamond on properties of epoxy matrix. Polymer Composites, 34(6), 811-818. http://dx.doi.org/10.1002/pc.22480. Received: Apr. 21, 2021 Revised: Oct. 12, 2021 Accepted: Oct. 19, 2021

Polímeros, 31(3), e2021026, 2021

ISSN 1678-5169 (Online)


Esterification of oleic acid employing sulfonated polystyrene and polysulfone membranes as catalysts Ana Paula de Lima1 , Andressa Tirone Vieira2 , Bárbara Nascimento Aud1,3 , Antonio Carlos Ferreira Batista2 , Luís Carlos de Morais4 , Anízio Márcio de Faria2 , Rosana Maria Nascimento de Assunção2  and Daniel Pasquini1*  Instituto de Química, Universidade Federal de Uberlândia – UFU, Uberlândia, MG, Brasil Instituto de Ciências Exatas e Naturais do Pontal, Universidade Federal de Uberlândia – UFU, Ituiutaba, MG, Brasil 3 Departamento de Áreas Acadêmicas, Instituto Federal de Goiás, Itumbiara, GO, Brasil 4 Instituto de Ciências Exatas, Naturais e Educação, Universidade Federal do Triângulo Mineiro – UFTM, Uberaba, MG, Brasil 1



Abstract In the present study, catalytic activity of dense, porous, electrospun membranes of polysulfone (PSF) and polysulfone with sulfonated polystyrene (PSF_PSS) have been evaluated in reactions of esterification of oleic acid with methanol, in times that varied from 10 to 480 minutes. Conversion to biodiesel has been confirmed by FTIR and quantified through gas chromatography. The results showed the catalysts used were effective in the esterification reaction studied and the PSF_PSS electrospun membrane has presented the best conversion to methyl oleate, reaching 70.5% in a 10-minute reaction and 95.8% in a 240-minute reaction, when methanol:oleic acid molar ratio of 10:1, 5% of catalyst and temperature of 100 °C were used. Considering the performance of solid catalysts described in literature, mainly related to reaction times and conversion of the process, this study reveals a promising feasibility of using electrospun membranes of PSF_PSS for developing a heterogeneous acid catalyst aimed to biodiesel synthesis. Keywords: biodiesel, esterification, membranes, polysulfone, sulfonated polystyrene. How to cite: Lima, A. P., Vieira, A. T., Aud, B. N., Batista, A. C. F., Morais, L. C., Faria, A. M., Assunção, R. M. N., & Pasquini, D. (2021). Esterification of oleic acid employing sulfonated polystyrene and polysulfone membranes as catalysts. Polímeros: Ciência e Tecnologia, 31(3), e2021027. https://doi.org/10.1590/0104-1428.20210067

1. Introduction Increasing population, urbanization and industrialization induce to a drastic need of energy. It is estimated that the global demand for energy will increase from almost 286 million barrels of oil equivalent per day (mboe/d) in 2018 to more than 357 mboe/d in 2040, with an average increase of around 1% per year[1]. The main energy source to the whole world is fossil biofuels, which includes all conventional sources of energy, such as petroleum-based, methane and coal. Due to the non-renewable nature of these resources, systematic rise in energy import prices and environmental factors, the interest in alternative sources of energy is increasingly growing[2]. In this context, biodiesel emerges as a biofuel that has similar properties to diesel fuel and has several benefits, such as being renewable, biodegradable, non-flammable, non-toxic, non-explosive and with low level of sulphur and pollutants emissions[3,4]. Alkyl fatty acid esters, or biodiesel, may be produced by transesterification of triglycerides or by esterification of free fatty acid with short chain alcohols in the presence of a catalyst[5]. The transesterification reaction performed with homogeneous alkaline catalysts is the technology

Polímeros, 31(3), e2021027, 2021

commonly applied in biodiesel industry, presenting high yields, moderate operation conditions and fast reaction rates. Nevertheless, when oils with high acidity and moisture are applied with basic catalysts, production effectiveness decreases due to saponification reactions[2,6,7]. Although homogeneous acid catalyst is insensitive to free fatty acid content and has the potential to simultaneously perform esterification and transesterification reactions, the process has some disadvantages, such as equipment corrosion, slow reaction rates and great amount of effluents to be treated, increasing environmental pollution and production costs[6,8]. Thus, solid acid catalysts are a promising replacement to the process, since they are easily recovered and reused in the reaction, avoid reactor corrosion problems and reduce the stages of product purification[3,4]. Recently, fatty acid esterification from long hydrocarbon chain fatty acids in the presence of acid catalysts has drawn great interest, since the esters produced can also be applied as biofuels[9]. Comprehensively considering, heterogeneous acid catalysts as sulfated zirconia[10], ion exchange resin[11], materials based on carbon[7], zeolite[12] and heteropoly acids[13] are good



Lima A. P., Vieira A. T., Aud B. N., Batista A. C. F., Morais L. C., Faria A. M., Assunção R. M. N. & Pasquini D. examples of solid catalysts to this reaction, once they display relatively high catalytic activity in moderate conditions[14]. When it comes to esterification reaction mechanisms, the process of biodiesel acquisition occurs with the formation of an oxonium ion through protonation of the fatty acid, thereby increasing the electrophilic character of carbonyl group and facilitating the nucleophilic attack from the alcohol in order to produce a tetrahedral intermediate. Subsequently, this intermediate goes through rearrangement, a water molecule is withdrawn, the ester is formed and the catalyst is recovered[2], as shown in Figure 1. Functional polymer materials, such as microspheres and membranes, chemically modified with acid groups are also viable options to overcome deficiencies and replace liquid acids in several organic reactions (transesterification and esterification, among them). The most used in heterogeneous acid catalysis are sulfonic resins (cation exchangers), Nafion non-porous resins and Amberlyst macroporous resins being

the most common[15]. In literature, there are many papers which report the use of acid polymer materials as solid catalysts in reactions aiming biodiesel production[16-21]. Polystyrene (PS), a thermoplastic polymer with good mechanic and insulating properties widely used in beaker production, transparent packing for foods, audio/video packaging, and light bulbs caps, among others[22], may be easily modified through sulfonation reactions to be applied as solid catalyst[18]. Meanwhile, a careful control of its degree of sulfonation (DS) is needed, considering some materials with high DS - for being soluble in water - would have little application in heterogeneous catalysis[23]. Also belonging to the thermoplastic class, there is polysulfone (PSF), a polymer formed by two monomers, bisphenol-A and diphenyl sulfone, which has great mechanic resistance, good dimensional and thermic stability, adequate flexibility, and high chemical and hydrolytic stability[24,25]. This polymer has been broadly used in resin manufacturing are allocated to the construction of

Figure 1. Esterification mechanism of fatty acids catalyzed by acid. 2/8

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Esterification of oleic acid employing sulfonated polystyrene and polysulfone membranes as catalysts parts and equipments in electrical and automotive engineering, household and medical equipments and also as membrane materials for being highly thermostable[26]. Considering this context, the preparation and characterization of dense (D), porous (P) and electrospun (E) membranes of polysulfone and polysulfone with sulfonated polystyrene (PSF_PSS) had been reported in a previous paper generated in our laboratory, aiming a preliminary evaluation of their catalytic properties in esterification reactions from oleic acid with methanol[27]. Due to the obtained results that illustrated the good performance of PSF_PSS membranes as heterogeneous acid catalysts in the reactional time of two hours, esterification experiments have been expanded in the work reported in the present paper, ranging reaction times, with the aim of improving the maximum obtainable conversion, linking the catalytic activity with physical and chemical properties from different membranes.

SEM micrographs in Figure 2. Potassium bromide (PA, Vetec) for infrared analysis and methyl oleate (Sigma Aldrich) for gas chromatography with flame ionization detector have also been used in the development of this paper. The sulfonation degree (SD) of PSF_PSS membranes was determined by titration with a standardized NaOH solution 0.01 mol L-1. The SD was calculated using Equation 1.

SD =

(104*M NaOH *VNaOH )

m − ( 81*M NaOH *VNaOH )


Where: 104 = Molar mass of the monomeric unit (g mol-1). MNaOH = Concentration of the standard NaOH solution (mol L-1). VNaOH= Volume of NaOH solution spent in the titration (L). m = Sample mass (g).

2. Materials and Methods

81= Molar mass of the SO3H group (g mol-1).

2.1. Materials

2.2. Esterification reaction

For the esterification reaction, oleic acid (Fluka Analytical) and methanol (PA, Sigma Aldrich) have been used. Dense, porous and electrospun membranes of polysulfone, with maximum degradation temperatures (Tmáx) of 531, 532 and 531 °C, respectively, and dense, porous and electrospun membranes of polysulfone with sulfonated polystyrene, with Tmáx of 526, 527 and 527 °C and superficial area of ~0, 5.4 and 184.4 m2 g-1, respectively, have been used as catalysts in this reaction. These membranes have been produced and characterized according to the previously mentioned paper[27] using the Carl Zeiss EVO MA10 and Tescam VEGA 3 LMU microscopes, which evidenced the homogeneous structures, materials with pores and materials formed by nanofiber networks. Dense, porous and electrospun membranes, respectively, have been produced as shown in

The catalytic performance of PSF and PSF_PSS membranes has been evaluated in the previously mentioned paper[27] in esterification reactions from oleic acid with methanol in a two-hour time. In this present paper, catalytic activity from membranes has been evaluated in different reactional times, ranging from 10 to 480 minutes. For this reaction, 5% of catalyst was added to methanol and the compound was kept resting for 24 hours. Afterwards, this solution was put into a Parr reactor, model 4848, together with oleic acid. The alcohol:oleic acid molar ratio was 10:1 and the reaction occurred at 100 °C. At the end of the process, the catalyst was removed from the reactional mean by simple filtration, and the product was submitted to evaporation under a reduced pressure process in order to withdraw excessive water and methanol.

Figure 2. SEM of membranes: (a) PSF Dense (magnification 2,000x); (b) PSF Porous (magnification 2,000x); (c) PSF Electrospun (magnification 250x); (d) PSF_PSS Dense (magnification 2,000x); (e) PSF_PSS Porous (magnification 2,000x); (f) PSF_PSS Electrospun (magnification 250x)[27]. The inset shows SEM images of the fractured surface on their respective membranes (magnification 5,000x). Polímeros, 31(3), e2021027, 2021


Lima A. P., Vieira A. T., Aud B. N., Batista A. C. F., Morais L. C., Faria A. M., Assunção R. M. N. & Pasquini D. The catalyst reuse test was performed for the PSF_PSS-E 10 min membrane. The catalysts were separated by simple filtration, washed with 10 mL anhydrous ethanol and 10 mL water, followed by a drying at 70 ºC for 3 h.

2.3. Characterization techniques Analyses of characterization performed to the obtained methyl esters included Fourier transform infrared (FTIR) and gas chromatography with flame ionization detector (GC-FID). FTIR spectra, using a Shimadzu IRPrestige-21 equipment, has been used to identify functional groups in the samples by KBr method and registered in the range of 4000 – 400 cm–1 [28]. For quantifying the methyl ester formation, an analytical curve from methyl oleate standard has been elaborated, being the analyses from the standard and the conversion products performed in a Thermo® gas chromatograph, model Focus GC. The temperature of the column oven was kept in isotherm to 190 °C. Other technique elements included: split injection, with flux division 1:10 and temperature of 250 °C; flame ionization detection, with temperature detection of 250 °C; injection of 1 μL sample; total analysis time for each sample of 5 minutes. Applied column: capillary column of stationary poly (ethylene glycol), Carbowax brand, 30 m length, 0.32 mm intern diameter and 0.25 μm film thickness.

3. Results and Discussions The polymeric catalyst activity is influenced by the incorporated sulfonic groups providing acidic sites for the esterification. Thus, the PSF_PSS membranes were evaluated for the sulfonation degree and obtained a SD value equal to 24%.

Figures 3 and 4 show spectra from products obtained by esterification using PSF and PSF_PSS membranes, respectively, in all reaction times studied. As the main characteristic changes of oleic acid conversion in biodiesel are focused in the 1800 to 1100 cm-1 region, only this will be presented to make the discussion easier. Analyzing the spectra, it is observed that the band in 1710 cm-1 corresponding to the stretching from C=O bond of acids is more pronounced in products that used PSF membranes, proving that there was not oleic acid full consumption in these esterification reactions. Band in 1742 cm-1 assigned to C=O bond from esters is more pronounced in products obtained when PSS_PSF membranes have been applied as catalysts, pointing an increase in biodiesel formation. Furthermore, stretching from O-CH3 bond located in 1196 cm-1, typical of methylic biodiesel, is more evident in product spectra obtained when PSF_PSS membranes catalyzed the reaction, showing greater oleic acid conversion to methyl oleate in these reactional systems[8,29,30]. Biodiesel quantification produced by the esterification reactions has been performed by GC-FID and an analytical curve has been built to determine the concentration of methyl oleate formed. Linear regression proved to fit well to data, with linear regression coefficient of 0.998, by the internal standard method. Biodiesel conversion values obtained in catalyzed reactions by dense, porous and electrospun membranes of PSF and PSF_PSS, after the reaction times of 10, 30, 60, 90, 120, 240, 360 and 480 minutes are listed in Table 1. Analyzing Table 1, it is observed that, in general, PSF membranes have presented low oleic acid conversion into methyl oleate, showing average results only in long reactional times, while PSF_PSS membranes have presented better results in all periods of time studied, reaching even satisfying

Figure 3. FTIR of membranes of PSF at all reaction times: (a) dense, (b) porous and (c) electrospun. 4/8

Polímeros, 31(3), e2021027, 2021

Esterification of oleic acid employing sulfonated polystyrene and polysulfone membranes as catalysts

Figure 4. FTIR of membranes of PSF_PSS at all reaction times: (a) dense, (b) porous and (c) electrospun.

conversions in reaction times relatively low. PSF electrospun membrane provided a conversion of 7.5% in 10 minutes, whereas PSF_PSS electrospun membrane converted 70.5% of fatty acid to methylic ester in the same time range. Comparing the performance of porous membranes in 60 minutes of reaction, it is observed that the PSF membrane resulted in a conversion of 18.0%, whereas PSF_PSS membrane led to a conversion of 65.5%. As an example of dense membranes, a conversion of 15.7% was obtained for the PSF membrane, whereas the catalyzed reaction with PSF_PSS converted 63.9% in 90 minutes of reaction. These results showed that the presence of sulfonated polystyrene in membranes has increased the conversion efficiency, as more active acid sites were available to catalyze the reaction. The best conversion rates found for a PSS_PSF electrospun membrane, even in short reaction times, may be attributed to its large superficial area, which allows reagents (methanol and oleic acid) to effectively spread in the catalytically active sites, increasing interaction due to better accessibility[5,11,20]. The superficial area of this membrane, for example, is approximately 34 times larger than the one from PSF_PSS porous membrane, whereas dense membrane has a superficial area near zero. In some cases, reaction conversion rates decreased while reaction time increased. For instance, the catalyzed esterification with PSF_PSS electrospun membrane, has converted 81.4% in a 360-minute reaction, while in 480 minutes this value has dropped to 68.1%. This may be attributed to the reversibility of esterification reactions, indicating that prolonged times may cause hydrolysis of the formed ester[7,8]. A pattern of similar results had been reported in the authors’ previously mentioned paper, which has evaluated the catalytic performance of these membranes in two reaction Polímeros, 31(3), e2021027, 2021

times as a function of oleic acid consumption. The values of fatty acid concentration found have shown that PSF membranes were not effective to catalyze the reaction, and better results were obtained when PSS_PSF porous and electrospun membranes were used as catalysts[27]. The use of polymer acids catalysts in esterification reactions is already the subject of study in literature. In the paper of Grossi et al.[18], oleic acid esterification has been performed using sulfonated polystyrene and a conversion of approximately 90% was found in 8 hours of reaction. Gomes et al.[20] have studied esterification of several fatty acids catalyzed by porous co-polymer polydivinylbenzene-co-triallylamine and it was found for the oleic acid methanolysis a yield of 92% after 10 hours of reaction. Dechakhumwat et al.[5] studied the catalytic activity from a sulfonated residue derived from corn on the cob in oleic acid esterification with methanol, and they obtained a methyl oleate yield next to 80% after 8 hours of reaction. Pan et al.[21] have carried out esterification reactions using an acid catalyst obtained by sulfonation of ethylenediamine in polydivinylbenzene, and after a 4-hour reaction at 100 °C they obtained 85% conversion. Considering the conversion results obtained in the present study, one can affirm that though heterogeneous catalysis usually requires more severe reactional conditions and longer periods of time to obtain a good performance, 10 minutes of reaction were enough to obtain a significant conversion when PSF_PSS electrospun membrane was applied as reaction catalyst. The catalyst reuse test was performed for the reaction condition: molar ratio oleic acid:methanol 1:10, temperature of 100 °C, 5% catalyst and 10 minutes of reaction for the electrospun membrane (PSF_PSS-E 10 min). Figure 5 show FTIR from products obtained by esterification in three reuse cycles. 5/8

Lima A. P., Vieira A. T., Aud B. N., Batista A. C. F., Morais L. C., Faria A. M., Assunção R. M. N. & Pasquini D. Table 1. Percentage conversion of oleic acid to methyl oleate, measured by GC-FID. Sample PSF-D 10 min PSF-D 30 min PSF-D 60 min PSF-D 90 min PSF-D 120 min PSF-D 240 min PSF-D 360 min PSF-D 480 min PSF-P 10 min PSF-P 30 min PSF-P 60 min PSF-P 90 min PSF-P 120 min PSF-P 240 min PSF-P 360 min PSF-P 480 min PSF-E 10 min PSF-E 30 min PSF-E 60 min PSF-E 90 min PSF-E 120 min PSF-E 240 min PSF-E 360 min PSF-E 480 min PSF_PSS-D 10 min PSF_PSS-D 30 min PSF_PSS-D 60 min PSF_PSS-D 90 min PSF_PSS-D 120 min PSF_PSS-D 240 min PSF_PSS-D 360 min PSF_PSS-D 480 min PSF_PSS-P 10 min PSF_PSS-P 30 min PSF_PSS-P 60 min PSF_PSS-P 90 min PSF_PSS-P 120 min PSF_PSS-P 240 min PSF_PSS-P 360 min PSF_PSS-P 480 min PSF_PSS-E 10 min PSF_PSS-E 30 min PSF_PSS-E 60 min PSF_PSS-E 90 min PSF_PSS-120 min PSF_PSS-E 240 min PSF_PSS-E 360 min PSF_PSS-E 480 min

Conversion (%) 10.4 7.8 8.3 15.7 12.2 21.2 54.7 42.4 6.6 12.7 18.0 22.3 19.3 31.3 52.1 59.6 7.5 11.2 11.3 14.4 16.6 26.4 51.2 39.9 30.0 56.2 87.0 63.9 91.6 85.1 72.5 95.4 38.9 40.0 65.5 56.3 73.4 62.8 63.3 70.7 70.5 63.4 70.6 69.0 59.1 95.8 81.4 68.1

Analyzing the FTIR spectra in Figure 5, it is observed that the band in 1710 cm-1 corresponding to the stretching from C=O bond of acids gradually increases and is more pronounced on the third reuse, proving that there was not oleic acid full consumption in these esterification reactions. The band in 1742 cm-1 assigned to C=O bond from esters is more pronounced in product obtained in first reuse and gradually decreases. The loss of catalyst activity may 6/8

Figure 5. FTIR from products obtained by esterification using PSF_PSS electrospun membranes in three reuse cycles.

be due to a reduced surface area and reduced acid site concentration[19]. In the first reuse, the catalyst showed a higher catalytic activity and this leads to the conclusion that this catalyst presents possibilities for reuse.

4. Conclusions In this paper, catalytic activity from six polymer membranes of PSF and PSF_PSS were investigated in esterification of oleic acids with methanol, in different reactional times. FTIR has confirmed the formation of methylic esters and the consumption of oleic acid. And by using gas chromatography, it was possible to quantify methyl oleate formation as a reaction product. The results have shown that PSF_PSS membranes, having more acidic sites in their structure due to the presence of sulfonated polystyrene, had a better catalytic performance than PSF membranes. Having a larger superficial area, the electrospun membrane resulted in a conversion of 70.6% with only 10 minutes of reaction and 95.8% in 240 minutes. Thus, results have shown that polysulfone with sulfonated polystyrene membranes, in particular electrospun, may be successfully applied as heterogeneous catalyst acids in esterification reactions from oleic acid with methanol, aiming biodiesel production, mainly due to the good conversion achieved in short reaction times with the possibility of reusing the catalyst.

5. Acknowledgments The authors wish to thank the Coordination for the Improvement of Higher Education Personnel - Graduate Support Program (CAPES/PROAP), Brazilian National Council for Scientific and Technological Development (CNPq), Brazilian Study and Project Funding Agency (FINEP), Federal Institute of Goiás (IFG) and Minas Gerais Research Foundation (FAPEMIG).

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Lima A. P., Vieira A. T., Aud B. N., Batista A. C. F., Morais L. C., Faria A. M., Assunção R. M. N. & Pasquini D. methyl esters in biodiesel produced by microwave-assisted transesterification. International Journal of Environmental Sciences and Development, 6(12), 964-969. http://dx.doi. org/10.7763/IJESD.2015.V6.730. 30. Ruschel, C. F. C., Huang, C. T., Samios, D., & Ferrao, M. F. (2014). Exploratory analysis applied to attenuated total reflectance fourier transform infrared (ATR-FTIR) of biodiesel/


diesel blends. Quimica Nova, 37(5), 810-815. http://dx.doi. org/10.5935/0100-4042.20140130.

Received: Aug. 31, 2021 Revised: Sept. 25, 2021 Accepted: Oct. 19, 2021

Polímeros, 31(3), e2021027, 2021

ISSN 1678-5169 (Online)


ABS/recycled PCTG blend compatibilized with ionomer: effect on impact resistance and morphology Juliana Augusto Molari1*  and Deborah Dibbern Brunelli1  Instituto Tecnológico de Aeronáutica – ITA, São José dos Campos, SP, Brasil



Abstract The effect of adding ionomer as a compatibilizing agent in ABS/recycled PCTG blend was the objective of this study. Design of experiments using extreme vertices was used, to obtain a mathematical equation to predict the result of the impact resistance of blends, within a pre-established interval. The sample that obtained the highest impact resistance was the 79/20/1 (ABS/PCTG/Ionomer) and was analyzed by DSC and SEM. The results showed partial compatibility. Through the analysis of the fracture surface of the Charpy test specimen, it was verified that the PCTG, as a dispersed phase, presented itself in the form of fibers and the ionomer acted as an emulsifier. All results showed that it is possible to reuse PCTG industrial waste by mixing it with ABS and Ionomer as compatibilizing agent. Keywords: ABS, PCTG, ionomer, polymeric blend, compatibilization. How to cite: Molari, J. A., & Brunelli, D. D. (2021). ABS/recycled PCTG blend compatibilized with ionomer: effect on impact resistance and morphology. Polímeros: Ciência e Tecnologia, 31(3), e2021028. https://doi.org/10.1590/01041428.20210070

1. Introduction Plastic waste has become a worldwide environmental problem once it has been accumulated on the ecosystems across the globe [1,2]. Several plastics waste streams come from packaging, construction, and automotive industry [1,2]. In order to reduce the environmental impact of plastics, different technologies for plastic mechanical recycling have been developed [3]. The main method of enhancing the properties of the plastic waste materials consists of adding new components in the mixture: virgin polymer content, compatibilizers and stabilizers [3-5]. The acrylonitrile-butadiene-styrene (ABS) polymer is a terpolymer widely used in the automotive and home appliance industries. The properties of the ABS polymer vary according to the proportion of its monomers [6]. The PCTG copolyester is a copolymer formed by the esterification and polycondensation reaction of cyclohexanedimethanol (CHDM), terephthalic acid (TPA) and ethylene glycol (EG). If the CHDM content in the copolyester is less than 50 %, the copolymer is called PETG, and when it is greater than 50 %, it becomes the PCTG. PCTG is widely used in the packaging industry and is very susceptible to hydrolytic thermal oxidation, which can cause discoloration/yellowing and chain scission. For this reason, the mechanical recycling of this material is not widely recommended, which makes it very difficult to reuse PCTG after its life cycle [7-9]. ABS and PCTG polymers have different polarities and are not fully compatible, therefore obtaining a blend between them requires the use of compatibilizing agents [10].

Polímeros, 31(3), e2021028, 2021

Ionomers are polymers that have ionic groups in their molecular structure resulting from the neutralization of sulfonic acid or carboxylic acid groups and can be used as compatibilizing agents in polymer blends [11]. This work aims the study of the effectiveness of the ionomer matching agent Surlyn® in the blend between ABS and PCTG, in order to develop a method of reusing industrial waste from PCTG. Strategic planning and execution of experiments (Design of Experiments - DOE) were used to search for the optimum conditions for this multivariable system. For the modeling of mixtures, the data can be adjusted by simplex-lattice, simplex-centroid or extreme vertices. Modeling by extreme vertices is the most appropriate method, using pseudocomponent, if there are upper and lower limit constraints on components.

2. Experimental 2.1 Materials Materials used in this work included ABS polymer Terluran® GP22, supplied by Styrolution™, PCTG SKYGREEN® JN400 industrial waste from SK Chemicals™ (obtained from leftover scrap resulting from injection molded parts from packaging industry) and the ionomer Surlyn® PC 2000, supplied by DuPont™.

2.2 Experimental design and blend preparation MINITAB® software was used to elaborate the modeling of the experiments by pseudocomponents and extreme



Molari, J. A., & Brunelli, D. D. vertices of third degree. The software was also used to complete the analysis of the data by analysis of variance (ANOVA) for the mixture of the components: ABS (X1), PCTG (X2) and ionomer (X3).

3. Results and Discussion

Pseudocomponent modeling and extreme vertices were used for compatibilized blends, with mass percentage ranges of 30 % ≤ ABS ≤ 80 %, 20 % ≤ PCTG ≤ 60 % and 1 % ≤ Ionomer ≤ 7 %. The MINITAB® software generated 13 runs for each formulation where the Charpy impact resistance was determined experimentally for each run prior to model. Each central point of the experiment was repeated at least three times. The total ratio for each formulation adds up to a total of 100 % for a mass of 100 g, as presented in Table 1.

Table 2 and Figure 1 show the results of Charpy impact strength obtained for blends compatibilized with Ionomer.

Samples were first weighed, cold mixed, and then extruded using a twin screw extruder in order to ensure an adequate homogeneity. Prior to the injection molding of the test specimens, all extruded mixtures were dried in a dehumidifying dryer at 80 °C for 4 hours.

2.3 Differential scanning calorimetry measurements (DSC) The DSC experiments were conducted using Mettler’s model 822e. Measurements were performed using the second heating from -120 °C to 250 °C at a ratio of 20 °C min-1, under the dynamic atmosphere of nitrogen (50 mL min-1). Samples that had the highest Charpy impact resistance were analyzed by this technique.

2.4 Charpy impact test Impact strength tests were performed according to ASTM D6110:17, using a Resil 25R instrumented impact machine from Ceast using a 1.0 J impactor in injected test specimens at 23 °C, with a pendulum velocity of 2.90 m/s.

2.5 Scanning electron microscopy (SEM) A low vacuum Scanning electron microscopy, model FEI Quanta 400, was used to evaluate the non-metallized fracture surface of the samples.

3.1 Charpy impact resistance

According to the results (Table 2), it is possible to notice that the impact resistance tends to increase as the percentage of PCTG decreases, and the impact properties tend to be lower when either ABS or PCTG are around 50 %. Regression and analysis of variance (ANOVA) were performed to determine an equation that could predict final impact resistance of the blend in different proportions, within the percentage range of each component previously stipulated.

3.2 Analysis of mixture design results Table 3 presents the statistical data collected regarding the regression performed on the mixtures considering the variable Y1 response (impact resistance). By adjusting the regression, with 95 % of confidence, the special quadratic model was postulated in Equation 1 for Charpy impact resistance (Y1). The p-value can be used to assess whether or not the term is significant for the model. If the p-value is less than 0.05, at a confidence level of 95 %, the model term is considered significant. The R2 value indicates the percentage of response variation around the mean that is explained by the regression, while the adjusted R2 value, despite being similar to the R2 value, does not increase with the inclusion of independent variables that are not significant. The experimental data were adjusted using the special quadratic model, represented by the Equation 1. Y 1 = 136 X 1 + 93 X 2 + 3185 X 3 – 263 X 1X 2 – 3915 X 1X 3 – 3918 X 2 X 3 + 3315 ( X 1) ² X 2 X 3 + 3370 X 1( X 2 ) ² X 3 – 17655 X 1X 2 ( X 3) ²


Table 1. Blend compositions. X1 1 0.88 0.18 0.06 0.94 0.59 0.12 0.47 0.77 0.70 0.36 0.30 0.53 0.53 0.53


Pseudocomponent X2 0 0 0.82 0.82 0 0.41 0.82 0.41 0.20 0.20 0.61 0.61 0.41 0.41 0.41

X3 0 0.12 0 0.12 0.06 0 0.06 0.12 0.03 0.09 0.03 0.09 0.06 0.06 0.06

ABS 79 73 39 33 76 59 36 53 68 65 48 45 56 56 56

Components (% m/m) PCTG 20 20 60 60 20 40 60 40 30 30 50 50 40 40 40

IONOMER 1 7 1 7 4 1 4 7 3 6 3 6 4 4 4

Polímeros, 31(3), e2021028, 2021

ABS/recycled PCTG blend compatibilized with ionomer: effect on impact resistance and morphology where, X1, X2 and X3 are, respectively, the percentages by weight of ABS, PCTG and Ionomer and, YI represents the value of impact resistance. The higher the value of R2 and adjusted R2, the better the model fits the data and, therefore the proposed model for impact resistance explains 97.41 % of the data and these values are reliable at a level of 93.95 %, respectively. According to the model adopted for impact resistance, there is a different effect for each variable and interaction, as shown in Table 4. According to the results presented in Table 4, the addition of ABS, PCTG and Ionomer, in isolation, contribute positively to the equations obtained for impact resistance. Regarding the interactions between the components, only the interactions (ABS)2*PCTG*Ionomer and ABS*(PCTG)2*Ionomer have positive effects. However, only the interactions ABS*PCTG, ABS*Ionomer, and PCTG*Ionomer have statistical significance. Figures 2 shows the influence on the impact strength of the percentage variation of the blend components: ABS, PCTG and Ionomer and Figure 3 shows (a) response surface and (b) contour plots, presenting the effects on the impact resistance of the ABS/PCTG/Ionomer blends. The component effect plot represents the influence of each element at the central point of the experimental region on the response values.

According to Figures 2 and 3, in the weighted percentage range of 30 % ≤ ABS ≤ 80 %, 20 % ≤ PCTG ≤ 60 %, and 1 % ≤ Ionomer ≤ 7 %, it is observed that an increase in PCTG in the mixture is harmful to the impact resistance, as well as Ionomer increasing. However, for extrapolated values of Ionomer, the effect can be the opposite.

Table 2. Average impact resistance values for the ABS/PCTG/ Ionomer blends. ABS (%)

PCTG (%)

79 73 39 33 76 59 36 53 67 64 47 44 56 56 56

20 20 60 60 20 40 60 40 30 30 50 50 40 40 40

IONOMER (%) 1 7 1 7 4 1 4 7 3 6 3 6 4 4 4

Charpy Impact (J/m) 136 ± 8 91 ± 1 67 ± 17 51 ± 9 93 ± 1 51 ± 3 52 ± 10 38 ± 9 80 ± 14 56 ± 1 53 ± 3 36 ± 12 44 ± 12 50 ± 8 39 ± 13

Table 3. Regression performed for the response variable Y1 (impact resistance) in the ABS/PCTG/Ionomer blend.

Figure 1. Impact resistance results of ABS/PCTG/Ionomer blends in different proportions.


Coef 136.0 93.0 3,185.0 -263.0 -3,915.0 -3,918.0 3,315.0 3,370.0 -17,655.0 97.41% 93.95%

SE Coef 6.46 10.11 1,348.10 41.73 1,548.01 1,567.44 1,465.30 1,582.85 8,636.61

p-value 0.001 0.045 0.047 0.064 0.077 0.087

Coef: Coefficient; SE Coef: Standard error of the coefficient; R2: Coefficient of determination; Adjusted R2: Adjusted coefficient of determination.

Table 4. Effects of variables with Ionomer compatibilizer.

Figure 2. Influence on the impact strength of the percentage variation of the blend components: ABS, PCTG and Ionomer. Polímeros, 31(3), e2021028, 2021



ABS PCTG Ionomer ABS*PCTG ABS*Ionomer PCTG *Ionomer (ABS)2 *PCTG* Ionomer ABS*(PCTG)2*Ionomer ABS*PCTG*(Ionomer)2

Positive Positive Positive Negative Negative Negative Positive Positive Negative

Statistical significance (p-value <0.05) Significant Significant Significant Not Significant Not Significant Not Significant


Molari, J. A., & Brunelli, D. D.

Figure 3. (a) Response surface and (b) contour plots showing the effects on the impact resistance of the ABS/PCTG/Ionomer blends. Table 5. ANOVA for impact resistance (Y1), ABS/PCTG/Ionomer blend. ANOVA Regression Linear Quadratic ABS*PCTG ABS*IONOMER PCTG*IONOMER Spacial Cubic (ABS)2*PCTG*IONOMER ABS*(PCTG)2*IONOMER ABS*PCTG*(IONOMER)2 Residual error Lack of adjustment Pure error Total

DF 8 2 3 1 1 1 3 1 1 1 6 4 2 14

Seq SS 9,896.0 5,520.2 4,100.7 4,020.3 23.3 57.1 275.2 68.9 22.9 183.3 263.3 205.3 58.0 10,159.3

Adj SS 9,896.05 983.48 1,949.28 1,740.36 280.69 274.17 275.15 224.52 198.94 183.34 263.25 205.26 57.99 -

Adj MS 1,237.01 491.74 649.76 1,740.36 280.69 274.17 91.72 224.52 198.94 183.34 43.88 51.32 28.99 -

F 28.19 11.21 14.81 39.67 6.40 6.25 2.09 5.12 4.53 4.18 1.77 -

p-value 0 0.009 0.004 0.001 0.045 0.047 0.203 0.064 0.077 0.087 0.392 -

DF: degree of freedom; Seq SS: sum of the sequential squares; Adj SS: adjusted sum of squares; Adj MS: mean of squares; F: F-test.

Table 5 presents the analysis of variance (ANOVA) for impact resistance. Each SS (Sequential or adjusted) is associated with a number of degrees of freedom (D.F), which indicates the number of independent values involving the “n” observations that are necessary to determine it. According to the results presented, the Adj SS value of the regression is higher than the Adj SS value of the residuals. It indicates that the fraction described by the regression is more representative than the fraction described by the residuals, since the Adj SS of the residuals provides information about the part of the response variation that the model cannot reproduce. The quadratic sum of the residuals can be decomposed into quadratic sum due to the pure error and also due to the lack of adjustment. 4/8

Thus, for the model proposed for impact resistance, the value of Adj SS pure error is lower than for the lack of adjustment, which demonstrates that the pure error associated with the results is small and the model has a variation in the adjustment. The adjusted quadratic mean of the residuals (Adj MS residuals) can be interpreted as an “average error” (quadratic) that is made when using the regression equation to predict a response. From the data presented in Table 5, it is possible to verify how much the proposed model fits with the collected data. For this, it is necessary to satisfy requirements 1-3 (Table 6), since Ftab is the tabulated value of Fisher-Snedecor that is 4.15 for F (8; 6) 19.25 for F (4; 2) in this case: It was found that the proposed model for impact resistance fits very well with the data collected, since all the requirements have been met and the residues are randomly dispersed. Polímeros, 31(3), e2021028, 2021

ABS/recycled PCTG blend compatibilized with ionomer: effect on impact resistance and morphology Table 6. Requirements for adjusting the proposed model to the data. Requirement 1)

AdjMS Regresion > FTab (D.F. AdjMS residue


AdjMS lack of adjustment < FTab (D.F. AdjMS pure error

Regression; D.F. residue) lack of adjustment; D.F. pure error)

Charpy Impact 28.19 > 4.15 True 1.77 < 19.25


3) Residue plot with random pattern

DF: degree of freedom

Figure 4. Experimental value vs. calculated value, using the proposed model for impact resistance with Ionomer compatibilizer.

Figure 4 summarizes the experimental values of the impact resistance and the calculated values using the model proposed in Equation 1. It can be concluded that, through the proposed model, it was possible to predict the final properties of the blend, since the theoretical values were very close to the experimental results.

3.3 Compatibility analysis ABS/PCTG/Ionomer 79/20/1 blend which presented the highest impact resistance result was selected to be analyzed by DSC and Scanning Electron Microscopy (SEM) in order to evaluate the compatibility between the components. Neat components, ABS and PCTG, as well as the blend 79/20/1 were analyzed by DSC in order to determine the glass transition temperature. The results are shown in Figure 5. One of the criteria to evaluate the miscibility in polymer blends is the analysis of the glass transition temperature [1214] . A blend can be considered miscible when there is only one glass transition temperature, depending on the blend composition. On the other hand, for a partially miscible blend, two or more glass transitions can be observed and attributed to the phases of the blend. It can be emphasized the glass transitions of the phases are shifted relative to that of the neat components. In this case, each phase consists of a miscible mixture containing different compositions [12-14]. Polímeros, 31(3), e2021028, 2021

Figure 5. DSC results for the analyzed samples.

Finally, for immiscible polymer blends, there is no intermediate glass transition temperature. In this case, the glass transition temperatures of the phases are close to those of the neat components [12-15]. According to the results shown in Figure 5, it is clear that the Tg of 79/20/1 ABS/PCTG/Ionomer blend, attributed to the PCTG-rich phase, was shifted to higher temperatures. The same behavior was observed for the ABS-rich phase. However, the displacement was more subtle, indicating that the ABS/PCTG/Ionomer blend is partially miscible. The fracture surface of the Charpy impact resistance specimen was analyzed to explore the compatibility mechanism. An incompatible blend consists of a continuous phase and a dispersed phase which presents larger particles the greater the immiscibility between the components of the blend. This phenomenon occurs due to the coalescence of dispersed phase when there is a high interfacial tension between the components [13,14]. On the other hand, when two immiscible polymers are mixed using compatibilizing agent, the interfacial tension between the dispersed and continuous phase is decreased and the coalescence phenomenon decreases. Compatibilizer can change the morphology of blend as smaller sites of the dispersed phase can be observed randomly distributed throughout the matrix that will improve final mechanical properties [13,14]. Figure 6 shows the micrographs of 79/21/1 ABS/PCTG/ Ionomer blend. It can be seen that PCTG-rich phase is dispersed 5/8

Molari, J. A., & Brunelli, D. D. as fibers, indicated by the arrows in the images. This may have occurred due to the presence of the compatibilizing agent that reduced the interfacial energy and prevented the particles from coalescing. As a result of the compatibilizing effect, the adhesion in the interface region became more effective, as it can be observed that most fibers are broken and not detached. This analysis corroborates the high impact resistance result discussed above. Although PCTG and PETG are different polymers, they have a chemical similarity. Zhang et al. [16] studied the compatibility between PP/PETG blends using different compatibilizing agents. In his work, the tendency for PETG to be dispersed in the form of fibers in the matrix was identified and the result of impact resistance was associated with this behavior. Similar morphology was observed in some other works [16-18]. This may have occurred due to the presence of the compatibilizing agent which reduced the interfacial energy and prevented the particles from coalescing. As a result of the

compatibilizing effect, the adhesion in the interface region became more effective, as it can be observed that most fibers are broken and not detached. This analysis corroborates the high impact resistance result discussed above. Regarding the mechanism of action of the ionomer as compatibilizer in the ABS/PCTG blend, it is suggested that it behaved as an emulsifier agent in the mixture and the compatibilization may have occurred through an acidolysis reaction, where R is the aliphatic carboxylic acid from the ionomer (Figure 7). Samios and Kalfoglou [19] and Dekoninck et al. [20] also observed this same mechanism, for blends between ABS/ PETG and ABS/PET, respectively (Equation 2). 2PCTG + R − COONa → PCTG − COONa + PCTG − R (2)

Regarding the interaction between ABS and the Ionomer, it is suggested that the group -COONa of the ionomer tends to remove electrons from the styrenic group of ABS, which tends to release electrons [19,20], as shown in Figure 8. This

Figure 6. Micrograph of the fracture surface of the ABS/PCTG/Ionomer sample (79/20/1) with magnification of (a) 5000 x (b) 10000 x.

Figure 7. Possible compatibility reaction by acidolysis between PCTG/Ionomer in the blend between ABS/PCTG/Ionomer. 6/8

Polímeros, 31(3), e2021028, 2021

ABS/recycled PCTG blend compatibilized with ionomer: effect on impact resistance and morphology

Figure 8. Possible reaction of compatibility between ABS/Ionomer in the blend between ABS/PCTG/Ionomer.

increasing in the mixture, as well as the increase of the ionomer proportion. The blend composition that presented the best performance regarding the impact resistance was 79/21/1 (ABS/PCTG/Ionomer). It was possible to establish a viable equation to predict Charpy impact resistance for the blend between the components within the concentration range used in this study through the modeling of the experiments. The equation can be useful in determining the most suitable proportions to mix the components in order to predict the final impact resistance of the blend and ensure the feasibility to reuse/ recycle industrial PCTG waste. The morphological analysis of the fracture surface and DSC analysis showed that the blend 79/20/1 had a partial miscibility. The PCTG as a dispersed phase presented the form of fibers with reduced sizes and homogeneous dispersion. It is suggested that the ionomer acted as an emulsifier, interacting with both components of the blend.

5. Acknowledgements Authors thank to ITA for the support provided for the preparation of the study.

6. References Figure 9. Scheme of the morphology and mechanism of action of the compatibilizer for 79/20/1 blend (ABS/PCTG/Ionomer).

theory was also observed by Ismail and Nasir [21] for blends between polystyrene and polypropylene, using ionomer as a compatibilizer. Figure 9 summarizes the compatibilization action of the Ionomer in the ABS/PCTG blend morphology. It can be concluded that the ionomer acted as a good compatibilizing agent both in the ABS-rich phase and in the PCTG-rich phase, since PCTG was presented as small-sized fibers homogeneously dispersed in the matrix. Therefore, it was possible to ensure partial miscibility of the ABS/PCTG/Ionomer blend improving its performance on impact resistance tests.

4. Conclusion According to the results, it is possible to conclude that the impact resistance decreases with the PCTG content Polímeros, 31(3), e2021028, 2021

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Molari, J. A., & Brunelli, D. D. 6. Joseph, S., Focke, W. W., & Thomas, S. (2010). Compatibilizing action of a poly(styrene–butadiene) Triblock co-polymer in ABS/PET-G blends. Composite Interfaces, 17(2-3), 175-196. http://dx.doi.org/10.1163/092764410X490590. 7. Heo, Y. M., Koo, J. M., Hwang, D. K., JaeGal, J. G., Hwang, S. Y., & Im, S. S. (2016). Synthesis and characteristics of biobased copolyester for thermal shrinkage film. RSC Advances, 6(62), 57626-57633. http://dx.doi.org/10.1039/C6RA10333B. 8. Chen, T., Jiang, G., & Zhang, J. (2014). Isothermal crystallization behavior and crystal structure of poly(ethylene terephthalateco-1,4-cyclohexylene dimethylene terephthalate) (P(ET/CT)) copolyesters. Crystal Research and Technology, 49(4), 232-243. http://dx.doi.org/10.1002/crat.201300369. 9. Cook, W. D., Moad, G., Fox, B., Deipen, G. V., Zhang, T., Cser, F., & McCarthy, L. (1996). Morphology-property relationships in ABS/PET blends. II. Influence of Processing conditions on structure and properties. Journal of Applied Polymer Science, 62(10), 1709-1714. http://dx.doi.org/10.1002/(SICI)10974628(19961205)62:10<1709::AID-APP22>3.0.CO;2-V. 10. Chen, T., & Zhang, J. (2018). Compatibilization of acrylonitrile-butadiene-styrene terpolymer/poly(ethylene glycol-co-1,4-cyclohexanedimethanol terephthalate) blend: effect on morphology, interface, mechanical properties and hydrophilicity. Applied Surface Science, 437, 62-69. http:// dx.doi.org/10.1016/j.apsusc.2017.12.168. 11. Koning, C., Duin, M., Pagnoulle, C., & Jerome, R. (1998). Strategies for compatibilization of polymer blends. Progress in Polymer Science, 23(4), 707-757. http://dx.doi.org/10.1016/ S0079-6700(97)00054-3. 12. Olabisi, O., Robeson, L. M., & Shaw, M. T. (1979). Polymerpolymer miscibility. New York: Academic Press. http://dx.doi. org/10.1016/B978-0-12-525050-4.X5001-X. 13. Utracki, L. A. (2003). Introduction to polymer blends. Dordrecht: Springer. http://dx.doi.org/10.1007/0-306-48244-4_1. 14. Paul, D. P., & Bucknall, C. B. (1999). Polymer blends: formulation and performance. New York: John Wiley & Sons.


15. Thirtha, V. M., Lehman, R. L., & Nosker, T. J. (2004). Morphological effects on glass transitions in immiscible polymer blends. MRS Online Proceedings Library, 856, 1112. http://dx.doi.org/10.1557/PROC-856-BB11.12. 16. Zhang, X., Li, B., Wang, K., Zhang, Q., & Fu, Q. (2009). The effect of interfacial adhesion on the impact strength of immiscible PP/PETG blends compatibilized with triblock copolymers. Polymer, 50(19), 4737-4744. http://dx.doi. org/10.1016/j.polymer.2009.08.004. 17. Li, B., Zhang, X., Zhang, Q., Chen, F., & Fu, Q. (2009). Synergistic enhancement in tensile strength and ductility of ABS by using recycled PETG plastic. Journal of Applied Polymer Science, 113(2), 1207-1215. http://dx.doi.org/10.1002/ app.30002. 18. Chen, T., & Zhang, J. (2016). Surface hydrophilic modification of acrylonitrile-butadiene-styreneterpolymer by poly(ethylene glycol-co-1,4-cyclohexanedimethanolterephthalate): preparation, characterization, and properties studies. Applied Surface Science, 388, 133-140. http://dx.doi.org/10.1016/j.apsusc.2016.02.242. 19. Samios, C. K., & Kalfoglou, N. K. (2001). Acrylic-modified polyolefin ionomers as compatibilizers for poly(ethylene-covinyl alcohol)/aromatic copolyester blends. Polymer, 42(8), 3687-3696. http://dx.doi.org/10.1016/S0032-3861(00)00432-8. 20. Dekoninck, J. M., Legras, R., & Mercier, J. P. (1989). Nucleation of poly(ethylene terephthalate) by sodium compounds: a unique mechanism. Polymer, 30(5), 910-913. http://dx.doi. org/10.1016/0032-3861(89)90191-2. 21. Ismail, H., & Nasir, M. (2002). The effect of various Compatibilizers on mechanical properties of polystyrene/ polypropylene blend. Polymer Testing, 21(2), 163-170. http:// dx.doi.org/10.1016/S0142-9418(01)00064-2. Received: Sept. 16, 2021 Revised: Oct. 18, 2021 Accepted: Oct. 24, 2021

Polímeros, 31(3), e2021028, 2021

ISSN 1678-5169 (Online)


Structural changes of polyethylene in blown films with different pro-oxidants João Augusto Osório Brandão1* , Fernando Dal Pont Morisso2 , Edson Luiz Francisquetti3  and Ruth Marlene Campomanes Santana1  Departamento de Materiais, Universidade Federal do Rio Grande do Sul – UFRGS, Porto Alegre, RS, Brasil 2 Laboratório de Estudos Avançados de Materiais, Universidade Feevale, Novo Hamburgo, RS, Brasil 3 Departamento de Materiais, Instituto Federal do Rio Grande do Sul, Farroupilha, RS, Brasil



Abstract The accumulation of polymeric residues has been one of the most impacting environmental problems in recent human history, coming, above all, from disposable artefacts, such as plastic bags. Processing polyolefins with pro-oxidant additives is an alternative to favour the abiotic degradation process of macromolecules, including thermooxidation, so that the oxygenated fragments produced can be assimilated by microorganisms. The objective of this work was to evaluate the process of thermomechanical oxidative degradation of polyethylene (PE) during tubular extrusion of HDPE/LDPE films, without and with 1% of two different pro-oxidants, d2wTM and benzoin. The results of viscosimetric and MFI analyses indicated smaller chain sizes in the additivated films. The FTIR spectra and contact angles indicate a higher presence of polar functional groups in the samples with pro-oxidants. The surface morphological analysis by SEM indicated difference of PE homogeneity in the films. Benzoin, however, proved to be a better pro-oxidant than d2wTM. Keywords: benzoin, d2wTM, polyethylene, pro-oxidants, thermooxidation. How to cite: Brandão, J. A. O., Morisso, F. D. P., Francisquetti, E. L., & Santana, R. M. C. (2021). Structural changes of polyethylene in blown films with different pro-oxidants. Polímeros: Ciência e Tecnologia, 31(3), e2021029. https://doi.org/10.1590/0104-1428.20210058

1. Introduction The use of polymeric materials has been growing worldwide since the 1940s, replacing the use of metals, ceramics and wood in many industrial branches[1-3]. Parallel to this, a major environmental problem regarding the use of this class of materials arises: the accumulation of plastic waste in the environment (soil, rivers and oceans) due to incorrect and unconscious disposal, especially of disposable items, such as plastic bags, sacks, cups and bottles[3-5]. Due to their low specific masses, easier processing and low cost, petrochemical resins are the most widely used and are also the most difficult to degrade, including PE[6-10], which plays an important role in the largest volume of plastic waste. An alternative to solve the environmental problem would be the use of biodegradable polymers, which are macromolecules that can be cleaved by the action of biological enzymes of microorganisms (fungi, bacteria and algae) and subsequently used as nutrients for the growth of colonies, provided in the appropriate environmental conditions[11,12]. Thus, biodegradable polymers are returned to the environment as gaseous compounds and salts, such as CO2, H2O, CH4, depending on the presence or absence of oxygen, in a process called mineralization[13]. Biodegradable polymers, however, are more expensive and difficult to process, which hinders their use when compared to petrochemical resins. Besides this, when they are low cost, they are not applicable to the

Polímeros, 31(3), e2021029, 2021

required purpose, due to the absence of some property, generally mechanical[14]. Another, more viable option would be the use of oxobiodegradable polymers, which are petrochemical resins processed with a pro-oxidant additive[15-17]. Pro-oxidant additives are responsible for favouring the abiotic degradation of the polymer, mainly by thermooxidation and photodegradation, with production of oxygenated fragments of lower molar mass that can be assimilated by microorganisms, a biotic process[18-20]. In general, they are organic salts of transition metals, mainly stearates of Fe, Co and Mn[20-25]. But organic pro-oxidants have been investigated, such as benzoin, which showed promising results in the abiotic degradation of polypropylene (PP)[26] and PE[27]. However, during polymer processing, as in the tubular extrusion process, the polymer is subjected to high shear rates and high temperatures when passing through the barrel, which can be initiators of the degradation process that, in the presence of O2 from air, can be called oxidative thermomechanical degradation. The present work is aimed at evaluating how the presence of pro-oxidant additives, one based on organic salts of transition metals (d2wTM) and another totally organic (benzoin), influence the thermomechanical oxidative degradation of PE during processing by tubular extrusion to obtain films. The films were evaluated by



Brandão, J. A. O., Morisso, F. D. P., Francisquetti, E. L. & Santana, M. C. dilute solution viscosimetry (DSV), with determination of viscosity average molar mass; Fourier-transform infrared spectroscopy (FTIR) and carbonyl index determination; flow index (MFI); contact angle and scanning electron microscopy (SEM).

2. Materials and Methods 2.1 Materials High density polyethylene (HDPE) grade HE-150, with MFI of 1.0 g/10 min (190 ºC/ 2.16 kg), and low density polyethylene (LDPE) grade EB-853/72, with MFI of 2.7 g/ min (190 ºC/ 2.16 kg), both produced by Braskem (Brazil), were used in this work. d2wTM pro-oxidant additive in masterbatch form with low density polyethylene (LDPE) as a base polymer, produced by Symphony Environmental (United Kingdom). Benzoin with a purity grade above 99%, produced by Merck KGaG (Germany). Decahydronaphthalene (Decalin) produced by Neon.

2.2 Obtaining films by tubular extrusion According to Table 1, the HDPE and LDPE blend, in mass proportion of 90 and 10%, respectively, was processed without and with 1 wt.% of pro-oxidants. The mixture was properly homogenized for processing. The blown films, with an average thickness of 30 μm, were obtained by tubular extrusion in a Seibt (Brazil) singlescrew extruder, model ES 35-FR, with 5 heating zones, using the following temperature profile: 120/150/175/185/210 ºC, from zones 1 (feed) to 5 (die).

2.3 Characterization 2.3.1 Dilute Solution Viscosimetry (DSV) and Viscosity Average Molar Mass ( Mv )

( )

In determining the viscosity average molar mass Mv , the dilute solution viscosimetry technique was used. Five PE solutions were prepared at concentrations of 0.2, 0.4, 0.6, 0.8 and 1.0 g/dL for each film analysed, using decalin as solvent. The solutions were obtained, one by one, by determining the mass needed to prepare 25 ml of each of the concentrations mentioned. The dissolution of the polymer took place under stirring and at 140oC for one hour. The viscosities of the solutions were measured in a CannonFenske viscometer (no 50) with an internal capillary diameter of 0.44 mm. The procedures were carried out according to ASTM D445 and ASTM D446. The analyses were performed with the viscometer immersed in a thermostatic silicone oil bath of SOLAB brand, Model SL 150, at 135.0 ± 0.1ºC[28]. First, the relative viscosities were calculated. Subsequently, the reduced specific (ηesp/C) and inherent viscosities (ln ηrel/C)

of each one of the solutions were determined. Plotting the graphs of such viscosities versus concentration, the intrinsic viscosity of the PE used in the preparation of the solutions was determined, from the extrapolation of the straight lines obtained by linear regression when the concentration tends to zero. The values found for the two straight lines tend to the same value and, for this reason, their average was used as intrinsic viscosity ([η]). For the determination of Mv of PE, the Mark-Houwink-Sakurada equation was used, which relates the intrinsic viscosity ([η]) and the viscosity average molar mass Mv , as presented in Equation 1:

( )

( )

η  = k . Mv



The α and K constants were 0.7 and 62 x 10-5 dL/g, respectively[28]. The films evaluated are made of a mixture of HDPE and LDPE, i.e., both PE. As per the literature reference, the viscosimetric constants used are for PE, without distinction among its variations (whether HDPE, LDPE or LLDPE, for example). 2.3.2 Melt Flow Index (MFI) The tests for determining the PE flow index of HDPE/ LDPE films with and without pro-oxidant additives were performed in the CEAST modular MeltFlow plastomer equipment, Model 7026.000, according to method A of ASTM D1238. The conditions used were 190 ºC/ 2.16 kg, with a residence time of 420 seconds and time between cuts of 60 seconds, with a total test time of 900 seconds for each sample. 2.3.3 Fourier Transform Infrared Spectroscopy (FTIR) and Carbonyl Index (CI) To investigate the changes in the chemical structure of the films, especially the appearance of new functional groups due to oxidation, the FTIR analysis was performed in Perkin Elmer equipment, Frontier model, according to ASTM E1131. The films were evaluated in the ATR (Total Attenuated Reflectance) mode and the spectra obtained at a controlled ambient temperature of 25ºC, air humidity of 30% and in an absorption region ranging from 4000 to 650 cm-1, with 10 scans for each sample analysed. From the spectra, the carbonyl indices (CI) were calculated. The poorly variable absorption band peak was adopted as that at 1463 cm-1[29], and the limits between 1468 and 1450 cm-1 (A1468-1450) were integrated. The absorption peaks of carbonyls adopted to verify the degree of PE oxidation were of esters and carboxylic acids (1300-1050 cm-1) and lactones (1780-1770 cm-1), and the limits between the adopted bands were integrated, named A1300-1050 and A1780-1770[25,30], respectively. The CI was calculated from Equation 2: = CI

( A1300−1050 +

A1780 −1770 ) / A1468−1450 (2)

Table 1. Mass composition of the HDPE/LDPE films. Sample

HDPE/LDPE (90/10) (wt. %)


100 99 99




2.3.4 Contact angle

(wt. %) 1

(wt. %) 1 -

The contact angle test was performed using deionized water as the liquid, according to ASTM D7334-08. The analysis allows the determination of the substrate (film) hydrophilicity or hydrophobia. For each film, 10 repetitions were made, in Polímeros, 31(3), e2021029, 2021

Structural changes of polyethylene in blown films with different pro-oxidants which drops of water were placed on the surface. The images were obtained from a Knup microscope, model Kp-8012, and the contact angles determined from the Surftens 4.3 software for 3 seconds and 3 minutes. 2.3.5 Scanning Electron Microscopy (SEM) The surface morphologies of the samples were obtained by scanning electron microscopy soon after processing by tubular extrusion, in a Jeol instrument, model JSM6510LV. The films were metallized with gold in a standard procedure in the Denton Caccum metallizer, model Desk. Micrographs were obtained with electron beams at 10 kV energy and magnifications of 1500 and 5000x.

3. Results and Discussions After processing, the samples were characterized to evaluate the influence of the pro-oxidant on the process of thermomechanical oxidative degradation of PE during extrusion of the blown films.

( )

3.1 Evaluation of Mv and MFI The intrinsic viscosity ([η]) of the PE was determined for each of the samples and is shown in Table 2. From the Mark-Howink-Sakurada equation (Equation 1), the Mv of the PE in the evaluated films were determined, presented in Figure 1-a. In Figure 1-b, the MFI results of the evaluated films are presented.

In Figure 1-a is possible to observe a decrease of Mv of the PE in the films extruded with the pro-oxidants if compared to the film without additives. The presence of the additives led to higher rates of chain scission and, consequently, a reduction in molar mass, indicating a greater degradation process of the polymer[31,32]. The PE scission process was more intense in the presence of benzoin than for d2wTM, since the PE molar mass reduction in the PE_ONM film was higher than in the PE_OM film. The MFI values, in Figure 1-b, indicate that there is a greater difficulty in the flow of the molten PE macromolecules in the PE_Pure film, possibly due to their larger size compared to the films containing pro-oxidants. The PE_ONM film containing benzoin, on the other hand, showed higher MFI, indicating higher fluidity and lower viscosity of the molten polymer, compatible with the lower Mv of the sample and for which there is higher mobility, possibly due to the presence of smaller chains size and less likelihood of entanglement. The PE_OM film showed MFI with intermediate value, compatible with its Mv . It can be seen that there is a coherent correlation between the Mv and the MFI, since the reduction of the Mv is followed by an increase in MFI, indicating with the PE with smaller chains size has higher flow when melted, observed for the PE_ONM film. The opposite is also true, i.e., the PE in the PE_Pure sample, with higher Mv , has the lowest flow when melted, hampered by the larger size of the polymer chains.

3.2 FTIR Table 2. Intrinsic viscosity ([η]) of PE in the samples evaluated. Sample PE_Pure PE_OM PE_ONM

[η]1* 1.5185 1.3812 1.2376

[η]2** 1.5651 1.4873 1.4336

[η]f*** 1.5418 1.4342 1.3356

*obtained of graph ηsp.red versus C; ** obtained graph ηinehr versus C; *** average intrinsic viscosity of [η]1 and [η]2.

Figure 2-a shows the FTIR spectra, between the 2000 and 1500 cm-1 bands, of the PE films just after processing by tubular extrusion. It is possible to observe that all samples present peaks between 1780 and 1700 cm1 , which indicates that the PE, in all samples, suffered oxidation during processing, with formation of oxidized products, such as esters, ketones, aldehydes, lactones and carboxylic acids(25, 34). Visually, the PE_Pure film presents lower peaks, mainly at 1780 and 1712 cm-1. The PE_OM

Figure 1. Properties of the evaluated samples: (a) Mv , (b) MFI. Polímeros, 31(3), e2021029, 2021


Brandão, J. A. O., Morisso, F. D. P., Francisquetti, E. L. & Santana, M. C. film presents well defined peaks at 1712, 1723, 1730 and 1780 cm-1, indicating that the d2wTM additive favoured autooxidation. The PE_ONM film presents a clear peak at 1780 cm-1, suggesting that benzoin favoured PE autooxidation, however, the peak at 1723 cm-1 is overlap with the 1730 cm-1 peak, emphasizing that benzoin has a ketone group in its structure and, for this reason, the production of fragments containing esters and carboxylic acids groups was considered, observable in Figure 2-b, FTIR spectra, between the 1400 and 1000 cm-1, in which it is possible to observe that there is production of these oxygenated fragments in all samples. Visually, the PE_ ONM film presents peak broadening between 1300 and 1250 cm-1, as well as between 1150 and 1100 cm-1, when compared to the other two films, which suggests a higher degree of oxidation. The PE_OM film, when compared to PE_Pure film, presents wider and deeper peak between

1100 and 1050 cm-1. However, it was decided to discuss the oxidation degree of the samples by the CI evaluation. Using Equation 2 and the Origin 8.5.1 software, the CI values for the samples were calculated and are shown in Figure 3. The different CI values indicate unequal degrees of oxidation of the PE in the samples during the extrusion process, favoured by exposure to high temperatures and high shear rates in the presence of O2 from the air[33,34].

In the case of the samples evaluated, the presence of carbonyl groups is evident in all, indicating that the PE suffered oxidation, producing, among others, carboxylic acids, esters, and lactones[15,19,20,34]. However, the CI value is higher for PE_ONM and PE_OM. Comparing only the CI values, benzoin and d2wTM showed equal efficiency in accelerating PE autooxidation by thermooxidation, since the index values are similar.

Figure 2. FTIR spectra of the evaluated samples: (a) between 2000 and 1500 cm-1, (b) between 1400 and 1000 cm-1.

Figure 3. Carbonyl index of the samples. 4/8

Polímeros, 31(3), e2021029, 2021

Structural changes of polyethylene in blown films with different pro-oxidants 3.3 Contact angle Another way to confirm the presence of functional groups is through the contact angle. Figure 4 shows the images of the water droplets on the surface of the films, in which it is possible to visualize the differences between the contours of the droplets and the contact angles for each of the films. It can be seen that the droplet on the PE_ONM sample (Figure 4-c) is more spread on the surface than the others, indicating that its surface is more hydrophilic. The average surface contact values of the films of the evaluated samples are shown in Figure 5. The results found indicated that the PE_Pure film is the most apolar, with the largest contact angle, consistent with its low oxidation rate, if considered the CI value. The PE_OM film, is the second

most apolar, with small reduction of the contact angle value, if compared to the PE_Pure film. The reduction of the contact angle is possibly due to the greater insertion of carbonyls in the polymer chain, which changed its polarity. The PE_ONM film showed a marked reduction in the contact angle value compared to the other two films. Besides the insertion of carbonyl groups, the film also presented a higher concentration of hydroxyls (-OH), capable of establishing hydrogen bonds with water and reducing the contact angle presenting, therefore, the lowest value. The contact angle reduction indicates the PE degradation process in films additivated with pro-oxidants, with insertion of hydrophilic functional groups[35]. Benzoin, besides catalyzing the degradation of PE, altered the film surface by the insertion of hydroxyls in its structure.

Figure 4. Images of deionized water droplets arranged on the analysed films: (a) PE_Pure; (b) PE_OM; (c) PE_ONM.

Figure 5. Contact angle of the evaluated samples. Polímeros, 31(3), e2021029, 2021


Brandão, J. A. O., Morisso, F. D. P., Francisquetti, E. L. & Santana, M. C.

Figure 6. SEM images of the samples: (a) PE_Pure – 1500 x; (b) PE_Pure – 5000 x; (c) PE_OM – 1500 x; (d) PE_OM – 5000 x; (e) PE_ONM – 1500 x; (f) PE_ONM – 5000 x.

3.4 SEM The micrographs obtained to evaluate the surface morphology of the samples evaluated are presented in Figure 6 and in them it is possible to observe differences that indicate that the processing of PE, with and without pro-oxidant additives, leads to structural changes in the extruded films. The film PE_Pure presents the highest surface homogeneity, while the film PE_ONM presents irregularities, possibly due to the presence of the polar chemical agent, benzoin, which hindered a better dispersion of the apolar polymer. The PE_OM film presents morphological variations with the appearance of agglomerations, which not indicate a total dissolution of the masterbatch containing the additive, whose processing took place in a single-screw extruder. Initially, these alterations only point out that the different 6/8

organizations of the macromolecules may favour an increase in the fragility of the films for subsequent uses.

4. Conclusions The tubular extrusion process is characterized by exposure of the polymer to high shear rates associated with high temperatures, and in this study, PE with pro-oxidants had its degradation accelerated, increasing the rate of scission and oxidation of the macromolecules. The benzoin pro-oxidant accelerates this degradation more markedly compared to the d2wTM additive, as it led to a greater reduction in the Mv of the polymer evaluated. The structural changes in PE observed in this work need to be considered when using pro-oxidant additives to obtain oxybiodegradable PE, since they can alter important characteristics of a plastic artifact produced from it, such as disposable bags, reducing its shelf life. Polímeros, 31(3), e2021029, 2021

Structural changes of polyethylene in blown films with different pro-oxidants

5. Acknowledgements Thanks to The National Council for Scientific and Technological (CNPq) for the financial support. Thanks to Federal Institute of Rio Grande do Sul (IF/RS) – Campus Farroupilha for carrying out FTIR analyses. Thanks to Feevale University for the SEM analyses.

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Journal of Polymers and the Environment, 19(3), 637-676. http://dx.doi.org/10.1007/s10924-011-0317-1. 14. Siracusa, V., Rocculi, P., Romani, S., & Rosa, M. D. (2008). Biodegradable polymer for food: a review. Trends in Food Science & Technology, 19(12), 634-643. http://dx.doi.org/10.1016/j. tifs.2008.07.003. 15. Ojeda, T. F. M., Dalmolin, E., Forte, M. M. C., Jacques, R. J. S., Bento, F. M., & Camargo, F. A. O. (2009). Abiotic and biotic degradation of oxo-biodegradable polyethylenes. Polymer Degradation & Stability, 94(6), 965-970. http://dx.doi. org/10.1016/j.polymdegradstab.2009.03.011. 16. Sen, S. K., & Raut, S. (2015). Microbial degradation of low density polyethylene (LDPE): a review. Journal of Environmental Chemical Engineering, 3(1), 462-473. http:// dx.doi.org/10.1016/j.jece.2015.01.003. 17. Santos, A. S. F., Freire, F. H. O., Costa, B. L. N., & Manrich, S. (2012). Sacolas plásticas: destinações sustentáveis e alternativas de substituição. Polímeros: Ciência e Tecnologia, 22(3), 228237. http://dx.doi.org/10.1590/S0104-14282012005000036. 18. Chiellini, E., Corti, A., D’Antone, S., & Baciu, R. (2006). Oxo-biodegradable carbon backbone polymers – Oxidative degradation of polyethylene under accelerated test conditions. Polymer Degradation & Stability, 91(11), 2739-2747. http:// dx.doi.org/10.1016/j.polymdegradstab.2006.03.022. 19. Liu, X., Gao, C., Sangwan, P., Yu, L., & Tong, Z. (2014). Accelerating the degradation of polyolefins through additives and blending. Journal of Applied Polymer Science, 131(18), 9001-9015. http://dx.doi.org/10.1002/app.40750. 20. Reddy, M. M., Gupta, R. K., Gupta, R. K., Bhattacharya, S. N., & Parthasarathy, R. (2008). Abiotic oxidation studies of oxo-biodegradable polyethylene. Journal of Polymers and the Environment, 16(1), 27-34. http://dx.doi.org/10.1007/s10924008-0081-z. 21. Roy, P. K., Surekha, P., Raman, R., & Rajagopal, C. (2009). Investigating the role of metal oxidation state on the degradation behavior of LDPE. Polymer Degradation & Stability, 94(7), 10331039. http://dx.doi.org/10.1016/j.polymdegradstab.2009.04.025. 22. Jakubowicz, I. (2003). Evaluation of degradability of biodegradable polyethylene (PE). Polymer Degradation & Stability, 80(1), 39-43. http://dx.doi.org/10.1016/S0141-3910(02)00380-4. 23. Corti, A., Muniyasamy, S., Vitali, M., Imam, S. H., & Chiellini, E. (2010). Oxidation and biodegradation of polyethylene films containing pro-oxidant additives: synergistic effects of sunlight exposure, thermal aging and fungal biodegradation. Polymer Degradation & Stability, 95(6), 1106-1114. http:// dx.doi.org/10.1016/j.polymdegradstab.2010.02.018. 24. Vogt, N. B., & Kleppe, E. A. (2009). Oxo-biodegradable polyolefins show continued and increased thermal oxidative degradation after exposure to light. Polymer Degradation & Stability, 94(4), 659-663. http://dx.doi.org/10.1016/j. polymdegradstab.2009.01.002. 25. Babetto, A. S., Agnelli, J. A. M., & Bettini, S. H. P. (2015). Evaluation of the pro-degradant systems in the thermooxidative degradation of HDPE. Polímeros: Ciência e Tecnologia, 25(No. esp.), 68-76. http://dx.doi.org/10.1590/0104-1428.2022. 26. Montagna, L. S., Catto, A. L., Forte, M. M. C., Chiellini, E., Corti, A., Morelli, A., & Santana, R. M. C. (2015). Comparative assessment of degradation in aqueous medium of polypropylene films doped with transition metal free (experimental) and transition metal containing (commercial) pro-oxidant/prodegradant additives after exposure to controlled UV radiation. Polymer Degradation & Stability, 120, 186-192. http://dx.doi. org/10.1016/j.polymdegradstab.2015.06.019. 27. Rosa, T. P. S. (2019). Polietileno modificado com pró-degradante orgânico para aplicação em embalagens flexíveis (Master’s Thesis). Universidade Federal do Rio Grande do Sul, Porto Alegre. 7/8

Brandão, J. A. O., Morisso, F. D. P., Francisquetti, E. L. & Santana, M. C. 28. Brandup, J., Immergut, E. H., & Grulke, E. A. (2009). Polymer Handbook. 2nd ed. New York: Wiley-Interscience Publication. 29. Gulmine, J. V., Janissek, P. R., Heise, H. M., & Akcelrud, L. (2002). Polyethylene characterization by FTIR. Polymer Testing, 21(5), 557-563. http://dx.doi.org/10.1016/S0142-9418(01)00124-6. 30. Sugimoto, M., Shimada, A., Kudoh, H., Tamura, K., & Seguchi, T. (2013). Product analysis for polyethylene degradation by radiation and thermal ageing. Radiation Physics and Chemistry, 82, 69-73. http://dx.doi.org/10.1016/j.radphyschem.2012.08.009. 31. Albertsson, A.-C., Erlandsson, B., Hakkarainen, M., & Karlsson, S. (1998). Molecular weight changes and polymeric matrix changes correlated with the formation of degradation products in biodegradable polyethylene. Journal of Environmental Polymer Degradation, 6(4), 187-195. http:// dx.doi.org/10.1023/A:1021873631162. 32. Khajehpour-Tadavani, S., Nejabat, G.-R., & Mortazavi, S. M. M. (2018). Oxo-biodegradability of high-density polyethylene films containing limited amount of isotactic polypropylene. Journal of Applied Polymer Science, 135(6), 45843. http:// dx.doi.org/10.1002/app.45843.


33. Agnelli, J. A. M., & Chinelatto, M. A. (1992). Degradação de polipropileno: aspectos teóricos e recentes avanços em sua estabilização. Polímeros: Ciência e Tecnologia, 2(3), 27-31. Retrieved in 2021, August 1, from https://revistapolimeros. org.br/journal/polimeros/article/588371317f8c9d0a0c8b4792 34. Gardette, M., Perthue, A., Gardette, J.-L., Janecska, T., Földes, E., Pukanszky, B., & Therias, S. (2013). Photo and thermaloxidation of polyethylene: comparison of mechanisms and influence of unsaturation content. Polymer Degradation & Stability, 98(11), 2383-2390. http://dx.doi.org/10.1016/j. polymdegradstab.2013.07.017. 35. Muthukumar, T., Aravinthan, A., & Mukesh, D. (2010). Effect of environment on the degradation of starch and pro-oxidant blended polyolefins. Polymer Degradation & Stability, 95(10), 1988-1993. http://dx.doi.org/10.1016/j. polymdegradstab.2010.07.017. Received: Aug. 05, 2021 Revised: Oct. 07, 2021 Accepted: Oct. 25, 2021

Polímeros, 31(3), e2021029, 2021

ISSN 1678-5169 (Online)


Antimicrobial activity of silver composites obtained from crosslinked polystyrene with polyHIPE structures Roberta Trovão Santos1 , Nathália Smith Santos1 , Mirian Araújo de Oliveira1 , Fernanda de Andrade Buás Campeão1 , Maria Aparecida Larrubia Granado Moreira Rodrigues Mandu2 , Mônica Regina Costa Marques3  and Luciana da Cunha Costa1*  Programa de Pós-graduação em Ciência e Tecnologia Ambiental, Centro Universitário Estadual da Zona Oeste – UEZO, Rio de Janeiro, RJ, Brasil 2 Departamento de Controle de Qualidade de Águas, Companhia Estadual de Águas e Esgoto – CEDAE/RJ, Rio de Janeiro, RJ, Brasil 3 Programa de Pós-graduação em Química, Universidade do Estado do Rio de Janeiro – UERJ, Rio de Janeiro, RJ, Brasil 1


Abstract The literature reports several potential applications of polymers prepared with high internal phase emulsions (HIPEs). However, the evaluation of these materials as supports for antimicrobial agents has not been explored. In this work, silver composites based on polyHIPEs were prepared. Initial studies indicated that these materials can be efficient to prevent biofilm formation. The silver composites were prepared in three steps. First, HIPEs based on styrene-divinylbenzene were polymerized by aqueous suspension polymerization. These particles showed surface areas of 18 and 48 m2/g. These polyHIPEs were sulfonated with concentrated sulfuric acid or acetyl sulfate and showed cation exchange capacities of 4.03 and 5.07 meq/g respectively. The sulfonated material was impregnated with silver ions, followed by reduction of the ions to prepare silver composites. These composites showed inhibition halos against E. coli and P. aeruginosa. and did not present adhesion of bacterial cells of K. variicola and S. aureus on their surface. Keywords: high internal phase emulsions, polyHIPEs, silver composites, styrene-divinylbenzene copolymers, biocidal polymers. How to cite: Santos, R. T., Santos, N. S., Oliveira M. A., Campeão, F. A. B., Mandu, M. A. L. G. M. R., Marques, M. R. C., & Costa, L. C. (2021). Antimicrobial activity of silver composites obtained from crosslinked polystyrene with polyHIPE structures. Polímeros: Ciência e Tecnologia, 31(3), e2021030. https://doi.org/10.1590/0104-1428.20210005

1. Introduction The literature indicates that high internal phase emulsions (HIPEs) (containing more than 74 vol% of internal phase) can be employed as templates for porous polymers, generating polyHIPE structures[1]. Despite a large number of publications and patents involving the preparation of polyHIPE structures, only a few of these works involve polymerization of these HIPEs by aqueous suspension polymerization (Hainey et al.[2]; Cameron et al.[3]; Desforges et al.[4]; Stefanec and Krajnc[5,6]; Yang et al.[7]; Mert and Hildirim[8]; Cui et al.[9]; Torquato et al.[10]) and US patents 5,583,162[11], 6,048,908[12], 6,100,306[13] and 6,218,440[14]). The polymerization of HIPEs by employing aqueous suspension polymerization, also referred as waterin-oil-in-water polymerization (W/O/W), can generate spherical particles with an open porous morphology with pores connecting larger cavities (windows)[10]. This type of morphology can favor access of reactants and products through the internal structure of the polymer[15,16], contributing to reduction of mass transfer limitations, especially intraparticle diffusion, and consequently allow the development of more

Polímeros, 31(3), e2021030, 2021

efficient materials. The application of these polyHIPEs is widespread in many areas[15-18], such as supports for catalysts[15], CO2 capture[16], stationary phase for preconcentration of PAHs in environmental water samples[17], and development of enzymatic reactors based on immobilization of enzymes on these polymers[18]. Crosslinked porous copolymers have been used as supports for several antimicrobial groups such as ammonium and phosphonium quaternary groups, charge transfer complexes involving iodine and quaternary ammonium groups, N-halamines, sulfo-derivatives and metal particles[19]. Antimicrobial polymers based on silver composites and nanocomposites have potential applications for the inhibition of biofilm formation[20]. Among the studies involving the preparation of silver composites from crosslinked copolymers can be cited the works of Gangadharan et al.[21] Santa Maria et al.[22], Mthombeni et al.[23] and Mandu et al.[20]. In these studies, the antimicrobial activity of the final products was evaluated by different strategies. Gangadharan et al.[21] developed a silver nanocomposite through impregnation



Santos, R. T., Santos, N. S., Oliveira, M. A., Campeão, F. A. B., Mandu, M. A. L. G. M. R., Marques, M. R. C., & Costa, L. C. of silver nanoparticles in microspheres of methacrylic acid and divinylbenzene, prepared by aqueous suspension polymerization. The antimicrobial activity of this material was evaluated against E. coli, P. aeruginosa, B. subtilis and S. aureus employing the plate method and in test tubes. Santa Maria et al.[22], Mthombeni et al.[23] and Mandu et al.[20] prepared silver composites from commercial sulfonic resins (Lewatit VPOC1800, Amberlyst 15WET, Amberlite IR-120, respectively). The method of impregnation of silver nanoparticles in these resins was based on exchange of H+ ions by Ag+ ions followed by chemical reduction of these ions. Santa Maria et al.[22] evaluated the antimicrobial activity of these materials against E. coli in tests employing micro syringes. In turn, Mthombeni et al.[23] evaluated the antimicrobial activity of these silver nanocomposites against E. coli in macroscale column experiments. Mandu et al.[20] evaluated the antimicrobial activity of silver composites against E. coli, P. aeruginosa and S. aureus through plate, batch, and column experiments. From the column experiments, also realized on a macro scale, they determined the breakthrough point and working biocidal activity. However, in these works, no tests were performed to simulate the ability of the final products to inhibit biofilm formation. Literature data show that polyHIPEs have not been investigated as supports for antimicrobial groups. Thus, here we report a strategy to prepare silver composites from polyHIPEs of styrene-divinylbenzene (Sty-DVB) containing sulfonic groups. The antimicrobial activity of these silver composites was evaluated by employing inhibition halo tests against two Gram-negative bacteria of medical importance (Escherichia coli and Pseudomonas aeruginosa). The silver composites were also evaluated regarding bacterial adsorption and adhesion employing the bacteria Klebsiella variicola (Gram-negative) and Staphylococcus aureus (Gram-positive) to simulate inhibition on biofilm formation. The sulfonation of polyHIPEs has been sparsely studied, by employing concentrated sulfuric acid[24,25] or lauroyl sulfate in cyclohexane[26] as sulfonating agents. Thus, in

the present work the sulfonation of the polyHIPE particles was studied by using concentrated sulfuric acid or acetyl sulfate as sulfonating agents (in a comparative study), aiming to learn the effect of the type of sulfonating agent on the morphology of the polyHIPEs and cation exchange capacity of the sulfonated particles. All steps of this work are shown schematically in Figure 1.

2. Materials and Methods 2.1 Materials Styrene (Sty) and divinylbenzene (DVB) (donated by Nitriflex Indústria e Comércio S.A.) were used after acid-base extraction by using 5% w/v NaOH aqueous solution. 2,2-azobis-isobutyronitrile (AIBN) was acquired from Mig Quimica and used after recrystallization with methanol. Poly(vinyl alcohol) (PVA) 224 (hydrolysis degree of 87-89%) was donated by Kurary Inc. and polyvinylpyrrolidone (PVP) (Mw 1,300.000) was acquired from Fluka. The microorganisms used in this work were Escherichia coli (ATCC 11229), Pseudomonas aeruginosa (ATCC 15442), Klebsiella variicola (ATCC 31488) and Staphylococcus aureus (ATCC 25923). The culture media used were Mueller-Hinton II broth (MHII) (BBL), Mueller-Hinton II agar (Kasvi), and plate count agar (PCA) (Neogen). Other reagents and solvents were purchased from Sigma-Aldrich and used as received.

2.2 Synthesis of Sty-DVB polyHIPEs First, high internal phase emulsions (HIPEs) of Sty-DVB were prepared in a round-bottom three-necked flask with 250 cm3 capacity, fitted with a mechanical stirrer reactor. The organic phase was composed of Sty (0.072 mol), DVB (0.029 mol), AIBN (0.0009 mol) and Span 80 (0.0053 mol). The aqueous phase (AP) was composed of NaCl (0.67 g), potassium persulfate (0.45 g) and distilled water (33 mL). The aqueous phase was dripped through an addition funnel into the organic phase (35 minutes) under mechanical stirring

Figure 1. Schematic representation of all stages of the experiment. 2/10

Polímeros, 31(3), e2021030, 2021

Antimicrobial activity of silver composites obtained from crosslinked polystyrene with polyHIPE structures at 250 rpm. In the second stage, HIPEs (as viscous white solutions) were polymerized by aqueous suspension, forming polyHIPEs structures. The aqueous phase was prepared using PVP (1.98 g) or PVA (1.98 g) as suspending agent, NaCl (1.49 g) and distilled water (99.4 cm3), and then transferred to a round-bottom three-necked flask with 500 cm3 capacity. The HIPE was added to the aqueous phase (pre-heated to 80 ºC) dropwise for 30 minutes under mechanical stirring at 150 rpm. The reaction mixture remained for 24 hours in these conditions at 80 ºC[4-6, 10]. After preparation, the particles were treated with hot water, ethanol and acetone to remove residues. After washing, the polyHIPEs were oven dried for 24 hours at 60 °C.

2.3 Sulfonation of Sty-DVB polyHIPEs The polyHIPE composed of Sty-DVB prepared with PVP as stabilizer was sulfonated with sulfuric acid or acetyl sulfate to promote silver ion anchorage in the silver impregnation step (Figure 1). The first sulfonation was carried out with sulfuric acid. PolyHIPE (4.0 g) was previously soaked in 1.2 dichloroethane (18.5 cm3) for 2.5 hours. After this period, the material was sulfonated with sulfuric acid (95-99%) (1.13 mol) in a 500 cm3 flask coupled to a mechanical stirrer and reflux condenser, and was stirred at a rate of 90 rpm at 90 °C for 24 hours. Previous swelling of the particles by soaking with 1.2 dichloroethane allowed greater access of the reactant sulfuric acid to the internal structure of the polymeric beads, aiming to preserve their spherical morphology[27]. The second sulfonation method was based on employment of acetyl sulfate as sulfonating agent[28]. This reagent was prepared employing acetic anhydride (1.13 mol), anhydrous dichloromethane (0.39 mol) and concentrated sulfuric acid (1.13 mol) under N2 atmosphere. PolyHIPE (4.0 g) was swollen in anhydrous dichloromethane (30 cm3) for 1 hour under N2 atmosphere in a round-bottom flask. After this period, the sulfonating agent prepared in the previous step was added to the flask. The reaction mixture was kept at a constant temperature of 40 °C for 24 hours under stirring at 90 rpm. The sulfonated polyHIPEs were filtered under reduced pressure, washed with distilled water to neutral pH and oven dried at 60 °C for 24 hours.

2.4 Impregnation of silver in sulfonated Sty-DVB polyHIPEs The impregnation of silver nanoparticles in polymers was performed according to the descriptions of Mandu et al.[20], Santa Maria et al.[22] and Gangadharan et al.[21], with modifications. Both sulfonated polyHIPEs (1 g) were treated with AgNO3 (aqueous solution of 0.1 mol L-1, 10 mL) for 48 hours in a sealed environment protected from light. After filtration and washing with water, the reduction of Ag+ to Ag0 ions was performed by applying 4.7 g of hydroxylamine hydrochloride and 2 mol L-1 NH4OH solution (to maintain the pH at 12.0) and a solution of distilled water (50 cm3), gelatin (1.5 g) and hydroxyethyl cellulose (1.5 g) (Figure 1). The mixture of gelatin and HEC was used as colloid protector, aiming to reduce the reaction of Ag+ ions in the polymer matrix in controlled form[29]. The composite produced was washed thoroughly with deionized distilled water (60 °C), then ethanol (50 cm3), and dried at 60 °C for 24 hours. Polímeros, 31(3), e2021030, 2021

2.5 Characterization of the polymers Sty-DVB polyHIPEs were characterized regarding apparent density (ASTM D1895-69)[30] by optical microscopy (Olympus BX51M) and scanning electron microscopy (SEM) (JEOL-JSM 6460 LV), and the specific area, pore volume and pore diameter were determined through nitrogen physisorption (Micromeritics ASAP 2020 apparatus) following the BET and BJH equations[31], and the particle size distribution was ascertained (Malvern Hydro 2000S model). Sulfonation was accompanied by determination of cation exchange capacity[32]. The FT-IR spectra of the all polymers (KBr pellets) were recorded with a Perkin-Elmer Spectrum One spectrometer (4 scans, 4 cm-1 resolution). TG and DTG curves of the copolymers were obtained using a TA Q50 instrument in the temperature range of 50–600 ºC (20 º C min-1) under N2 atmosphere (60 mL min-1).

2.6 Evaluation of antimicrobial activity Antimicrobial evaluation of silver composites was performed by inhibition halo tests to verify the ability of polymers to inhibit growth of the bacteria E. coli and P. aeruginosa[33]. Petri dishes (100 mm diameter) were filled with 30 cm3 of MH II agar. Bacterial suspensions in MHII medium with 1x108 CFU/mL were used to inoculate the entire sterile surface of the agar in the Petri dishes using sterile swabs. The polymers were then inserted into 5 mm diameter holes in the inoculated MHII agar. The dishes were incubated in a culture oven at 37 ºC for 24 hours. The test was performed in triplicate. After the incubation period, the halos of the microbial inhibition zones were measured, corresponding to the shortest distance between the outer surface of the well and the beginning of the microbial growth region. They were measured with a millimeter ruler. Silver composites also were evaluated in relation to bacterial adsorption and adhesion[21] aiming to evaluate potential applications to inhibit biofilm formation. This study was conducted by employing suspensions of K. variicola and S. aureus. Suspensions of these bacteria (106 CFU/mL) were inoculated using nutrient medium in test tubes. Silver composites (0.2 g) were also transferred to these tubes and kept in contact with the bacterial suspensions for 24 h at 45 ºC. After this, the silver composite was filtered by using sterile syringes containing filter paper and washed with sterile deionized water. The silver composite was plated on nutrient agar (PCA medium) and incubated for 48 h at 45 ºC to check bacterial adsorption/adhesion. For statistical analysis, the PAST v2.17 data analysis package was used and the Kruskal-Wallis test for nonparametric data was applied, with p <0.05 being considered statistically significant.

3. Results and Discussions 3.1 Synthesis of polyHIPEs HIPEs based on styrene-divinylbenzene (Sty-DVB) were suspended in an aqueous suspension phase containing as stabilizer poly(vinyl alcohol) (PVA) or polyvinylpyrrolidone (PVP) to prepare polyHIPEs. Optical microscopic images of these polyHIPEs (Figure 2) showed the presence of spherical particles and undefined shapes formed by using both stabilizers. 3/10

Santos, R. T., Santos, N. S., Oliveira, M. A., Campeão, F. A. B., Mandu, M. A. L. G. M. R., Marques, M. R. C., & Costa, L. C. The formation of agglomerates by using both stabilizers was also observed, even after intense washing of the particles with hot water. The polyHIPE N3, prepared by using PVA as stabilizer, showed a larger distribution of particles sizes (Figure 3c) than the polyHIPE N2, synthesized by using PVP as stabilizer. The optical microscopic images of the polyHIPE N3 revealed the presence of smaller particles and agglomerations, and also a small group of larger particles, some spherical and others stick-shaped (Figure 2d and Figure S1). However, the particle size curve for this material (Figure 3c) did not show a bimodal size distribution for these particles, not corroborating the optical microscopic data. According to Budhlall et al.[34] PVA and PVP adopted varied conformations on the aqueous phase, related to the mechanism of solvating the polymeric chains, which is dependent on molecular mass and hydrolysis degree of these stabilizers. These variations cause significant modifications in particle nucleation mechanisms during the polymerization process. Considering this, it is also possible to suppose that both the process of dispersion up and coalescing in the early stage of the polymerization, and the agglomeration process during formation of the macroscopics structures (generated by w/o/w polymerization) were influenced by the conformations of the stabilizers in the aqueous phase. From a simpler point of view, it is possible to suppose that higher viscosity of the aqueous solution prepared using PVP

as a stabilizer (Mw 1,300.000) generated more compact droplets during early stages of the polymerization, in turn generating more compact particles. SEM micrographs of these materials (Figure 2) indicated that the polymer prepared with the stabilizer PVP had a more porous structure. This observation was confirmed by data of surface area and pore volume of these two materials (Table 1). Probably more voids between polymeric domains were formed when PVP was used as the stabilizer, generating more porous structures[35]. The influence of the type of stabilizer on the morphology of the polyHIPEs prepared by aqueous suspension has not yet been explored in literature. Thus, more studies are necessary to confirm these data. The polyHIPE particles showed type IV isotherms (Figure 3)[31,36]. It was visually clear that hysteresis was low between the adsorption and desorption processes, associated with capillary condensation on mesopores structures. The initial portion of this isotherm is related with monolayermultilayer adsorption[31]. The low content of N2 adsorbed at low relative pressure indicates that this material did not contain micropores in its structure[36]. The hysteresis loop profiles can be classified as H3 type (or B type) (considered intermediate between hysteresis type H1 and H4). This type of hysteresis loop is characteristic of aggregates of plate-like particles, giving rise to slit-shaped pores[31].

Figure 2. (a, b) SEM micrographs of the polyHIPEs N2 and N3, respectively. Magnifications of 250x; (c, d) Optical microscopic images of the polyHIPEs N2 and N3, respectively. Magnifications of 50x.

Figure 3. (a) N2 adsorption-desorption isotherms of polyHIPE N2; (b) N2 adsorption-desorption isotherms of polyHIPE N3; (c) Particle size distribution of polyHIPE N2 and N3. Table 1. Data on apparent density, surface area, pore volume and pore diameter of polyHIPEs of Sty-DVB N2 and N3. PolyHIPEs N2 N3

Stabilizer PVP PVA

dap (cm3/g) 0.36 0.27

S (m2/g) 48.4 18.3

Vp (cm3/g) 0.32 0.25

D (Å) 214.9 272.5

dap: apparent density; S: specific surface area; Vp: pore volume; D: pore diameter; PVA: polyvinyl alcohol; PVP: polyvinyl pyrrolidone.


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Antimicrobial activity of silver composites obtained from crosslinked polystyrene with polyHIPE structures The morphology of the polyHIPEs prepared by polymerization of HIPEs via aqueous suspension has not been sufficiently studied by N2 adsorption. Comparison between the results found in this work and data reported by Cui et al.[9] shows similarity between the N2 adsorptiondesorption isotherms. However, the microspheres prepared by Cui et al.[9] had specific surface areas varying between 3 and 4 m2/g whereas the microspheres synthesized in this study had surfaces areas of 18 and 48 m2/g. Probably this difference was related to variations in monomeric composition. PolyHIPEs obtained in this work were composed of Sty and DVB while the polyHIPEs prepared by Cui et al.[9] were composed of Sty, DVB, stearyl methacrylate (SMA) and n-butyl acrylate (BA). It is plausible to suppose that alterations of monomer polarity will generate modifications in the phase separation mechanism and the phase separation stage, generating beads with lower porosity[35].

3.2 Sulfonation of polyHIPEs PolyHIPE N2 prepared with PVP as stabilizer was chosen for sulfonation reactions due to its higher specific surface area and pore volume. It was submitted to sulfonation with concentrated sulfuric acid or acetyl sulfate in order to introduce SO3H groups in the polymeric matrix. After sulfonation reaction the particles showed color ranging from dark brown to black. Oliveira et al.[28] reported that acetyl sulfate is a better sulfonating agent than sulfuric acid because it contributes to maintaining the spherical morphology of

the particles. In this work, alterations of morphology of the particles related to the type of sulfonating agent were not clearly observed by optical microscopy, since both stabilizers generated undefined particles in addition to spherical particles and agglomerates. We observed that part of the polyHIPEs kept their spherical morphology even after sulfonation with sulfuric acid (Figure S2). FT-IR analysis of polyHIPEs after sulfonation with both sulfonating agents (Figure 4) revealed the presence of new bands at 1216 cm-1, attributed to asymmetric stretching vibration of the S=O bonds in the sulfonate group (H3O+SO3- group)[28,37], besides a pronounced increase of bands at 3400 cm-1 due to hydrogen bonding between sulfonic groups (polar) and water molecules (moist)[37], confirming the occurrence of both sulfonation reactions. These new bands were more accentuated in the FT-IR spectrum of polyHIPE sulfonated with sulfuric acid, indicating that this sulfonation method was more thorough. After sulfonation with both sulfonating agents, the polyHIPEs showed three stages of decomposition (Table 2), related to water loss, decomposition of sulfonic groups and degradation of the entire carbon chain[38]. The content of residue was higher in polyHIPE sulfonated with sulfuric acid, also indicating that this sulfonation was more efficient. PolyHIPE sulfonated with sulfuric acid showed higher cation exchange capacity (CEC) than polyHIPE sulfonated with acetyl sulfate (Table 3). This result agrees with those obtained by FT-IR and thermogravimetry and indicates that

Figure 4. FT-IR spectra of unmodified polyHIPE (polyHIPE N2), polyHIPE sulfonated with acetyl sulfate (polyHIPE N2-AS (a)) and polyHIPE sulfonated with sulfuric acid (polyHIPE N2-SA (b)). Table 2. Thermal characteristics of the polyHIPE N2, polyHIPE N2 sulfonated with sulfuric acid (polyHIPE N2-SA), polyHIPE N2 sulfonated with acetyl sulfate (polyHIPE N2-AS), silver composites derived from these two materials (based on work of Simplicio et al.[38]). Material

Tonset(1) (º C)a

Tonset(2) (º C)b

Tonset(3) (º C)c

RCd (%)




polyHIPE N2








polyHIPE N2-SA








polyHIPE N2-AS







polyHIPE N2-SA-Ag





polyHIPE N2-AS-Ag














Tonset related with the first stage of decomposition, b Tonset related with the second stage of decomposition, c Tonset related with the third stage of decomposition, d RD: residue content after last stage of the decomposition, e RD: Residue difference, difference of residue content between modified and unmodified polyHIPE, f Inorganic residue: Content of inorganic residue determined by difference between residue of composite and sulfonated materials, g Resistance to degradation%: Increase of thermal resistance at 700 ºC (comparing sulfonated polyHIPEs and silver composites derived from these materials): resistance %: (residue% composite – residue% sulfonated material)/(residue% sulfonated material) x 100. a

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Santos, R. T., Santos, N. S., Oliveira, M. A., Campeão, F. A. B., Mandu, M. A. L. G. M. R., Marques, M. R. C., & Costa, L. C. the sulfonation reaction with concentrated sulfuric acid generated structures with higher content of accessible sulfonic groups. Considering that the sulfonation with acetyl sulfate was carried out in a milder condition than the sulfonation with sulfuric acid, it is possible to suppose that the degree of sulfonation with acetyl sulfate was lower, i.e., sulfonation with sulfuric acid enabled introducing more sulfonic groups in the innermost layers of the particles. Oliveira et al.[28] also observed that polymers sulfonated with sulfuric acid showed higher CEC than polymers sulfonated with acetyl sulfate. Table S1 shows CEC values of sulfonic resins synthesized by other authors with commercial sulfonic resins[28,32,39-41]. The CEC value found for polyHIPE sulfonated with acetyl sulfate (PolyHIPE N2-AS) was more than three times higher than the CEC of the Sty-DVB copolymer sulfonated with the same sulfonating agent prepared by Oliveira et al. (RM-10)[28]. The CEC of polyHIPE sulfonated with sulfuric acid (PolyHIPE N2-AS) was higher than that of sulfonic resins prepared by Reis et al.[32], Aguiar et al.[39], Rezende et al. [40] and equivalent to the CEC of the sulfonic resin prepared by Coutinho et al.[37] and several commercial sulfonic resins used in industrial processes, such as Amberlyst 18[39], Amberlyst XN 1010[39], Amberlyst 35[39,41] and Amberlyst 36[32,39,40]. Sulfonation reactions were conducted after swelling of the beads in 1,2-dichloroethane (a good solvent of the polymeric chains) in order to avoid cracking the beads during sulfonation and also to favor diffusion of the sulfonating agents through the structure of the polymers[27]. PolyHIPE structures are characterized by the presence of interconnecting pores and larger cavities formed during polymerization of an external phase containing monomers dispersed in an aqueous phase[10]. Probably this type of structure favored the diffusion of the sulfonating agents through the internal structure of polymers, contributing to the higher CEC values observed.

nanoparticles into the polymeric structure[20-23]. EDS spectra of silver composites showed peaks associated with elemental silver located between 2.6 KeV and 3.6 KeV (Figure 5). The presence of these peaks confirmed the reduction of Ag+ ions present in [Ag(NH3)2]+ groups, associated with SO3- groups, to Ag0[21,22]. It was possible to observe small variations in the degradation profile of the composites polyHIPE N2-SA-Ag (derived from polyHIPE sulfonated with sulfuric acid) and polyHIPE N2-AS-Ag (derived from polyHIPE sulfonated with acetyl sulfate) (Table 2). For polyHIPE N2-AS-Ag, the first stage of degradation occurred at 116 ºC (1.6% of mass lost) while for polyHIPE N2-SA-Ag, the first step of decomposition occurred at 74 ºC (4.8% of mass lost). This first stage of decomposition can be related with water loss, while the second and third stages are related to the degradation of sulfonic groups and polymeric matrix, respectively[38]. The silver composite polyHIPE N2-SA-Ag had more residue (40.4%) than polyHIPE N2-AS-Ag (24.4%). Resistance to degradation of these composites was calculated considering the difference between residue content of the composites and sulfonated polyHIPEs. The values found were 41.3% (composite polyHIPE N2-SA-Ag) and 34.1% (composite polyHIPE N2-AS-Ag)[38]. These results allow suggesting that the composite derived from polyHIPE sulfonated with sulfuric acid contains more inorganic residue (Ag0 content) resistant to decomposition.

3.4 Evaluation of antibacterial activity of polymers According to Table 4, all silver composites produced inhibition halos against Pseudomonas aeruginosa and Escherichia coli (Gram-negative bacteria), demonstrating the composites’ ability to inhibit bacterial growth around silver particles. Statistical analysis of these data, performed by one-way analysis of variance, demonstrated that the

3.3 Preparation of silver composites PolyHIPEs sulfonated with acetyl sulfate and sulfuric acid were submitted to impregnation with silver particles in two steps: treatment of sulfonic particles with aqueous solution of AgNO3 and reduction of Ag+ to Ag0 employing hydroxylamine hydrochloride and NH4OH, in the presence of gelatin and hydroxyethyl cellulose (Figure 1). Some researchers have indicated that this method introduces silver

Table 3. Cation exchange capacity (CEC) of the polyHIPE sulfonated with acetyl sulfate (polyHIPE N2-AS), polyHIPE sulfonated with sulfuric acid (polyHIPE N2-SA). Supports

Sulfonating agents

CEC (meq/SO3H/g)


acetyl sulfate



Sulfuric acid


Figure 5. (a) EDS spectra of polyHIPE sulfonated with sulfuric acid; (b) EDS spectra of the polyHIPE. 6/10

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Antimicrobial activity of silver composites obtained from crosslinked polystyrene with polyHIPE structures Table 4. Halo values of polymers tested on Mueller-Hinton agar against Pseudomonas aeruginosa and Escherichia coli and variable correlations by applying the Kruskal-Wallis test. Halo values

Polymer Pseudomonas aeruginosa Escherichia coli

N2 N2-AS-Ag N2-SA-Ag N2 N2-AS-Ag N2-SA-Ag

Inhibition halo (mm) 0 2 2 0 2 2

0 3 2 0 1 2

Variable correlations Mean

0 3 2 0 2 2

0 2.7 2.0 0 1.7 2.0

Standard deviation 0.0 0.6 0.0 0.0 0.6 0.0

N2* 0.178 0.1406 0.178 0.1406

N2-AS-Ag N2-SA-Ag 0.05935 0.5629 0.05935 1

0.04685 0.1876 0.04685 0.505 -

N2: PolyHIPE of styrene-divinylbenzene, N2-AS-Ag: Composite derived from polyHIPE N2 sulfonated with acetyl sulfate and impregnated with silver, N2-SA-Ag: Composite derived from polyHIPE N2 sulfonated with sulfuric acid and impregnated with silver. * p<0.05: statistically significant.

Figure 6. Silver composite derived from the commercial resin Dowex: (a) control; (b) presence of Staphylococcus aureus; (c) presence of Klebsiella variicola in bacterial adsorption and adhesion tests. Silver composites PoliHIPE N2: (d) control; (e) absence of Staphylococcus aureus; (f) absence of Klebsiella variicola in bacterial adsorption and adhesion tests.

difference in zone diameter was significant (p<0.05) only in the comparison between the composites derived from polyHIPE sulfonated with sulfuric acid and unmodified polyHIPEs. Tests showed that unmodified polyHIPEs and sulfonated materials did not have any antimicrobial activity, corroborating results previously published by Maria et al.[22]. Silver composites prepared from polyHIPEs sulfonated with sulfuric acid showed equal inhibition halos for the two bacteria (E. coli and P. aeruginosa), both Gram-negative. Data from thermogravimetry indicated that the composite derived from polyHIPE sulfonated with sulfuric acid presented higher content of silver than the composite derived from polyHIPE sulfonated with acetyl sulfate. The difference in zone diameter was not significant (Table 4). Silver composites also were evaluated in relation to bacterial adsorption and adhesion by employing K. variicola (Gram-negative) and S. aureus (Gram-positive) (Figure 6). A silver composite derived from the commercial resin Dowex was used as reference in this test. It was possible to observe bacterial growth in the dishes containing composites derived Dowex. Dishes containing unmodified polyHIPEs and sulfonated materials also showed bacterial growth. Polímeros, 31(3), e2021030, 2021

On the other hand, this bacterial growth was not observed in dishes containing silver composites, indicating absence of adsorption/adhesion of bacterial cells on the surface these materials. This test indicated that this material could prevent biofilm formation.

4. Conclusions Silver composites were prepared by impregnation of silver particles in polyHIPEs of styrene-divinylbenzene. These polyHIPEs, with spherical shape, were prepared by aqueous suspension polymerization of high internal phase emulsions (HIPEs). Both particles produced had structures with macropores and low specific surface areas (18 and 48 m2/g). PolyHIPEs were sulfated with sulfuric acid or acetyl sulfate and the sulfonation was confirmed by EDX, FT-IR and TGA. The cation exchange capacity (CEC) values of polyHIPEs sulfonated with sulfuric acid and acetyl sulfate were 4.03 and 5.07 meq g-1 respectively, indicating that sulfonation with sulfuric acid was more efficient. These sulfonated polymers presented similar CEC values in relation to commercial sulfonic resins 7/10

Santos, R. T., Santos, N. S., Oliveira, M. A., Campeão, F. A. B., Mandu, M. A. L. G. M. R., Marques, M. R. C., & Costa, L. C. with macroreticular structures. Also, the silver composite prepared from polyHIPE sulfonated with sulfuric acid had higher content of silver than that prepared with polyHIPE sulfonated with acetyl sulfate. Both composites presented similar inhibition halos against the bacteria E. coli and P. aeruginosa (Gram-negative bacteria). These composites also did not present adsorption/adhesion of bacterial cells of K. variicola (Gram-negative) and S. aureus (Gram-positive) on their surface, indicating these materials can be efficiently used to prevent biofilm formation.

5. Acknowledgements We thank Fundação de Amparo a Pesquisa do Estado do Rio de Janeiro (FAPERJ), Nitriflex, Engepol/COPPE/ UFRJ, CETEM/UFRJ and Laboratório de Caracterização Instrumental III/IQ/UERJ.

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Santos, R. T., Santos, N. S., Oliveira, M. A., Campeão, F. A. B., Mandu, M. A. L. G. M. R., Marques, M. R. C., & Costa, L. C.

Supplementary Material Supplementary material accompanies this paper. Figure S1. Optical microscopy of polyHIPE N3. Magnification of 50x Figure S2. Optical microscopies of the Sty-DVB polyHIPE N2 (a) before (b) after sulfonation reaction with sulfuric acid. Magnification of 50x Table S1. Cation exchange capacity (CEC) of the polyHIPE sulfonated with acetyl sulfate (N2-AS), polyHIPE sulfonated with sulfuric acid (N2-SA) and other sulfonic resins reported in the literature This material is available as part of the online article from https://www.scielo.br/j/po


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ISSN 1678-5169 (Online)


Effects of gamma radiation on nanocomposite films of polycaprolactone with modified MCM-48 Marcos Vinícius Paula1* , Leandro Araújo de Azevedo2 , Ivo Diego de Lima Silva3 , Glória Maria Vinhas3 , Severino Alves Junior2  Faculdade de Engenharia de Materiais, Universidade Federal do Pará – UFPA, Campus Ananindeua, Belém, Pará, Brasil 2 Laboratório de Terras Raras, Departamento de Química Fundamental, Universidade Federal de Pernambuco – UFPE, Recife, Pernambuco, Brasil 3 Laboratório de Materiais Poliméricos e Caracterização, Departamento de Engenharia Química, Universidade Federal de Pernambuco – UFPE, Recife, Pernambuco, Brasil



Abstract The aim of this investigation was to assess the effects of gamma radiation on nanocomposite films (NC films) formed by PCL (polycaprolactone) with MCM-48 nanoparticles (PCL/MCM-48) and PCL with MCM-48 NPs modified with (3-aminopropyl)triethoxysilane (APTES) (PCL/MCM-48-NH2). The nanocomposite films were obtained using the solvent casting method. After preparing the films, they were irradiated at 25 kGy in the presence of air and analyzed by X-ray diffraction (XRD), Fourier transform infrared spectroscopy (FT-IR), differential scanning calorimetry (DSC), thermogravimetric analysis (TGA), and scanning electron microscopy (SEM), transmission electronic microscopy (TEM), as well as for their mechanical properties. The exposure of NC films to gamma radiation at 25 kGy did not cause major changes in either thermal or mechanical properties such as tensile strength and modulus of elasticity. The results revealed that gamma radiation was a successful choice for the sterilization of these materials. Keywords: gamma radiation, MCM-48, mechanical properties, polycaprolactone. How to cite: Paula, M. V., Azevedo, L. A., Silva, I. D. L., Vinhas, G. M., & Alves Junior, S. (2021). Effects of gamma radiation on nanocomposite films of polycaprolactone with modified MCM-48. Polímeros: Ciência e Tecnologia, 31(3), e2021031. https://doi.org/10.1590/0104-1428.20210044

1. Introduction Polycaprolactone (PCL) is a synthetic polymer formed of hexanoate units linked by ester linkages[1,2]. It is a hydrophobic, semi-crystalline polymer, being biocompatible and biodegradable[3-5]. PCL is a biocompatible polymer[6], which is interesting for its use for implants[7], controlled drug release systems[8], scaffolds[9] and in food packaging[10]. PCL is degraded by microorganisms and enzymes[2,11]. This degradation may take years, depending on its molecular mass, crystallinity content and degradation conditions[2]. In the human body, PCL is degraded via hydrolysis of the ester bonds and eliminated from the body through the citric acid cycle[2]. For some applications, PCL does not have satisfactory thermomechanical properties[12]. One way to solve this lack of satisfactory thermomechanical properties is to add nanomaterials to the PCL, producing a final material called a nanocomposite[13,14]. Nanometric materials have a high surface area and volume, presenting distinct chemical and physical properties when they are on a micro or macro scale[15]. Nanosized particles known as nanoparticles (NPs) may also be used as fillers[16,17]. However, these materials have a high tendency to aggregate, which may render the dispersion in the polymer unfeasible, compromising the

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anticipated final properties of the material[18]. This dispersion, however, can be improved by changing the surface of the NPs from surfactant molecules or other modifiers by stabilizing and improving the dispersion of the NPs in the matrix, which provides strong repulsion between the NPs[15]. Ordered mesoporous silica are among the materials that have nanometric dimensions that can be added to the polymers. They were first synthesized in 1992 through the M41S[19] family of materials and used to promote the increase of thermal stability of polymers[20]. They have porous ordered structures of silicon oxide (SiO2)[21], involving the silanols (Si-OH) groups[22]. One of the most popular members of this family is the MCM-48[19]. The nanoparticles of MCM-48 (MCM-48 NPs) have a cubic structure indexed in the spatial group Ia3d with 2 interpenetrating networks , biocompatibility and high thermal stability, which justifies their use as a thermal stabilizing agent for polymers[19,20,21,23]. Some authors attribute this effect to the mechanism of radical scavenging by MCM-48 NPs during thermal degradation[20]. This indicates a promising use of MCM-48 NPs to form NCs with PCL to obtain materials with better thermomechanical properties. Lopez-Figueras and co-workers, modified the



Paula, M. V., Azevedo, L. A., Silva, I. D. L., Vinhas, G. M., & Alves Junior, S. surface of MCM-41 with N-[3-trimethoxysilyl) propyl] ethylenediamine (EPTES), to increase the dispersibility of MCM-41 in the PCL[12]. The composites obtained with the modified silica had better thermal properties than the polymer, besides showing greater dispersion of the NPs in the PCL. The improvement in thermal properties and dispersibility was attributed to surface modification with EPTES and the performance of NPs as a barrier. Mallakpour and Khani, using the solvent evaporation method, produced films of NCs between PCL and amorphous SiO2 nanoparticles modified with vitamin B1. The films exhibited better thermomechanical properties than PCL, in addition to exhibiting in vitro bioactivity with the formation of hydroxyapatite, which reveals the potential of the materials obtained to be used in bone tissue engineering[24]. Most of the academic and manufacturing interest in PCL materials involves the use of these materials in medical applications such as drug delivery[25], implants[26], tissue engineering[27] and in food packaging[28]. For these applications, the materials need to be sterilized to kill microorganisms such as fungi, bacteria and viruses[29]. The standard procedure adopted for sterilization is the exposure of these materials to gamma radiation at 25 kGy[30]. This exposure kills microorganisms that can modify their chemical and physical properties, due to the origin of effects such as scission and cross-linking of polymer chains[30]. Thus, this study sought an innovative means to identify the effects of gamma radiation at 25 kGy on the structural, thermal and mechanical properties of NCs formed by PCL with MCM-48 NPs modified with APTES. The choice of dose at 25 kGy was established based on the concept of verification dose maximum (VDmax)[31], to achieve an sterility assurance level (SAL) of 10-6 or better for bacteria. However for virus sterilization doses greater than 25 kGy are required, but the use of these doses is beyond the scope of our investigation[31].The films formed by PCL with 0.5% modified and unmodified MCM-48 NPs were produced by the solvent casting method and exposed to gamma radiation at 25 kGy in the presence of air.

2. Materials and Methods 2.1 Materials All the reagents used were analytical grade and used as received. Tetraethylorthosilicate (98%), cetyltrimethylammonium bromide (98%), a non-ionic detergent (Pluronic® F-127), 3-(Aminopropyl)triethoxysilane and ammonium hydroxide were acquired from Sigma Aldrich. Chloroform, ethanol, toluene and were purchased from Dinâmica.

2.2 Preparation of MCM-48 nanoparticles The MCM-48 NPs were obtained according to the literature[21] with some modifications. To obtain MCM48, 0.5 g of cetyltrimethylammonium bromide, 2.05 g of F127 (Pluronic® F-127) was diluted in 96 mL of distilled water, then 43 mL of ethanol was added, together with 10.05 g of 29% ammonium hydroxide solution. The mixture was stirred until homogeneous, followed by addition of 1.8 grams of tetraethylorthosilicate. This mixture was stirred 2/10

for 10 minutes and kept standing for 12 hours at room temperature. Thereafter, the mixture was centrifuged at 6000 rpm for 30 minutes. A white solid was obtained which was then washed with distilled water and centrifuged again, under the same conditions as above, then dried at 70 °C for 12 hours. This solid was added to the Teflon reactor with 8.5 mL of distilled water, subjected in a sealed system to a temperature of 140 °C for 48 hours, with heating and cooling rates of 10 °C min-1 in a programmable oven. After this step, the solid obtained was heated at 550 °C for 4 hours in a programmable oven, with alterations in heating and cooling at a rate of 10 °C min-1.

2.3 Functionalization of MCM-48 nanoparticles The MCM-48-NH2 NPs were synthesized according to the literature, with some modifications[32]. Basically 3 mL of APTES was added dropwise to a suspension formed by 1 gram of MCM-48 NPs and 30 mL of dry toluene under nitrogen atmosphere. The mixture obtained was refluxed for 24 hours. Then the suspension was filtered and rinsed with ethanol. The resulting solid was dried under vacuum at room temperature and stored for further use.

2.4 Preparation of PCL, PCL/MCM-48 and PCL/MCM48-NH2 NCs films Films of PCL, PCL/MCM-48 and PCL/MCM-48NH2 with 0.5% concentration of silica NPs relative to the polymer mass were produced. Predetermined amounts of MCM-48 and MCM-48-NH2 nanoparticles were added to 5mL of chloroform and sonicated for 30 minutes. A solution of 2.5 g of PCL with 50 ml of chloroform was then added to this and the mixture stirred for 48 hours. After this step, the sample was poured into a petri dish and the chloroform was removed by evaporation in air at room temperature.

2.5 Methods and Analysis 2.5.1 Irradiation of samples All samples were exposed to gamma radiation from a source of 60Cobalt (Gammacell GC220 Excel irradiator - MDS Nordion, Canada) at a dose of 25 kGy (rate of 2.157 kGy h-1, during 11 h, 35 min and 24 s), in the presence of air at room temperature. The dose rate in the center of the gammacell irradiator was determined using a Radcal ionization chamber of 0.3cc volume, previously calibrated in the Ionizing Radiation Metrology Laboratory, located in the Department of Nuclear Energy-Federal University of Pernambuco. 2.5.2 Fourier Transform Infrared Spectroscopy FT-IR spectra was measured with a FT-IR/FT-NIR Perkin Elmer Spectrum 400 Bruker spectrometer in the range of 4000–520 cm-1. Analyses were performed in the attenuated total reflectance mode by direct analysis of samples on ZnSe crystal. 2.5.3 X-ray diffraction Diffractograms were acquired in a Bruker D8 Advance X-ray diffractometer with Cu Kα (0.15 nm), at the speed of 0.02◦ min-1. Polímeros, 31(3), e2021031, 2021

Effects of gamma radiation on nanocomposite films of polycaprolactone with modified MCM-48 2.5.4 Thermogravimetric analysis The thermogravimetric tests were obtained using a SHIMADZU DTG-60H instrument, between room temperature to 600 °C, under nitrogen atmosphere (100 mL.min-1) at a rate of 10 ºC min-1. 2.5.5 Differential Scanning Calorimetry Heat flow curves were carried out in a differential scanning calorimeter, model 1 Star* system (Mettler Toledo) under a nitrogen atmosphere with the following steps: 1) 0 °C to 80 °C, at a rate of 10 °C min-1; 2) cooling to 0 °C, at a rate of 20 °C min-1; and 3) 0 °C to 80 °C, at a rate of 10 °C min-1[13]. The degree of crystallinity Xc of the films was determined based on the equation: X c =

∆H m 0 ∆H m

0 , where ∆H m ,

equals the heat needed for a melting temperature for 100% crystalline PCL. The value used for the heat of fusion of the fully crystalline polymer was 139.3 g-1 [33].

3430 cm-1, ascribed to Si-OH stretching[35], while the peaks at 1052 cm-1 and 800 cm-1 are related to Si-O-Si asymmetric and symmetric stretching vibration modes[35], respectively. In addition, in MCM-48-NH2, a new peak at 1555 cm-1 can be observed, due to N-H stretching, which suggests the surface modification of MCM-48 NPs with APTES[35]. FT-IR spectra of films are presented in Figure 2. PCL exhibited a peak at 1722 cm-1, ascribed to the carbonyl in the PCL, the peaks at 2948 and 2868 cm-1 originated from the C-H bonds of the polymer[36]. NCs films exhibited the same spectra of PCL before and after exposure to gamma radiation. In PCL membranes exposed to gamma radiation at 35 kGy and 65 kGy, the formation of peaks attributed to OH (hydroxyl) and COOH (carboxyl) groups has been observed[30], which is indicative of the radiolytic oxidation of PCL[37]. These findings in general has not been reported for PCL exposed at 25 kGy[29,30]. Mallakpour and Khani

2.5.6 Scanning Electron Microscopy (SEM) The samples were sputtered on carbon tape on an aluminum support and coated with gold, using a Bal‑Tec SCD 050 sputter coater. Images were recorded by a scanning electron microscope (Tescan Mira3) with 10 kV accelerating voltage. 2.5.7 Transmission Electronic Microscopy (TEM) TEM images for the NC films were obtained using a transmission electron microscope (Jeol, model JEM-2100), with 200 kV accelerating voltage. Drops of the NC films suspended in dichloromethane were deposited on copper grids, with slow evaporation of the solvent. 2.5.8 Mechanical measurements Tensile tests were determined on an Instron machine EMIC, DL-500 N, using the ASTM D-882, crosshead speed of 5 mm min-1, at room temperature. For the tensile test, three samples of each film were analyzed. Duncan’s statistical test was used to identify the presence of significant statistical variations.

Figure 1. FT-IR spectra of (a) MCM-48 NPs; (b) MCM-48-NH2 NPs.

3. Results and Discussions 3.1 FT-IR analysis of NCs The modification of MCM-48 NPs with the APTES (to obtain a better distribution in the PCL), forms a siloxane bond[34,35] between the APTES and the surface of MCM48 NPs. After modification with the APTES, the NPs and the PCL chain interaction occurs via hydrogen bonds, established with the amine group present in the NPs with the carbonyl groups of the PCL. The characterization of effects of gamma radiation on the structure, mechanical and thermal properties of all films will be presented in the next sections. FT-IR spectroscopy was used to access the effects of gamma radiation on the chemical structure of films. Figure 1 displays the spectra for MCM-48 NPs and MCM-48-NH2 NPs. MCM-48 exhibited a peak around Polímeros, 31(3), e2021031, 2021

Figure 2. FT-IR spectra of (a) PCL; (b) PCL/MCM-48; (c) PCL/ MCM-48-NH2; (d) PCL-25 kGy; (e) PCL/MCM-48-25 kGy; (f) PCL/MCM-48-NH2-25 kGy. 3/10

Paula, M. V., Azevedo, L. A., Silva, I. D. L., Vinhas, G. M., & Alves Junior, S. obtained films of NCs formed from PCL with amorphous SiO2 nanoparticles and reported changes in peak position of C=O[24]. In view of the foregoing, it is possible to state that the chemical structure of the polymer is preserved after the exposure of PCL and NC films to gamma radiation at 25 kGy.

3.2 Thermogravimetric analysis TGA and its derivative (DTG) curves were acquired to investigate the thermal stability of all samples. Figures 3a, b display the TGA and DTG curves for MCM-48 NPs and MCM-48-NH2 NPs. The MCM-48 NPs exhibited a mass loss around 100 °C, referring to the loss of water molecules adsorbed on the surface of the material. While the MCM-48-NH2 NPs presented additional mass losses between 200 to 600 °C, ascribed to the decomposition of the organic part, demonstrating the surface modification of MCM-48 NPs with APTES[35]. PCL and NCs films had one event of mass loss, between 200 to 500 °C, as observed in the DTG curves, which was ascribed to the thermal decomposition of the polymer[9,36] (Figures 4a, b). Table 1 reports 5% (T5), 10%

(T10), 50% (T50) mass loss temperature for PCL and NCs films and also Tonset and Tmax. Table 1 displays that T5 and T10 of PCL occurred at 361 °C and 370 °C, respectively, while T10 for PCL/MCM-48-NH2 was 10 °C higher than PCL. The PCL/MCM-48 had a similar valor for T10 as compared to polymer. T5 exhibited few variations for NCs as compared to PCL matrices. This improvement in the T10 of the polymer in PCL/MCM-48-NH2, can be ascribed by a better distribution of nanoparticles in the PCL due a surface modification of MCM-48 with APTES, which prevented the aggregation of nanoparticles. This stability can be attributed to the interactions between polymers and nanoparticles through the hydrogen bonds formed with groups of NPs with PCL carbonyl groups. Lopes-Figueras and co-workers observed an increase in thermal stability of composites obtained with PCL and MCM-41 nanoparticles modified with EPTES, but these results were observed only in composites with content above 2% of modified MCM-41[12]. PCL and NCs exposed to gamma radiation had the same thermal degradation profile, with similar values for T5 and T10, as compared to the PCL and NCs films before irradiation

Figure 3. (a) TGA curves of MCM-48 NPs, MCM-48-NH2 NPs; (b) DTG curves of MCM-48 NPs, MCM-48-NH2 NPs.

Figure 4. (a) TGA curves of PCL, PCL/MCM-48, PCL/MCM-48-NH2, PCL-25 kGy, PCL/MCM-48-25 kGy, PCL/MCM-48-NH2-25 kGy; (b) DTG curves of PCL, PCL/MCM-48, PCL/MCM-48-NH2, PCL-25 kGy, PCL/MCM-48-25 kGy, PCL/MCM-48-NH2-25 kGy. 4/10

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Effects of gamma radiation on nanocomposite films of polycaprolactone with modified MCM-48 (Figure 4a). In another study different from our results Lyu and co-workers found 3 mass loss events for composite films formed by PCL with grapefruit seed extract, this thermal behavior profile was attributed to the characteristic of the grapefruit seed extract used to obtain the films[38]. These results indicate that gamma radiation did not change the thermal properties of PCL and its NCs films.

subsequent gamma radiation exposure of the NCs films did not significantly alter the crystalline structure of the PCL in the irradiated films, as observed in the section on X-ray diffraction (Section 3.3) also identified by Horacova and

3.3 X-ray diffraction patterns X-Ray diffraction was used to determine the crystalline profile of NPs and NCs before and after exposure to gamma radiation. Figure 5 shows the diffractograms for NPs. The peak observed for MCM-48 NPs at 2.6° is attributed to plane (211), characteristic of the spatial group Ia3d of the cubic structure of silica[23,35,39]. The diffractogram for MCM-48-NH2 NPs exhibited one reflection angle, indicating that the structure of the silica had been preserved[35] (Figure 5). The intensity of the diffractogram for modified NPs decreased, which may be due to the organic groups on silica after modification with APTES[35]. The PCL has three reflections at 21.5◦, 22◦ and 23.8◦, relative to (110), (111) and (200) planes, respectively, of the orthorhombic crystal structure of PCL[40] (Figure 6). NCs films showed only the semi-crystalline behavior of PCL. The same results were observed for the PCL and NCs films after exposure to gamma radiation (Figure 6), indicating that the gamma radiation at 25 kGy had preserved the crystalline structure of polymer[30].

‘Figure 5. XRD curves of (a) MCM-48 NPs; (b) MCM-48-NH2 NPs.

3.4 Differential Scanning Calorimetry (DSC) DSC analyses were conducted to access the crystallinity of PCL, PCL/MCM-48 and PCL/MCM-48-NH2. Table 2 presents the values for Tm (melting temperature), Tc (crystallization temperature), ΔHm (melt enthalpy) and Xc (crystallinity percentage) for all films analyzed. Addition of MCM48 NPs modified with APTES increased the Tm of the PCL to 65.70 °C, with the same behavior being verified for the Tm of PCL/MCM-48-NH2-25 kGy. The results for Tm obtained by DSC indicate that the addition of NPs and the

Figure 6. XRD curves of (a) PCL; (b) PCL/MCM-48; (c) PCL/MCM-48-NH2; (d) PCL-25 kGy; (e) PCL/MCM-48-25 kGy; (f) PCL/MCM-48-NH2-25 kGy.

Table 1. Thermal properties of PCL, PCL/MCM-48, PCL/MCM-48-NH2 NCs films before and after irradiated at 25 kGy. Sample

T5 (◦C)

T10 (◦C)

T50 (◦C)

Tonset (◦C)

Tmax (◦C)

PCL PCL/MCM-48 PCL/MCM-48-NH2 PCL-25 kGy PCL/MCM-48-25 kGy PCL/MCM-48-NH2-25 kGy

361 364 374 359 357 374

370 372 380 377 374 385

401 388 395 408 394 409

381 385 389 389 389 390

408 397 399 409 405 407

Table 2. Melting temperature, crystallization temperature, enthalpy of melting and percentage of crystallinity for PCL, PCL/MCM-48, and PCL/MCM-48-NH2 NCs films before and after irradiated at 25 kGy. Sample


Tc (◦C)

∆Hm (J/g)

Xc (%)

PCL PCL/MCM-48 PCL/MCM-48-NH2 PCL-25 kGy PCL/MCM-48-25 kGy PCL/MCM-48-NH2-25 kGy

61.33 61.40 65.70 61.91 56.46 65.85

20.46 22.80 23.42 23.46 20.64 23.00

61.14 29.34 49.69 44.82 37.15 51.78

43.89 21.06 35.67 32.18 26.67 37.18

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Paula, M. V., Azevedo, L. A., Silva, I. D. L., Vinhas, G. M., & Alves Junior, S. co-workers[41]. The increase in Tm for NCs with modified NPs may be assigned to the wide dispersion of the NPs in the PCL, which facilitates the generation of thick crystals, which does not occur with unmodified NPs [42]. The crystallization temperature of PCL in the films increased with the addition of NPs. For irradiated samples, marginal variations were observed in Tc. The crystallinity content of PCL (43.89%) is in accordance with the reported[13,43]. For the irradiated PCL and PCL/MCM-48-25 kGy films, however, a reduction in the crystallinity content was observed. These results agree with the findings reported by Foggia and co-workers where they found that gamma radiation did not cause major changes in crystallinity and composition on biomedical poly-(e-caprolactone)/hydroxyapatite composites[44]. These

findings indicate that the exposure of the PCL and NCs films to gamma radiation at 25 kGy did not produce significant changes in Tm, Tc and Xc for the analyzed samples.

3.5 Scanning Electron Microscopy (SEM) The surface morphology of PCL and NCs films was evaluated by SEM analysis. Figure 7a shows the smooth surface of PCL. Figures 7b, c show the surface and cross-section morphology of PCL/MCM-48 and PCL/PCL-48-NH2 NCs films, respectively, where the images revealed that spherical NPs were randomly distributed on the PCL. Figure 8 shows that the surface morphology of nancomposite films irradiated at 25 kgy is preserved.

Figure 7. SEM images of (a) surface of PCL; (b) surface (left) and cross-section of PCL/MCM-48 (right); (c) surface (left) and cross-section (right) of PCL-MCM-48-NH2. 6/10

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Effects of gamma radiation on nanocomposite films of polycaprolactone with modified MCM-48 3.6 Transmission Electron Microscopy (TEM)

3.7 Mechanical properties

Figure 9a, b exhibit the TEM images for the PCL/MCM-48 and PCL/MCM-48-NH2 films, histograms for NPs size distribution and the mean value of the NPs in the films. The images obtained by TEM revealed that MCM-48 NPs and MCM-48-NH2 NPs maintained a spherical shape and random dispersion in the PCL. The mean particle size calculated for MCM-48 NPs and MCM-48-NH2 NPs on the films was 171.2 nm and 148.6 nm, respectively. NPs modified with APTES showed a reduction around 20 nm for the mean particle size. This reduction was attributed to the alteration of the NPs with APTES, which reduced the aggregation of the NPs and favored their interaction with the polymer, reducing the aggregation of NPs in the PCL[24].

The mechanical properties of PCL, PCL/MCM-48 and PCL/MCM-48-NH2 films are shown in Table 3. The parameters evaluated were tensile strength (σ), modulus of elasticity (ε) and elongation at break (Eb). The mean values were compared by the Duncan test, with a significance level of 5%. The Duncan test found that the tensile strength for irradiated and non-irradiated NC films did not have significant statistical differences. Duncan’s test also did not observe statistically significant variations for the modulus of elasticity for all samples analyzed. However, Duncan’s test indicated an increase in Eb for samples PCL/MCM48 and PCL/MCM-48-NH2, which is consistent with the literature[12,24,45]. The improvement in the mean Eb value

Figure 8. SEM images of surface of PCL/MCM-48-25 kGy (left) and ) PCL/MCM-48-NH2-25 kGy (right).

Figure 9. TEM images of (a) PCL/MCM-48; (b) PCL/MCM-48-NH2. Polímeros, 31(3), e2021031, 2021


Paula, M. V., Azevedo, L. A., Silva, I. D. L., Vinhas, G. M., & Alves Junior, S. Table 3. Mechanical properties obtained for PCL, PCL/MCM-48, PCL/MCM-48-NH2 NCs films before and afterirradiation at 25 kGy. Sample PCL PCL/MCM-48 PCL/MCM-48-NH2 PCL-25 kGy PCL/MCM-48-25 kGy PCL/MCM-48-NH2-25 kGy

σ (MPa) 10.18 ± 0.62a 9.98 ± 1.15a 9.28 ± 0.76a 10.34 ± 1.17 a 9.38 ± 0.97a 9.21 ± 1.43a

ε (MPa) 169.86 ± 7.01a 174.36 ± 9.58ª 170.30 ± 9.26a 167.70 ± 9.98a 164.16 ± 4.06a 182.2 ± 17.02a

Eb(%) 76.29 ± 5.79a 376,70 ± 9.04b 137,40 ± 15.51c 44,55 ± 0.98d 123,63 ± 9.51c 50,72 ± 7,71d

The mean with the same letter (i.e. a) in the same column did not differ at p <0.05 for the Duncan’s test.

of these samples is attributed to the formation of hydrogen bonds between the MCM-48 and the PCL chains, favoring the interfacial adhesion between the components of the nanocomposites[24]. The reduction of Eb in the irradiated samples is attributed to the effects of gamma radiation on the films. These results indicate that gamma radiation at 25 kGy did not modify the mechanical properties of tensile strength and modulus of elasticity of PCL and NCs films. This indicates that the use of gamma radiation at 25 kGy is a suitable method for the sterilization of the nanocomposite films formed by PCL with MCM-48 NPs modified with APTES.

4. Conclusions PCL nanocomposites films with 0.5% of MCM-48 and MCM-48-NH2 NPs, were satisfactorily obtained by the solvent evaporation method and exposed to gamma radiation at 25 kGy. The images obtained by MEV and TEM revealed the presence of aggregates of NPs randomly dispersed in the films. TEM images exhibited that the MCM-48-NH2 NPs presented smaller particle size in the PCL, which was attributed to modification with the APTES. Samples irradiated at 25 kGy had the same spectral behavior in the infrared region as the non-irradiated samples, which is indicative of the maintenance of the chemical structure of the PCL. The same results were observed in the diffractograms of the irradiated films, indicating the maintenance of the semicrystalline character of PCL. The PCL MCM-48-NH2 film exhibited an enhancement to the thermal stability of PCL, due to modification with the APTES. The exposure of the PCL/MCM-48 and PCL/MCM-48-NH2 films to gamma radiation at 25 kGy did not cause major changes in the mechanical and thermal properties of NC films, demonstrating success of the sterilization method.

5. Acknowledgements The authors thank Fundação de Amparo Ciência e Tecnologia do Estado de Pernambuco (FACEPE) for the provided scholarships.

6. References 1. Labet, M., & Thielemans, W. (2009). Synthesis of polycaprolactone: a review. Chemical Society Reviews, 38(12), 3484-3504. http:// dx.doi.org/10.1039/b820162p. PMid:20449064. 2. Woodruff, M. A., & Hutmacher, D. W. (2010). The return of a forgotten polymer Polycaprolactone in the 21st century. 8/10

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Paula, M. V., Azevedo, L. A., Silva, I. D. L., Vinhas, G. M., & Alves Junior, S. 39. Solovyov, L. A., Belousov, O. V., Dinnebier, R. E., Shmakov, A. N., & Kirik, S. D. (2007). X-ray diffraction structure analysis of MCM-48 mesoporous silica. The Journal of Physical Chemistry B, 109(8), 3233-3237. http://dx.doi.org/10.1021/ jp068521h. PMid:16851346. 40. Augustine, R., Malik, H. N., Singhal, D. K., Mukherjee, A., Malakar, D., Kalarikkal, N., & Thomas, S. (2014). Electrospun polycaprolactone/ZnO nanocomposite membranes as biomaterials with antibacterial and cell adhesion properties. Journal of Polymer Research, 21(3), 347. http://dx.doi.org/10.1007/ s10965-013-0347-6. 41. Horakova, J., Klicova, M., Erben, J., Klapstova, A., Novotny, V., Behalek, L., & Chvojka, J. (2020). Impact of Various Sterilization and Disinfection Techniques on Electrospun Poly-ε-caprolactone. ACS Omega, 5(15), 8885-8892. http:// dx.doi.org/10.1021/acsomega.0c00503. PMid:32337451. 42. Augustine, R., Kalarikkal, N., & Thomas, S. (2016). Effect of zinc oxide nanoparticles on the in vitro degradation of electrospun polycaprolactone membranes in simulated body fluid. International Journal of Polymeric Materials and


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Polímeros, 31(3), e2021031, 2021

ISSN 1678-5169 (Online)


Effect of drying different inclusion plasters on the mechanical properties of thermoactivated acrylic resins Tarcisio José de Arruda Paes-Junior1 , Natália Rivoli Rossi1 , Tayná Mendes Inácio de Carvalho1 , Vanessa Cruz Macedo1 , Michelle de Sá dos Santos Gomes1 , Leonardo Jiro Nomura Nakano1  and Cristiane Mayumi Inagati1* Departamento de Materiais Odontológicos e Prótese, Instituto de Ciência e Tecnologia, Universidade Estadual Paulista “Júlio de Mesquita Filho”– UNESP, São José dos Campos, SP, Brasil



Abstract This article aimed to evaluate some mechanical and chemical properties of acrylic thermoactivated resins by microwave energy, varying the condition and type of plaster. The groups were divided into Lucitone and Vipi-Wave groups, with or without previous treatment (drying) of type III and IV plasters. It was evaluated flexural strength, microhardness, roughness, porosity, residual monomer, and also, time and temperature relationship of plaster and acrylic resin during the polymerization cycles. The data were analyzed using Analysis of Variance (ANOVA) 5%, followed by Tukey’s test. The results showed that the drying of the plaster influenced the results and the groups with dry plaster maintained a higher temperature permanence. Therefore, changes in the water condition contained in the inclusion plaster showed effects on the final properties of the acrylic resin, which may be a technical indicator for laboratory procedures in the manufacture of prosthetic devices. Keywords: acrylic resin, dental plaster, flexural strength, hardness, roughness. How to cite: Paes-Junior, T. J. A., Rossi, N. R., Carvalho, T. M. I., Macedo, V. C., Gomes, M. S. S., Nakano, L. J. N., & Inagati, C. M. (2021). Effect of drying different inclusion plasters on the mechanical properties of thermoactivated acrylic resins. Polímeros: Ciência e Tecnologia, 31(3), e2021032. https://doi.org/10.1590/0104-1428.20210049

1. Introduction Polymethylmethacrylate (PMMA) or better known as acrylic resin is a polymer widely used in Dentistry. Although there are several clinical applications of acrylic resin, from temporary crowns to total prosthesis, PMMA has some limitations, such as low mechanical resistance[1] and dimensional change[2]. To improve these properties, several studies have been studying materials that can reinforce acrylic resin, as well as nylon, fiberglass, nanoparticles, among others[1,3], as well as ways to improve the processing technique of acrylic resin[4]. The mechanical properties can be influenced by the type of polymerization, for instance, by the water bath or by the autoclave[5, 6]. However, according to some studies, the best technique used is microwave polymerization, which produces greater mechanical resistance and less dimensional change when compared to the heat technique, but unlikely, a lower surface hardness[4,7]. The microwave is an alternative to the conventional method being easier to handle and clean because it warms up faster. Also, the activation is homogeneous, leading to a better finish of the prostheses[4,8]. The dimensional change occurs during the entire process, including as a result of choosing the type of plaster used during obtaining the work[2]. Another concern is the amount of residual monomer from the polymerization which irritates

Polímeros, 31(3), e2021032, 2021

the patient’s mucosa[9]. However, according to the work of Paes-Junior et al.[10] it was possible to reduce the amount of residues by pre-drying the plaster before curing the resin. During the separation of the plaster, depending on the agent used, a smoother and brighter surface is obtained[11]. Besides, to reduce the undesirable effects of PMMA, a light- and heat-cured urethane dimethacrylate (UDMA) emerged, which is characterized by its biocompatibility, low bacterial adherence, greater resistance to masticatory loads, and thus providing greater satisfaction for the patient[12]. Finally, the objective of the present study was to evaluate the mechanical properties of the acrylic resin by drying two types of plasters (III and IV) which were used in the inclusion phase, and also the activation method, by microwave oven or water bath. The null hypothesis was that the type and performance of dried plaster and the type of thermal activation did not influence the mechanical and chemical properties of the acrylic resin.

2. Materials and Methods The materials used in the study were the following acrylic resins, Lucitone 550 (Dentsply Sirona, United States) and Vipi-Wave (Dental Vipi, Brazil) and also, the dental plasters, Gypsum stone type III Herodont (Coltene, Brazil)



Paes-Junior, T. J. A., Rossi, N. R., Carvalho, T. M. I., Macedo, V. C., Gomes, M. S. S., Nakano, L. J. N. & Inagati, C. M. and Gypsum stone type IV Velmix (Kavo Kerr, Brazil). The groups of the present study can be seen in Table 1. Three types of metallic patterns of stainless steel with geometric shape of a parallelogram and dimensions respectively 2.0X2.0X2.0cm, 2.0X2.0X1.0cm, and 2.0X2.0X0, 5cm were used to perform the specimens. These patterns were copied by a laboratory silicone-based impression material (Rodorsil-VWL, Brazil), and from these replicas were made in wax 7 (Classic, Brazil). The wax replicas were included in the muffle varying the type of the plaster and its manipulation. The type III plaster was mixed in the proportion of 180g of powder to 60ml of water. For the type IV plaster, the proportion was 180g of powder to 36 ml of water. Six wax patterns were positioned equidistantly and parallel to each other. The muffle was closed to leak the plaster into the counter muffle. The counter muffle was filled with plasters type III and IV provided in 240g of powder for 80ml of water and 240g of plaster for 50ml of water respectively, through the opening in the upper part of the muffle. For the groups where the plaster was previously dried, the procedure adopted was based on the methodology described by Canay et al.[13] and adapted by Paes-Junior et al.[10] who dried the inclusion plaster in a microwave oven for 10 min at 600W. Then, the muffles were left at room temperature for a period of one hour until completely cooled, and then with their parts open, they were stored for a minimum of 24 hours in a dry oven under temperature of 37ºC, before pressing and polymerizing acrylic resins. After one hour, the muffle parts were opened and the wax patterns were removed under immersion in heated water. Then, a thin layer of insulator for acrylic resin Al-Cote (Dentsply, Brazil) was applied with a brush over the entire surface area of the ​​ plaster. For the pressing procedure, Lucitone 550 or Vipi-Wave acrylic resin was used, being accommodated during the plastic phase in the spaces left by the wax. Then, the muffle was closed and placed in a hydraulic press (Techno Máquinas, Brazil), until the final pressure of 1000Kgf maintained for 30 minutes before polymerization. For the microwave polymerization cycle was used a Continental AW-30 oven (BS Continental da Amazônia Ind. e Com. Ltda, Brazil) with a rotating plate and frequency of 2450 megahertz (MHz) with maximum power of 900W. The utilized cycle was 20% of the power of the device for 20 minutes plus 5min to 60% of the power. Then, the muffle was kept for about two hours at room temperature until cooled. To finish the pieces, a rotary sander (Panambra, Brazil) was used under constant refrigeration, and the wear

was performed by the aluminum oxide sandpaper Norton (Norton), with decreasing weight (180, 320, 600#).

2.1. Fabrication of specimens for flexural strength and FTIR Metallic patterns made in bars shape, rectangular, with sharp edges, in the dimensions of 27 x 12.60 x 3.10 mm were included in the muffle as described before. The only difference was the use of the Zetalabor silicone (Zermack, Brazil) which was applied around the metallic patterns. After 30 minutes, the counter muffle was settled and the plaster was poured. One hour after the final inclusion phase, the muffle was opened and the metallic patterns were removed. A thin layer of insulator was applied over the silicone and the acrylic resin was included in the muffle. After its polymerization, the specimens were removed from the muffle and placed in a water container, then the finishing was performed as already described. The FT-IR was performed in a PerkinElmer Spectrum One spectrophotometer, using the attenuated total reflectance (ATR) technique, in the 5002000 cm-1 region, with a resolution of 4 cm-1 and 16 scans, in which are obtained the quantified MR performing the respective scans per specimen. The spectra obtained were sent to a software for graph analysis (Origin 7.0 with PeakFitting), where they were imported, analyzed, plotted and the comparative graphs established[14,15].

2.2. Microhardness analysis Before the test, each group was stored in distilled water at a temperature of 37ºC+2ºC, for a period of 48h+2h. After this period, the microhardness analysis was performed by Vickers indentation test (VHN), using a microdurometer (FM 700, Japan), with a load of 10gf for 5 seconds, with 3 indentations in each specimen.

2.3. Porosity analysis The results were obtained through the visual inspection by three previously calibrated observers who made a ranking and assigned scores based on the following criteria: 0 (without porosity), 1 (minimum amount of porosity), 2 (average amount of porosity), and 3 (large amount of porosity). After the evaluation, the arithmetic means of the scores attributed individually by the examiners were obtained.

2.4. Roughness analysis The roughness analysis was performed over one of the faces of the specimen. However, before the analysis, the samples were ultrasonically cleaned with distilled water

Table 1. Groups analyzed according to the treatment of choice. Groups LP LPS VP VPS L4 L4S V4 V4S


Treatment Lucitone 550 acrylic resin, microwave polymerization without prior treatment of type III plaster. Lucitone 550 acrylic resin, microwave polymerization with pre-treatment of type III plaster (dissected plaster). Vipi-Wave acrylic resin, microwave polymerization without prior treatment of type III plaster. Vipi-Wave acrylic resin, microwave polymerization with pre-treatment of type III plaster (dissected plaster). Lucitone 550 acrylic resin, microwave polymerization without previous treatment of type IV plaster. Lucitone 550 acrylic resin, microwave polymerization with previous treatment of type IV plaster (dissected plaster). Vipi-Wave acrylic resin, microwave polymerization without prior treatment of type IV plaster. Vipi-Wave acrylic resin, microwave polymerization with pre-treatment of type IV plaster (dissected plaster).

Polímeros, 31(3), e2021032, 2021

Effect of drying different inclusion plasters on the mechanical properties of thermoactivated acrylic resins for 380s, and dried with an air jet, free of water and oil. For the roughness test, the specimens were analyzed in a contact roughness (Mitutoyo SJ – 400), with 3 readings of 4 mm, in two directions perpendicular to the sample surface. Roughness analysis was performed on one of the faces of the cubes used in the porosity analysis on one of the faces.

2.8 Statistical analysis The obtained values were subjected to the normality and homogeneity test. For flexural strength, microhardness, and surface roughness, it was performed the ANOVA variance analysis with the Turkey test with a significance level of 5%. The porosity analysis did not follow a normal distribution so the non-parametric Kruskal-Wallis test was performed.

2.5. Flexural strength analysis The samples were placed in distilled water and kept at a temperature of 37 ºC + 2ºC, for a period of 48h + 2h. After this period, the three-point flexural test was performed by a universal testing machine – EMIC (Model DL-1000, Brazil), with a load cell of 100Kgf and an application speed of 5 mm/min. The flexural strength values for the groups were obtained in mega Pascals (MPa) and subjected to statistical analysis.

3. Results and Discussions 3.1. Porosity analysis Regardless of the plaster, the type of resin, and drying (volume of 2 cm3), the amount of porosity was practically non-existent for all groups. For the drying (volume of 4 cm3), it was found slightly ​​ higher values for the LC and L4S groups. While for the drying (volume of 8 cm3), the LC, LCS, and VCS groups had higher porosity than the other groups (Table 2). The groups that used Vipi-Wave resin had lower porosity values, being that the VP group showed absence of porosities regardless of the specimen volume. The VC and VPS groups showed minimal bubbles for the 8cm3 volume and absence for the other volumes. This group was in the opposite direction to the others, where plaster drying, especially for specimens of greater volume, increased the occurrence of bubbles, which may perhaps be explained by the characteristics of the type III alpha hemidrate plaster used in this group, however there was no statistical difference between drying and the type of plaster.

2.6. Measurement of temperature in the polymerization cycle A muffle from each group was subjected to the analysis of the temperature gradient. For this purpose, a device in the form of a metallic tablet was inserted into one of the thickest specimens (8 cm3), surrounded by a plastic for its protection. According to the study by Savirmath and Mishra[16], dimensional change occurs with temperature change. These were positioned during the resin inclusion phase. The temperature monitoring was done 20 times, minute by minute of phase 1 of the polymerization cycle and, five times of phase 2 of the cycle. An iButton DS1922E temperature sensor (Maxim integrated, USA) was used to perform the temperature measurement, which is a resistant and self-sufficient system that measures temperatures between 15ºC and 140ºC. This sensor is configured to communicate with a computer using the 1-Wire serial protocol, which connects to a USB port. The sensor is connected to this device which can program and activate according to the registration data, then the results are transferred to a computer in graphs and tables. This analysis was qualitative and determined a time/temperature graph for each group.

3.2. Microhardness analysis Despite dried processo, the groups that used type IV plaster showed lower values ​​of microhardness compared to type III plaster. Regarding resin, Lucitone presented higher microhardness values ​​(Table 3).

3.3. Roughness analysis Type III plaster showed the lowest roughness, and when the plaster was previously dried, it also found small roughness values (Table 4), with the Lucitone resin showing lower values ​​(Table 5).

2.7. Remaining residual monomer (RM) analysis The acrylic resin samples used in the flexural test were used. The specimens were submitted to finishing with 600 grit sandpaper and, five of each group had the MR quantified, through FT-Raman spectroscopy, with the respective scans per specimen. The obtained spectra were sent to a software (Origin 7.0 with Peak Fitting) which graphs were imported, analyzed, and established.

3.4. Flexural strength analysis Lucitone 550 resin presented greater flexural strength than Vipi- Wave resin and the drying of the plaster also influenced the increase of its resistance (Table 6).

Table 2. Kruskal-Wallis for the groups and volumes observed. Volume 2mm 4mm 8mm

LC 1 2 3

LCS 1 1 3

LP 0 0 2

LPS 0 1 2

VC 0 0 1

VCS 0 1 3

VP 0 0 0

VPS 0 0 1

L4 0 1 3

L4S 1 2 2

V4 0 1 2

V4S 0 0 2

LC: Lucitone, without prior treatment of type II plaster. LCS: Lucitone, with pre-treatment of type II plaster. LP: Lucitone, without prior treatment of type III plaster. LPS: Lucitone, with pre-treatment of type III plaster. VC: Vipi-Wave, without prior treatment of type II plaster. VCS: Vipi-Wave, with pre-treatment of type II plaster. VP: Vipi-Wave, without prior treatment of type III plaster. VPS: Vipi-Wave, with pre-treatment of type III plaster. L4: Lucitone, without prior treatment of type IV plaster. L4S: Lucitone, with pre-treatment of type IV plaster. V4: Vipi-Wave, without prior treatment of type IV plaster. V4S: Vipi-Wave, with pre-treatment of type IV plaster.

Polímeros, 31(3), e2021032, 2021


Paes-Junior, T. J. A., Rossi, N. R., Carvalho, T. M. I., Macedo, V. C., Gomes, M. S. S., Nakano, L. J. N. & Inagati, C. M. Table 3. Vickers microhardness averages, considering only the type of plaster. Plaster Type IV Type III

Hardness 24.82 ± 2.93 27.49 ± 4.11

Resin Vipiwave Lucitone


Hardness 25.85 ± 3.14 27.74 ± 4.09


Table 4. Roughness averages (Ra), considering only the type of plaster. Plaster Type III Type IV

Roughness (Ra) 0.14 ± 0.03 0.16 ± 0.03

Drying Yes No


Roughness (Ra) 0.15 ± 0.03 0.16 ± 0.04


Table 5. Average roughness of the type of resin and whether or not drying was performed. Drying Yes Yes No No

Resin Lucitone Vipiwave Vipiwave Lucitone

0.14 ± 0.03 0.16 ± 0.02 0.16 ± 0.04 0.16 ± 0.04


Table 6. Flexural strength averages in MPa of the types of resin. Resin Vipi-Wave Lucitone

Strength 102.80 ± 15.33 121.16 ± 31.06


Drying No Yes

Strength 107.55 ± 17.50 116.41 ± 32.02


3.5. FTIR spectroscopy analysis The FTIR curves of the Vipi- Wave and Lucitone resins can be interpreted in Figure 1 and 2, respectively, where it was observed that all the graphics obtained, regardless of the group, were coherent and were in the same observation ranges. In Figure 1, the curve corresponding to the band at 1150-1cm was related to vibration and it was possible to infer that the VP group had a lower conversion when compared to the other groups. While in Figure 2 in the 750-1cm band, the L4 and L4S groups had lower values.

3.6. Temperature gradient Table 7 shows that for all groups there was a fast temperature rise of the resin when activated in the microwave, maintaining the temperature at the maximum peak, and decreasing quickly. When the plaster was dried, all the groups had a significant increase in the maintenance temperature, except for the LC group.

Figure 1. FTIR curve for the Vipi-Wave group.

3.7. Discussion The type of plaster, plaster drying method, and type of activation influenced the mechanical and chemical properties of the acrylic resin so the null hypothesis was rejected. Type III and IV plaster differ from each other in terms of the amount of water needed for their handling, so their mechanical strength was also different. In the present work, the plaster drying method was effective as it promoted a better removal of the water, without interfering with the subsequent inclusion and polymerization processes, in agreement with the works of Canay et al.[13]. This study evaluated the effect of the drying method of a type III plaster in a microwave and a conventional activation, which obtained an increase in resistance when the drying method was performed, however, there was no difference between the activation method of choice. In contrast, Paes4/6

Figure 2. FTIR curve for the Lucitone group.

Junior et al.[10] analyzed whether the performance of drying plaster previously would influence the amount of residual monomer and the study found that the specimens where the Polímeros, 31(3), e2021032, 2021

Effect of drying different inclusion plasters on the mechanical properties of thermoactivated acrylic resins Table 7. Maximum temperature data (ºC) for the cycles and time of permanence for the experimental groups. Temperature (ºC) Temperatura peak Permanence time

LC 123°C 30s

LCS 126°C 30s

LP 126°C 150s

LPS 126°C 300s

VC 120°C 30s

VCS 126°C 300s

VP 126°C 90s

VPS 126°C 360s

LC: Lucitone, without prior treatment of type II plaster. LCS: Lucitone, with pre-treatment of type II plaster. LP: Lucitone, without prior treatment of type III plaster. LPS: Lucitone, with pre-treatment of type III plaster. VC: Vipi-Wave, without prior treatment of type II plaster. VCS: Vipi-Wave, with pre-treatment of type II plaster. VP: Vipi-Wave, without prior treatment of type III plaster. VPS: Vipi-Wave, with pre-treatment of type III plaster.

drying method was performed obtained a lower amount of residual monomer. The samples that had a thickness of 8 cm3 showed a significant difference in porosity when compared to the volumes of 2 cm3 and 4 cm3, regardless of the plaster choice and drying. This probably occurred, because heat dissipation is harder on thick specimens thus causing internal porosities, also found in the work of Kimpara et al. [17] who evaluated four different polymerization cycles and found that performing short cycles has a greater amount of porosity. While Neisser et al.[18] also studied four different polymerization cycles and observed that the shorter cycles with lower energy, raised the temperature, thus causing a higher index of porosity. The Ibutton wireless device showed a fast increase of temperature, with a constant temperature peak of the specimens for all groups, especially in the groups where plaster was dried due to the slow heat dissipation. Regarding the FTIR test to evaluate the degree of polymerization, a qualitative aspect was considered where all the obtained graphs, regardless of the group, had the same observation ranges. In the analysis of roughness, the LPS group showed a higher level of roughness, regardless of the volume of the material used. In the study performed by Rizzatti-Barbosa and Ribeiro-Dasilva[19], it was compared the roughness of the simultaneous polymerization or double vial of the upper and lower total dentures that were made using the microwave or conventional technique. It was observed that there were no differences between the groups studied. For the polymerization of acrylic resin, it is necessary to activate free radicals, this can occur through microwave waves, heat by hot water, or even by photoactivation. In the literature, the use of the microwave technique for polymerization reduced the amount of residual monomer present in total dentures[20]. In the present study, a significant change was observed in the flexural strength for the drying method performed in the plaster, when compared to the conventional cycle in a heated water bath, in agreement with the research by Silva et al.[7]. This study evaluated the flexural strength of four heat cycles by microwave energy and the authors concluded that when the microwave was used, the flexural strength was greater when compared to the other groups.

4. Conclusions Within the limitations of this study, it was possible to conclude that the drying process of the plaster, regardless of the type of plaster used, could influence the mechanical properties of the thermoactivated resins presented in this study. Polímeros, 31(3), e2021032, 2021

5. Acknowledgements The authors would like to thank FAPESP Process nº 2014/08408-1 for their assistance in this research.

6. References 1. Al-Harbi, F. A., Abdel-Halim, M. S., Gad, M. M., Fouda, S. M., Baba, N. Z., AlRumaih, H. S., & Akhtar, S. (2019). Effect of nanodiamond addition on flexural strength, impact strength, and surface roughness of PMMA denture base. Journal of Prosthodontics, 28(1), e417-e425. http://dx.doi.org/10.1111/ jopr.12969. PMid:30353608. 2. Baydas, S., Bayindir, F., & Akyil, M. S. (2003). Effect of processing variables (Different Compression Packing Processes and Investment Material Types) and time on the dimensional accuracy of polymethyl methacrylate denture bases. Dental Materials Journal, 22(2), 206-213. http://dx.doi.org/10.4012/ dmj.22.206. PMid:12873123. 3. Karci, M., Demir, N., & Yazman, S. (2019). Evaluation of Flexural Strength of Different Denture Base Materials Reinforced with Different Nanoparticles. Journal of Prosthodontics, 28(5), 572-579. http://dx.doi.org/10.1111/jopr.12974. PMid:30298558. 4. Durkan, R., & Oyar, P. (2018). Comparison of mechanical and dynamic mechanical behaviors of different dental resins polymerized by different polymerization techniques. Nigerian Journal of Clinical Practice, 21(9), 1144-1149. http://dx.doi. org/10.4103/njcp.njcp_423_17. PMid:30156199. 5. Gad, M. M., Fouda, S. M., ArRejaie, A. S., & Al-Thobity, A. M. (2019). Comparative Effect of Different Polymerization Techniques on the Flexural and Surface Properties of Acrylic Denture Bases. Journal of Prosthodontics, 28(4), 458-465. http://dx.doi.org/10.1111/jopr.12605. PMid:28543925. 6. Bural, C., Aktaş, E., Deniz, G., Ünlüçerçi, Y., Kızılcan, N., & Bayraktar, G. (2011). Effect of post-polymerization heattreatments on degree of conversion, leaching residual MMA and in vitro cytotoxicity of autopolymerizing acrylic repair resin. Dental Materials, 27(11), 1135-1143. http://dx.doi. org/10.1016/j.dental.2011.08.007. PMid:21920593. 7. Silva, L. H., Tango, R. N., Kimpara, E. T., Saavedra, G. S. F. A., & Paes-Junior, T. J. A. (2011). Flexural strength and microhardness of a chemically activated acrylic resin after microwave energy treatment. Revista Gaucha de Odontologia, 59(2), 237-242. 8. Figuerôa, R. M. S., Conterno, B., Arrais, C. A. G., Sugio, C. Y. C., Urban, V. M., & Neppelenbroek, K. H. (2018). Porosity, water sorption and solubility of denture base acrylic resins polymerized conventionally or in microwave. Journal of Applied Oral Science, 26(0), e20170383. http://dx.doi. org/10.1590/1678-7757-2017-0383. PMid:29742260. 9. Melilli, D., Curró, G., Perna, A. M., & Cassaro, A. (2009). Cytotoxicity of four types of resins used for removable denture bases: in vitro comparative analysis. Minerva Stomatologica, 58(9), 425-434. PMid:19893467. 10. Paes-Junior, T. J. A., Carvalho, R. F., Cavalcanti, S. C. M., Saavedra, G. S. F. A., & Borges, A. L. S. (2013). Influence 5/6

Paes-Junior, T. J. A., Rossi, N. R., Carvalho, T. M. I., Macedo, V. C., Gomes, M. S. S., Nakano, L. J. N. & Inagati, C. M. of plaster drying on the amount of residual monomer in heatcured acrylic resins. Brazilian Journal of Oral Sciences, 12(2), 84-89. http://dx.doi.org/10.1590/S1677-32252013000200003. 11. Kim, T. H., Ahn, T. J., Enciso, R., & Knezevic, A. (2014). Effect of gypsum separating media on the appearance of stone cast surfaces. The Journal of Prosthetic Dentistry, 112(4), 1001-1005. http://dx.doi.org/10.1016/j.prosdent.2014.06.011. PMid:25134996. 12. Ali, I. L., Yunus, N., & Abu-Hassan, M. I. (2008). Hardness, flexural strength, and flexural modulus comparisons of three differently cured denture base systems. Journal of Prosthodontics, 17(7), 545-549. http://dx.doi.org/10.1111/j.1532849X.2008.00357.x. PMid:18761582. 13. Canay, S., Hersek, N., Çiftçi, Y., & Akça, K. (1999). Comparision of diametral tensile strength of microwave and oven-dried investment materials. The Journal of Prosthetic Dentistry, 82(3), 286-290. http://dx.doi.org/10.1016/S0022-3913(99)70082-X. PMid:10479254. 14. Rodriguez, L. S., Paleari, A. G., Giro, G., Oliveira, N. M., Jr., Pero, A. C., & Compagnoni, M. A. (2013). Chemical characterization and flexural strength of a denture base acrylic resin with Monomer 2-Tert-Butylaminoethyl Methacrylate. Journal of Prosthodontics, 22(4), 292-297. http://dx.doi. org/10.1111/j.1532-849X.2012.00942.x. PMid:23106690. 15. Türkcan, I., Nalbant, A. D., Bat, E., & Akca, G. (2018). Examination of 2-methacryloyloxyethyl phosphorylcholine polymer coated acrylic resin denture base material: surface characteristics and candida albicans adhesion. Journal of Materials Science. Materials in Medicine, 29(7), 107. http:// dx.doi.org/10.1007/s10856-018-6116-7. PMid:29971499.


16. Savirmath, A., & Mishra, V. (2016). A comparative evaluation of the linear dimensional changes of two different commercially available heat cure acrylic resins during three different cooling regimens. Journal of Clinical and Diagnostic Research : JCDR, 10(11), 50-54. http://dx.doi.org/10.7860/JCDR/2016/22066.8903. PMid:28050504. 17. Kimpara, E. T., Silva, L. H., Costa, C. B., Borges, A. L. S., Tango, R. N., & Paes-Junior, T. J. A. (2009). Acrylic resin for complete denture: effect of polymerization cycles at residual monomer released and porosity evidence. Revista da Faculdade de Odontologia: UPF, 14(1), 37-41. 18. Neisser, M. P., Hilgert, E., Cavalcanti, B. N., Barros, E. A., & Magalhães, O., No. (2005). Thermal curves of acrylic resins in microwave curing. Brazilian Dental Science, 8(2), 25-30. http://dx.doi.org/10.14295/bds.2005.v8i2.385. 19. Rizzatti-Barbosa, C. M., & Ribeiro-Dasilva, M. C. (2009). Influence of double flask investing and microwave heating on the superficial porosity, surface roughness, and knoop hardness of acrylic resin. Journal of Prosthodontics, 18(6), 503-506. http://dx.doi.org/10.1111/j.1532-849X.2009.00469.x. PMid:19432756. 20. Urban, V. M., Machado, A. L., Oliveira, R. V., Vergani, C. E., Pavarina, A. C., & Cass, Q. B. (2007). Residual monomer of reline acrylic resins: effect of water-bath and microwave postpolymerization treatments. Dental Materials, 23(3), 363-368. http://dx.doi.org/10.1016/j.dental.2006.01.021. PMid:16620950. Received: June 21, 2021 Revised: Oct. 18, 2021 Accepted: Nov. 05, 2021

Polímeros, 31(3), e2021032, 2021

ISSN 1678-5169 (Online)


ABS/Recycled PCTG blend compatibilized with SBS: effect on mechanical properties and morphology Juliana Augusto Molari1* , Deborah Dibbern Brunelli1  Instituto Tecnológico de Aeronáutica – ITA, São José dos Campos, SP, Brasil



Abstract The reuse of plastic polymers is one of the ways to reduce the negative environmental impact caused by these products. This work presents a study of mechanical and morphological properties of ABS and PCTG residue blend using SBS as compatibilizing agent to make copolyester recycling process feasible. It was observed that the incorporation of SBS in the mixture decreased the stiffness and increased the impact resistance compared to the results obtained in the noncompatible mixture, indicating that the SBS acted as a toughening agent in the mixture. Additionally, according to the results obtained by DSC and SEM, the blends obtained can be considered partially miscible, since two glass transition temperatures were evidenced shifted by a few degrees from neat components. Micrograph suggests that there are SBS small domain inclusions dispersed in the PCTG matrix and partial compatibility occurred by partial interaction of the SBS in the interface. Keywords: ABS, PCTG, SBS, polymeric blends, compatibilization. How to cite: Molari, J. A., & Brunelli, D. D. (2021). ABS/Recycled PCTG blend compatibilized with SBS: effect on mechanical properties and morphology. Polímeros: Ciência e Tecnologia, 31(3), e2021033. https://doi.org/10.1590/01041428.20210074

1. Introduction The manufacture of plastic products has undergone a great expansion recently due to global demand [1]. In general, the increase in production contributes to increased waste. Advanced materials with specific and sustainable properties have been extensively developed through the manufacturing of engineering polymer commodities and/ or blends. The great advantage of this type of mixture is the possibility of reusing residues whose original application does not allow the return of the reprocessed material [2]. Evaluating the feasibility of obtaining blends of two polymers for industrial applications involves studying their miscibility and mechanical properties. However, if the polymers are immiscible, the mixture between them will be brittle and morphologically unstable. Thus, the use of compatibilizers that stabilize the phases and promote synergy between them is recommended. Acrylonitrile-butadiene-styrene (ABS) copolymer is widely used in the automotive, aeronautics, home appliances and packaging industries, among others, precisely because of its unique characteristics of mechanical resistance and brightness [3]. Poly (1,4-cyclohexylene di-methylene terephthalate glycol) (PCTG) is a copolymer formed by the esterification and polycondensation reactions (with metallic catalyst) of cyclohexane di-methanol (CHDM), terephthalic acid (TPA) and ethylene glycol (EG). PCTG has excellent processing and optical properties that allow the use of this polymer to

Polímeros, 31(3), e2021033, 2021

obtain transparent thermoformed products with complex shapes, being widely used in the packaging industry [4]. Once copolyesters such as PCTG can undergo hydrolytic thermal oxidation when reprocessed, which can cause discoloration or yellowing, chain splitting, and molar mass reduction, it is difficult to recycle or reuse them [4]. The proposed alternative to use PCTG residues is to blend it with ABS, using block copolymer styrene-butadienestyrene (SBS) as a compatibilizing agent. As ABS is opaque, it contributes to minimize visual effects of the yellowing process that occurs to PCTG in the proposed blend. This work aimed to develop a new route to reuse industrial PCTG waste. For this, blending PCTG residues with virgin ABS was evaluated, using block copolymer butadienestyrene (SBS) as a compatibilizing agent. The influence on the mechanical properties were evaluated by tensile and Charpy impact tests and morphology of the blend was evaluated by scanning electron microscope (SEM) and DSC.

2. Experimental 2.1. Materials The materials used included ABS polymer, Terluran® GP22, compatibilizing agent, SBS Styroflex® 2G66, both produced by Styrolution ™ and PCTG waste obtained from grinded injection molding channels (SKYGREEN® JN400, from SK Chemicals).



Morali, J. A. & Brunelli, D. D. 2.2. Preparation methods

2.3.3. Scanning Electron Microscopy (SEM)

The blends were prepared according to the proportions shown in the Table 1, varying the concentration of the blend components and compatibilizing agent. For the compatibilized samples, different levels of SBS were added and the amount sum 100 %. Polymers were first weighted, cold mixed and then extruded using a twin screw extruder to ensure adequate homogeneity of the blend. Prior to the injection of the test specimens, all extruded mixtures were maintained in the dehumidifier at 80 °C for 4 hours.

Scanning Electron Microscopy was used to evaluate the fracture surface of the samples that had the best impact resistance (with and without SBS). A low vacuum Scanning electron microscope was used, model FEI Quanta 400.

Table 1. Compositions of the ABS/PCTG blends without SBS (control samples) and with SBS (compatibilized samples). Sample Control samples

Compatibilized samples

ABS (%) 67 33 50 67 64 47 44 36 33

PCTG (%) 33 67 50 30 30 50 50 60 60

SBS (%) 0 0 0 3 6 3 6 4 7

2.3. Characterization 2.3.1. Mechanical properties Tensile tests were performed according to ASTM D638, using test specimens of ASTM D638-Type I. Samples were tested on an Instron Model 5569 universal test machine using a load cell of 5,000 N at room temperature (23 °C). For impact strength, tests were performed on a Resil 25R instrumented impact machine from Ceast using a 1.0 J impactor in test specimens injected at 23 °C with a pendulum velocity of 2.90 m/s. The samples that presented the best performance on Impact Strength test were selected to be analyzed by DSC and SEM. 2.3.2. Differential Scanning Calorimetry measurements (DSC) For DSC, measurements were performed in the second heating from -120 °C to 250 °C at a ratio of 20 °C min-1 in dynamic nitrogen atmosphere (N2) with gas flow of 50 mL min-1 in Mettler’s model 822e equipment.

3. Results and Discussion 3.1. Mechanical properties Table 2 summarizes mechanical properties results of the ABS/PCTG blends with and without the addition of compatibilizing agent SBS. It can be observed that the sample 67/33 had the highest values of elastic modulus and Charpy impact strength, however it had the lowest elongation at break. This behavior was expected, since neat ABS terpolymer is more rigid than PCTG. On the other hand, increasing of the proportion of PCTG in blends with SBS causes an increase in elongation at break and in Charpy impact strength, but there is a decrease of the rigidity in relation to control samples, due the highest toughness of the elastomer copolyester. Figure 1 shows elastic modulus results obtained for the control and compatibilized blends with similar proportions. As it can be seen for the control blends, as the content of ABS in the blend increases, the elastic modulus also increases, indicating a higher rigidity. This behavior may be associated to greater rigidity of neat ABS compared to neat PCTG. Regarding the compatibilized blends, all the results showed that the addition of SBS increased the flexibility of the blends compared to the control samples, since the elastic modulus of the compatibilized blends decreased. Comparing 33/67/0 blend with the ABS/PCTG/SBS 36/60/4 and 33/60/7 blends, it can be observed a reduction on stiffness of approximately 5 % and 10 %, respectively. Therefore, the increase on the SBS content increases the flexibility of the final blend. All the compatibilized samples presented the same behavior. Figure 2 shows the elongation at break results of control and compatibilized blends. Comparing the blends ABS/PCTG 33/67 and ABS/PCTG/SBS 36/60/4, it is possible to notice a small increase of 5 % of elongation at break. On the other hand, 33/60/7 blend showed an increase of approximately 65 % in the elongation at break. The 47/50/3 blend had a greater increase in the elongation at break (189 %) compared

Table 2. Elastic modulus, elongation at break and Charpy impact strength of the ABS/PCTG blends with and without SBS. Sample Control samples

Compatibilized samples


ABS (%)

PCTG (%)

SBS (%)

67 50 33 67 64 47 44 36 33

33 50 67 30 30 50 50 60 60

0 0 0 3 6 3 6 4 7

Elastic modulus (MPa) 1,989 ± 7 1,904 ± 38 1,821 ± 20 1,924 ± 42 1,840 ± 39 1,826 ± 48 1,735 ± 24 1,732 ± 39 1,623 ± 14

Elongation at break (%) 11.0 ± 2 14.5 ± 6 43.0 ± 12 21.0 ± 3 31.0 ± 6 42.0 ± 14 28.0 ± 1 45.0 ± 23 71.0 ± 18

Charpy impact strength (J/m) 66 ± 17 63 ± 17 51 ± 12 65 ± 15 77 ± 17 78 ± 18 88 ± 18 108 ± 23 115 ± 18

Polímeros, 31(3), e2021033, 2021

ABS/Recycled PCTG blend compatibilized with SBS: effect on mechanical properties and morphology

Figure 1. Elastic modulus (MPa) of ABS/PCTG and ABS/PCTG/SBS blends at similar contents of ABS and PCTG.

Figure 2. Elongation at break (%) of ABS/PCTG and ABS/PCTG/SBS blends at similar contents of ABS and PCTG. Polímeros, 31(3), e2021033, 2021


Morali, J. A. & Brunelli, D. D. to the ABS/PCTG 50/50 mixture. The 33/60/7 blend also had an expressive increase of 181 % compared to the 33/67/0 control blend. It can be concluded that the increase in the ABS content in the control samples decreased the toughness of the blends. However, SBS compatibilizer attenuated this effect since there was an increase on the elongation at break of the compatibilized blends. Figure 3 shows Charpy impact strength results obtained for the compatibilized and control blends. It can be observed that the compatibilized blends produced using the following proportions of ABS/PCTG/SBS, 36/60/4 and 33/60/7, presented an increase on the Charpy impact strength of 111 % and 125 %, respectively, compared to 33/67/0 blend. However, the blends with a higher percentage of ABS, such as 47/50/3 and 44/50/6 compatibilized blends showed increase on this property of approximately 23 % and 33 %, respectively, when compared to the 50/50/0 control blend. The comparison of the 67/33/0 control sample to 67/30/3 and 64/30/6 blends showed smaller increases of approximately 2 % and 16 %, respectively. Therefore, results of Charpy impact strength test indicated that the increase on the ABS content resulted on the increase of this mechanical property, considering the non-compatibilized blend. Blends containing SBS presented an enhance on this mechanical property compared to control samples, corroborating the results of elastic modulus and elongation at break. It can be concluded that this effect is related to the elastomeric characteristic of the SBS compatibilizer that promotes better performance in blends with higher PCTG content. Among the control samples, the 67/33/0 blend was the one that presented the highest impact strength result whereas, among the compatibilized blends, the mixture

33/67/7 presented the best result for this property. These two compositions were evaluated by DSC and SEM techniques in order to evaluate the compatibility mechanism between the components.

3.2. Miscibility of polymeric blend 3.2.1. DSC analysis In order to evaluate the miscibility between the components of the blends with and without compatibilizer agent, samples that presented the highest results on Charpy impact strength tests were selected to be analyzed by DSC and SEM. Thus, the samples analyzed were: 67/33/0 and 33/60/7. For DSC analysis, samples of neat ABS and PCTG were also evaluated. According to Olabisi et al. [5], a blend can be considered miscible when the glass transition temperature (Tg) is unique and is in an intermediate range between the temperatures of the neat components. For a partially miscible blend, the glass transition temperature is given by two or more transitions that are corresponding to the blend phases and are shifted in relation to the Tg of the neat components. In this case, each phase is formed by a miscible mixture containing different compositions [5-8]. Table 3 shows the values o​​ f the glass transition temperature obtained for the analyzed samples. The 67/33/0 blend presented two different glass transition temperatures that were remarkably close to the neat ABS and PCTG components shifted by a few degrees. The same occurred to the 33/60/7 blend.

Figure 3. Charpy impact strength (J/m) of ABS/PCTG and ABS/PCTG/SBS blends at similar contents of ABS and PCTG. 4/8

Polímeros, 31(3), e2021033, 2021

ABS/Recycled PCTG blend compatibilized with SBS: effect on mechanical properties and morphology Table 3. Glass transition temperature of polymeric blends and their neat constituents. Sample (ABS/ PCTG/SBS) 100/0/0 67/33/0 33/60/7 0/100/0

Tg1 (°C)

Tg2 (°C)

Tg2 – Tg1

-81.9 74.3 77.0 83.1

102.6 107.5 110.2 -

33.2 33.2 -

Chen and Zhang [9,10] observed the similar behavior for ABS/PETG blends without compatibilizer. PCTG differs from PETG by the amount of the CHDM co-monomer. PETG is defined as a copolymer containing less than 50% (in weight) of CHDM comonomer; for contents higher than 50%, the material is defined as PCTG. In their work, the 70/30 ABS/PETG blend presented Tg1 of 74.3 °C and Tg2 of 110.4 °C while for the 30/70 ABS / PETG blend, the Tg results obtained were, respectively, 76.2 °C and 111.4 °C, corresponding to ABS and PETG phases, respectively. Considering Tg difference (Tg2 - Tg1), it is possible to notice that the 33/60/7 blend presented the same difference as the control blend. This leads to the conclusion that the 33/60/7 blend was not totally miscible. 3.2.2. Morphology and interface observations The morphology of a fracture surface may provide information about the compatibilization mechanism of polymeric blends. The component that is presented in the smallest proportion is called the dispersed phase while the component in the highest proportion is the continuous phase. When the dispersed phase is incompatible with the continuous phase, or matrix, it appears in spherical forms, large and with a well-defined interface [7,8]. Figures 4a and b show the fracture region of the impact specimen at 41 and 38x magnification, respectively, while Figures 4c and d show the micrographs at 5000x magnification and Figures 4e and f, with 10000x magnification, for 67/33/0 and 33/60/7 blends. It can be observed that the fracture surface of the compatibilized blend with SBS is rougher (Figure 4b), with several cracks, but with well-defined reliefs that indicates a ductile failure compared to the blend produced without the addition of compatibilizer. Figure 4c shows a good dispersion of the PCTG in the ABS matrix that has “fibers” form. The opposite is seen in Figure 4d, where the PCTG is the continuous phase (matrix), and the ABS is regularly dispersed in the mixture with different domain sizes. ABS, as an acrylonitrile, butadiene, and styrene terpolymer, presents a more heterogeneous aspect, due to its different co-monomers, while the PCTG has a smoother and more homogeneous surface. This aspect can be seen in Figures 4c and 4d, respectively. Micrographs of the 67/33/0 and 33/60/7 blends show an interface that do not present any adhesion between the phases as indicated by the circles in Figures 4 c-d. In addition, there are some regions, represented by the diamonds in Figures 4 c-d, with better adhesion in the interface. Micrograph of the 67/33/0 blend (Figure 4e) shows regions where PCTG-rich phase is dispersed in the ABS-rich Polímeros, 31(3), e2021033, 2021

matrix with a larger domain size and regions with a smaller domain size with a tendency to appear in the form of fibers, as indicated by the arrows in the micrograph. According to Zhang et al. [11], the smaller the diameter of the fiber, the better the adhesion at the interface and, consequently, the greater the tendency of the fibers to break or not to detach from the matrix. Thus, the 67/33/0 blend showed a lower result of impact strength due to the weak adhesion in the interface region between the components and due to the larger diameter of the dispersed phase in the matrix. However, the interface was not weak enough to allow the complete detachment of the PCTG fibers from the ABS matrix. Joseph et al. [12] studied ABS/PETG blends using SBS as a compatibilizer; the mixture of ABS in PETG as a dispersed phase presented larger domain sizes than when PETG is the dispersed phase. The phenomenon of coalescence is more pronounced in high concentrations of the dispersed ABS phase, due to the high mobility of the ABS domains in the PETG matrix. This phenomenon can also explain the morphology of the 33/60/7 ABS/PCTG/SBS blend. Analysis of the micrograph of the 33/60/7 blend (Figure 4f) suggests that there are rubber small domain inclusions dispersed in the PCTG-rich matrix (arrows) and regions of irregular sizes of the dispersed phase, ABS, suggesting phase distortion and co-continuous morphology. Bo Li et al. [13] studied the toughening mechanism of the blend ABS/PETG and a co-continuous morphology was also evidenced. The toughening mechanism with this co-continuous morphology can be explained by energy dissipation with the interfacial debonding, energy absorption with the distortion of ABS phase [13] and the presence of elastomeric phase, from SBS, which influences in the crack propagation process, turning the material more ductile (Figure 5). According to the model proposed by Macosko [14], a co-continuous morphology should enhance the fluidity, toughening and rigidity of the final blend, depending on the properties of the components and the dispersion of energy, corroborating the results of mechanical properties obtained previously for impact resistance. Figure 6 summarizes how the blends components acted in the interface and on the compatibilization of the systems ABS/PCTG and ABS/PCTG/SBS, at the proportions 67/33/0 and 33/60/7, respectively. Therefore, the evaluation of all micrographs suggests that the main mechanism of action of the SBS in the blend was by dispersion and increase of rubber content, which attenuates the transference of mechanical stresses to the matrix and increases the impact strength and by the cocontinuous morphology of the ABS-rich phase, which can enhance the toughness of the material. According to the theory proposed by Joseph et al. [11], if the compatibilization of ABS in the PCTG had occurred by emulsification, the morphology would be more homogeneous and with smaller domains. Considering the present study, for regions where it is possible to identify a better affinity between the ABS-rich phase and PCTG-rich matrix, there may have been a small 5/8

Morali, J. A. & Brunelli, D. D.

Figure 4. Micrographs of the fracture surface of the ABS/PCTG/SBS samples at different magnifications.

Figure 5. Crack propagation process in a brittle matrix and in a tenacified matrix. 6/8

Polímeros, 31(3), e2021033, 2021

ABS/Recycled PCTG blend compatibilized with SBS: effect on mechanical properties and morphology

Figure 6. Scheme of the morphology and mechanism of action of the compatibilizers for the ABS/PCTG and ABS/PCTG/SBS blends at the following contents: (a) 67/33/0; (b) 33/60/7, respectively.

interaction between the butadiene phase of SBS and PCTG, as well as an interaction of the styrene phase with the ABS. Nevertheless, this was not the main mechanism responsible for the results that were found.

4. Conclusion According to the results, it was found that the inclusion of SBS as a compatibilizing agent in the blend between ABS and PCTG, promoted an increase on the toughness of the blend and it enhanced impact strength, when compared to the control blend. It was observed that, in the 67/33 ABS/PCTG blend, PCTG is dispersed in the ABS matrix in the form of fibers of different sizes and homogeneous distribution with regions where the interface is easily observed and regions with a certain affinity of the phases. For 33/60/7 ABS/PCTG/SBS blend, it was verified the coalescence of the dispersed ABS phase in the PCTG matrix, since that the SBS acted, primarily, as an impact modifier in the matrix, but there may also exist a small interaction between them, this corroborates with the results of glass transition temperature obtained, in which there was a slight displacement in relation to the neat polymers. Therefore, it is suggested that SBS promoted a toughening effect due to the increase in the rubber content in the final blend, since a significant increase in elongation at break was observed, although there was a decrease in the elastic modulus. Polímeros, 31(3), e2021033, 2021

Thus, it can be concluded that the blend between virgin ABS and residues of PCTG using SBS as compatibilizer, in the proportion of 33/60/7, can be used in several applications, like packaging and home appliances, for example, once that the impact properties have improved with the toughening.

5. Acknowledgements Thanks to ITA for the support provided for the preparation of the study.

6. References 1. López, R. (1994). The environment as a factor of production: the effects of economic growth and trade liberalization. Journal of Environmental Economics and Management, 27(2), 163-184. http://dx.doi.org/10.1006/jeem.1994.1032. 2. Dewil, R., Everaert, K., & Baeyens, J. (2006). The European plastic waste issue: trends and toppers in its sustainable re-use. In: Proceedings of the 17th International Congress of Chemical and Process Engineering (pp. 27-31). Prague: CHISA Secretariat. 3. Olabisi, O., & Adewale, K. P. (2016). Handbook of thermoplastics. USA: CRC Press, Taylor & Francis Group. http://dx.doi. org/10.1201/b19190. 4. Granado, A., Eguiazabal, J. I., & Nazabal, J. (2006). High Compatibility and improved barrier performance in blends based on a copolyester modified with a poly(amino ether) resin. Macromolecular Materials and Engineering, 291(9), 1074-1082. http://dx.doi.org/10.1002/mame.200600159. 5. Olabisi, O., Robeson, L. M., & Shaw, M. T. (1979). Polymerpolymer miscibility. USA: Academic Press, Inc. 7/8

Morali, J. A. & Brunelli, D. D. 6. Koning, C., Van Duin, M., Pagnoulle, C., & Jerome, R. (1998). Strategies for compatibilization of polymer blends. Progress in Polymer Science, 23(4), 707-757. http://dx.doi.org/10.1016/ S0079-6700(97)00054-3. 7. Utracki, L. A. (2003). Introduction to Polymer Blends. In:, L. A. Utracki. Polymer blends handbook (pp. 1-122). Netherlands: Kluwer Academic Publishers. 8. Utracki, L. A. (2002). Compatibilization of Polymer Blends. Canadian Journal of Chemical Engineering, 80(6), 1008-1016. http://dx.doi.org/10.1002/cjce.5450800601. 9. Chen, T., & Zhang, J. (2018). Compatibilization of acrylonitrile-butadiene-styrene terpolymer/poly(ethylene glycol-co-1,4-cyclohexanedimethanol terephthalate) blend: effect on morphology, interface, mechanical properties and hydrophilicity. Applied Surface Science, 437, 62-69. http:// dx.doi.org/10.1016/j.apsusc.2017.12.168. 10. Chen, T., & Zhang, J. (2016). Surface hydrophilic modification of acrylonitrile-butadiene-styrene terpolymer by poly(ethylene glycol-co-1,4-cyclohexanedimethanol terephthalate): preparation, characterization, and properties studies. Applied Surface Science, 388(Pt A), 133-140. http://dx.doi.org/10.1016/j. apsusc.2016.02.242.


11. Zhang, X., Li, B., Wang, K., Zhang, Q., & Fu, Q. (2009). The effect of interfacial adhesion on the impact strength of immiscible PP/PETG blends compatibilized with triblock copolymers. Polymer, 50(19), 4737-4744. http://dx.doi. org/10.1016/j.polymer.2009.08.004. 12. Joseph, S., Focke, W. W., & Thomas, S. (2010). Compatibilizing Action of a Poly(styrene–butadiene) Triblock Co-polymer in ABS/PET-G Blends. Composite Interfaces, 17(2-3), 175-196. http://dx.doi.org/10.1163/092764410X490590. 13. Li, B., Zhang, X., Zhang, Q., Chen, F., & Fu, Q. (2009). Synergistic enhancement in tensile strength and ductility of ABS by using recycled PETG plastic. Journal of Applied Polymer Science, 113(2), 1207-1215. http://dx.doi.org/10.1002/ app.30002. 14. Macosko, C. W. (2000). Morphology development and control in immiscible polymer blends. Macromolecular Symposia, 149(1), 171-184. http://dx.doi.org/10.1002/15213900(200001)149:1<171::AID-MASY171>3.0.CO;2-8. Received: Sep. 17, 2021 Revised: Nov. 12, 2021 Accepted: Nov. 13, 2021

Polímeros, 31(3), e2021033, 2021

ISSN 1678-5169 (Online)


Role of cellulose nanocrystals in epoxy-based nanocomposites: mechanical properties, morphology and thermal behavior Nayra Reis do Nascimento1 , Ivanei Ferreira Pinheiro1 , Guilherme Fioravanti Alves1 , Lucia Helena Innocentini Mei1 , José Costa de Macedo Neto2*  and Ana Rita Morales1*  Laboratório de Blendas e Compósitos Poliméricos, Departamento de Engenharia de Materiais e Bioprocessos, Universidade Estadual de Campinas - UNICAMP, 13083-852, Campinas, SP, Brasil 2 Laboratório de Materiais e Processamento, Departamento de Engenharia de Materiais, Universidade do Estado do Amazonas - UEA, 69050-020 Manaus, AM, Brasil


*morales@unicamp.br; jmacedo@uea.edu.br

Abstract This study evaluated the influence of cellulose nanocrystals (CNC) content on the properties of epoxy nanocomposites. The CNC were obtained from microcrystalline cellulose by acid hydrolysis. 4.0, 5.5 and 7.0% of untreated CNC were incorporated into epoxy resin. Sonication was used to disperse the CNC in the resin. The thermal stability, the glass transition temperature and the degree of conversion were reduced as observed by Thermogravimetry and Differential Scanning Calorimetry, respectively. The tensile and bending modulus showed no significant improvement and the impact resistance showed a slight reduction due to the non-uniform dispersion of the CNCs, as observed by Transmission Electron Microscopy. Analysis of Scanning Electron Microscopy showed a change of the fracture mechanism of the epoxy resin: the CNCs increased the elastic modulus by reinforcement, but accelerated the fracture by acting as defects. The Halpin-Tsai model was applied to predict the elastic modulus of the epoxy/CNC system. Keywords: cellulose nanocrystals, nanocomposites, epoxy resin. How to cite: Nascimento, N. R., Pinheiro, I. F., Alves, G. F., Mei, L. H. I., Macedo Neto, J. C., & Morales, A. R. (2021). Role of cellulose nanocrystals in epoxy-based nanocomposites: mechanical properties, morphology and thermal behavior. Polímeros: Ciência e Tecnologia, 31(3), e2021034. https://doi.org/10.1590/0104-1428.20210057

1. Introduction The development of high-performance polymers has focused on materials that are reinforced with nanoparticles that can improve the properties in relation to those made with pure materials. Polymeric nanocomposites are materials that contain fillers, in which at least one of their dimensions has measurements that is less than 100 nm, dispersed in a polymer matrix[1]. They are part of a class of engineering materials that are very attractive because they provide better mechanical and thermal properties than conventional composites. In addition, only small amounts of reinforcement are required[2,3]. Cellulose nanocrystals (CNCs) are one of these particles that are used because they have a set of peculiar characteristics that characterize them as a good polymer matrix-reinforcing agent. These characteristics are the following: a high amount of hydroxyl groups that can provide the possibility of very good interaction in the polymer/filler interface; a high specific surface area and high aspect ratio that is important to stress transfer; high crystallinity and elastic modulus, which are responsible for the increase in stiffness; as well as high thermal stability, low cost, low density and a peculiarly interesting morphology[4].

Polímeros, 31(3), e2021034, 2021

Due to the excellent properties of CNCs, many studies have been developed that show their use as nanofillers in polymers. The high polarity and high hydrophilicity nature of CNCs causes the interaction and dispersion in some types of polymeric matrices to be compromised[4]. Polar matrices can present better dispersion of the CNCs due to the high interaction of hydrogen bonds, thus obtaining better mechanical properties[5], while nonpolar matrices can present worse dispersion and reduced mechanical properties. Many studies have shown that CNC surfaces treated with organic compounds have a better interaction with non-polar polymers[6]. However, in order to promote better interaction with the polymeric matrix, treated CNCs are used in liquid suspension[7]. Another factor that influences the dispersion of CNCs in the polymer matrix is the method used for obtaining it. Some methods of obtaining nanocomposites use complex equipment such as extrusion and hot pressing[7]. In this study, the CNCs were inserted in the polymeric matrix in the form of a dry powder and the nanocomposites were obtained by a simple method of manual mixing at room temperature. CNCs are extracted from cellulose, which is the most abundant natural polymer in nature. Among other sources,



Nascimento, N. R., Pinheiro, I. F., Alves, G. F., Mei, L. H. I., Macedo Neto, J. C., & Morales, A. R. they can be extracted from microcrystalline cellulose (MCC), which is a substance obtained from partially depolymerized cellulose. It can be dried to a fine, powdery particle form or co-processed with a water-soluble polymer to form a colloid. MCC is widely used because it is chemically stable and physiologically inert[8]. As such, there has been increased interest in its use in the pharmaceutical[9], food industries and the energy sector[10]. Epoxy resin was chosen as the nanocomposite matrix for this study. Epoxy resin is part of an important class of highperformance thermosetting polymers that are widely used in automotive, construction and aerospace applications[11]. It is considered one of the best matrices since it has excellent properties such as good adhesion[12], good mechanical properties such as high elastic modulus, shear strength, corrosion resistance, thermal stability, low shrinkage after the curing process, favorable viscosity properties, light weight, low cost, resistance to environmental degradation, as well as good friction and wear resistance[13]. The incorporation of CNCs in the epoxy resin can result in a high stiffness and good fatigue strength, which, for example, is very desirable in aerospace applications. Studies in which CNCs from tunicates and cotton were incorporated into the epoxy matrix showed good dispersion and consequent improvement in mechanical properties due to the increased in interfacial adhesion between the CNCs and the epoxy matrix[14]. Roszowska-Jarosz et al.[15] describe the production of epoxy/ CNC nanocomposites by adding a suspension of CNCs and acetone in the quantities of 0.5%, 1% and 1.5%wt in epoxy resin, for which they obtained superior values for impact resistance and flexural strength. Nissilä et al.[16] produced nanocomposites by impregnating cellulose nanofibers (CNF) in bio-epoxy resin. For this, they produced suspensions in water in the quantities of 1.0 and 1.5%wt of CNF and then they prepared aerogels by freezing the suspensions in molds. The frozen suspensions were impregnated with the bio-epoxy resin using a vacuum and its curing was carried out at 80 °C. Data on the mechanical properties were obtained by tensile and flexion tests in which the nanocomposites showed an increase in the flexural modulus of approximately 50%, and of 20% in the tensile strength in relation to the polymer without CFN. Qi et al.[17] produced CNC/epoxy nanocomposites using superficially modified CNCs by grafting with poly (n-vinylpyrrolidone), which significantly increased the toughness of epoxy resin and is explained by the excellent dispersion of the CNCs and its effect in promoting plastic deformation in the system. Some other works mention the modification of epoxy resin with CNCs identifying, in addition to changes in mechanical properties, changes in thermal characteristics[18]. The objective of this study was to characterize the CNCs obtained from MCC and evaluate the mechanical properties of epoxy/CNC nanocomposites considering the influence of the CNC volumetric fraction, taking into account fractions in above, and below the percolation threshold. The HalpinTsai modulus of elasticity prediction model was applied as a tool to understand the interactions of the components and the mechanical behavior of the system under study. 2/13

2. Material and Methods 2.1 CNC production Commercial microcrystalline cellulose (MC 500, Blanver, Brazil) was used as a source of cellulose for CNC extraction. Sulfuric acid (Synth) was used in the acid hydrolysis process. The nanocomposites were produced using DER 331 bisphenol A diglycidyl ether-based (DGEBA) epoxy resin and the hardener DEH 24, a triethylenediamine (Dow Chemical Company), which were mixed in the ratio of 100:13, as indicated by the supplier. Acid hydrolysis consists of a process that removes the amorphous phase of cellulose so that the nanocrystals are isolated. Microcrystalline cellulose was hydrolyzed with sulfuric acid solution (55% v/v) at 45 °C for 30 min under constant stirring, then, 300 mL of distilled water was added to stop the reaction, following the Hassan method[19]. The ratio was 10 mL of sulfuric acid solution to 1 g of microcrystalline cellulose. Centrifugation was applied to partially remove sulfuric acid and separate the nanocrystals from the solution. The supernatant that appeared after 10 minutes of each centrifugation cycle at 6000 rpm was removed and replaced with distilled water. This exchange procedure was performed repeatedly until the solution reached a pH close to 7. The resulting suspension was submitted to dialysis. The material was placed on the dialysis membrane so that the exchange with the aqueous medium occurred. The water, where the membranes were immersed, was replaced once after 24 h and dialysis continued for a further 24 h. As the last step, lyophilization consisted of the total removal of water using a freeze dryer, which left only the powdered nanocrystals. From this step, the CNCs were ready to be incorporated into the epoxy matrix. During this stage, CNCs tend to cluster together in structures that are stabilized by hydrogen interactions. Freeze-drying was the method of drying used, and this allows the dispersed CNC material structure to be maintained.

2.2 Nanocomposites production Nanocomposites with a CNC volumetric concentration of 4.0, 5.5 and 7.0% were prepared. These concentrations were determined from the theoretical calculation of the percolation threshold[20]. The CNCs were weighed and carefully added to the epoxy resin, then manually mixed at room temperature, according to the work of Omrani et al. [21] . An ultrasonic probe (QSONICA, Q700, USA) was used with a maximum power of 700 W and a 6.4 mm diameter sonotrode for 10 min. to fully disperse the CNCs in the polymer matrix under 30 W power conditions. Then, the hardener was added and mixed evenly with the resin. Figure 1 shows a schematic of the steps in the production of nanocomposites. The mold used was produced in aluminum alloy with dimensions of 30 × 30 × 1.5 cm and its cavities in the shapes and dimensions of specimens for mechanical tests were produced by machining (Figure 2). As a release agent, PVC film was used inside the mold. The mixture was poured into the mold and kept at room temperature at atmospheric pressure for curing. Polímeros, 31(3), e2021034, 2021

Role of cellulose nanocrystals in epoxy-based nanocomposites: mechanical properties, morphology and thermal behavior

Figure 1. (a) Manual mixing of CNCs and epoxy resin; (b) ultrasonic CNC dispersion; (c) manual addition and mixing of hardener; (d) the mixture poured into the mold; (e) specimen after curing and (f) mechanical testing.

Figure 2. Metal mold and specimens produced.

2.3 CNCs and nanocomposite characterization The samples of CNCs obtained were analyzed by using Fourier transform infrared spectroscopy (FTIR) (Nicolet 6700, Thermo Scientific, USA). The spectra were obtained in the ATR mode, with the SMART OMNI-SAMPLER accessory operating in the 4000-675 cm-1 range, with 4 cm-1 resolution and 128 scans. Transmission electron microscopy (TEM) (JEM 2100, JEOL, Japan) was used to evaluate the morphology of the CNCs. To obtain the TEM image, a drop of the diluted CNC suspension was allowed to dry on a carbon-coated copper grid (400 mesh). To improve contrast, a 2% by weight solution of uranyl acetate was placed on the grid for 30 seconds. For determination of the aspect ratio (L/D), dimensions of the CNCs were obtained from TEM images with the aid of Image J software. For this calculation, about 60 units of CNCs were used. Polímeros, 31(3), e2021034, 2021

The crystal structure of the CNCs was evaluated using X-ray diffraction analysis (XRD), (X’Pert-MPD, Philips, USA), operating with CuKα (λ = 1.54060 Å), 40 kV voltage and 40 pA current, 2θ between 5 and 50º. The samples were placed in the XDR equipment in dry powder form. The degree of crystallinity was estimated following the method cited by Martin et al.[22] as it is widely used for natural fibers. The index (Ic) was estimated by means of Equation 1. I Ic = 1 −  am  I002

  x100 


Using the peaks (I002, 2θ = 22.4 º) and (Iam, 2θ = 18.8º). I002 represents both the crystalline and the amorphous material, while Iam represents amorphous material. The nanocomposites obtained were analyzed using FTIR (Nicolet 6700, Thermo Scientific, USA). Infrared spectra were obtained in ATR mode, with the SMART OMNI3/13

Nascimento, N. R., Pinheiro, I. F., Alves, G. F., Mei, L. H. I., Macedo Neto, J. C., & Morales, A. R. SAMPLER accessory operating in the 4000-675 cm-1 range, with 4 cm-1 resolution and with 128 scans. To evaluate the thermal behavior, DSC equipment (DCS1, Mettler Toledo, Switzerland) was used with a heating rate from 10 °C/min from -30 to 300 °C and an inert atmosphere (N2) with a flow of 50 ml/min. Firstly, the total cure enthalpy of the base system was obtained from the DSC analysis of the resin + hardener in the range of -30 to 300 °C in the proportions indicated by the manufacturer, which was found to be 4592 mJ. The nanocomposites were also evaluated using a thermogravimetric analyzer (TG) (TGA-50M, Shimadzu, Japan) and a microanalytical scale (MX5, Mettler Toledo, Switzerland). Analyses were performed in an inert atmosphere (N2), in a temperature range of 25 to 600 °C, and a heating rate of 10 °C/min. The morphology and size distribution of the particles in the CNC filled epoxy nanocomposites were studied by using TEM (TECNAI T20, FEI, USA) with 200 kV. In this analysis, the samples with 4.0, 5.5 and 7.0% of CNCs was cut in thin sections (around 120 nm in thickness) using an ultramicrotome (UltraCut, Leica, Austria) and placed on a copper grid of 3 mm in diameter. A universal testing machine was used to obtain the tensile mechanical properties of modulus of elasticity (ET), tensile strength (σT) and strain at break (εT). Samples were prepared according to ASTM D638-10, Type I, and an extensometer (High Elongation extensometer AHX850, MTS, USA) was used. For each formulation, five specimens were tested and the average value was calculated. The universal test machine (Alliance RT/5, MTS, USA) operated at a crosshead speed of 1 mm/min and a 5 kN load cell was used. Flexural tests of the samples were performed and flexural modulus (EF) and flexural strength (σF) were evaluated. For each formulation, five specimens were tested and the average value was calculated. The 3-point bending tests were performed at room temperature according to ASTM D790-03. A universal testing machine (5984, Instron, USA) was used with a 150 kN load cell and at speed of 2 mm/min. For determination of the impact strength (IS) the samples were prepared according to ASTM D4812-19 and

the notched Izod test was performed using an impact tester (Impact 104, Tinius Olsen, USA) with capacity of 25 J and a 10 º impact hammer. The nanocomposite fractured surface morphology was studied using scanning electron microscopy (SEM) (LEO 440i, Leica, England) at 10 mA operating current. The samples from the tensile test were cut at 5.0 mm (manual cutter) below and parallel to the fracture. The cut piece was fixed with carbon adhesive on the SEM sample holder and was gold sputter coated with a sputter coater (SC7620, Polaron VG Microtech, UK) with a film of approximately 20-50 nm.

Figure 3. CNC and MCC FTIR spectra.

Figure 4. Diffractogram of the CNCs and the MCC.


3. Results and Discussion 3.1 CNC characterization Figure 3 shows the comparative FTIR spectra of the MCC and the CNCs. In the MCC spectrum, the band around 3400 cm-1, which is due to the OH-stretching vibration, is typical for adsorbed water and gives considerable information concerning the hydrogen bonds. The band at 2900 cm-1 corresponds to the stretching of the C-H bonds. The bands at 1430, 1059 and 897 cm-1 are typical of pure cellulose and can be seen in both the MCC and CNC spectrum. In general, the 897 cm-1 band is assigned to vibrations of the C-O-C elongations of the characteristic β (1→4) glycosidic bond[23]. Therefore, typical cellulose bands can be seen in both the MCC and CNCs[24]. Figure 4 shows the results of X-ray diffraction of the CNCs and the MCC. A peak at 22.42° can be observed and indicates the presence of crystallinity that is characteristic of this nanoparticle. This peak refers to cellulose type I and approaches the value of 22.7º found in the literature[25]. This means that the CNCs were isolated from other parts of the cellulose and that the acid hydrolysis process did not change their structure during treatment. The degree of crystallinity obtained was 82.7%. The degree of crystallinity was calculated following the method cited by Martin et al.[22] and is widely used for natural fibers. One shoulder can also be observed on the diffractogram around 20º, which refers

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Role of cellulose nanocrystals in epoxy-based nanocomposites: mechanical properties, morphology and thermal behavior to cellulose type II that could be present in low amounts. At peak 22.5° of the MCC, a lower intensity is observed and this may indicate the presence of amorphous regions[26]. Figures 5a-c show the image from the TEM that was performed after the CNCs were lyophilized. It can be observed that the nanocrystals were obtained successfully. The image shows the needle-like nanoparticles that correspond to the characteristic and expected morphology of the CNCs. It can also be noted that there was some level of agglomeration of particles and consequent stacking on each other. The CNC agglomeration phenomenon is mainly due to their very small size and their very high surface energy. In addition, the hydroxyl group on the surface of the CNCs facilitates the formation of hydrogen bonds, which facilitates their approximation. The morphology of the CNCs is shown in Figure 5. The length (L) and diameter (D) of the CNCs were (142±36) nm and (11±2.5) nm, respectively. A total of 60 length and diameter measurements were taken and then the aspect ratio (L/D) was calculated from each measurement. Thus, the aspect ratio obtained was 12.5. A high aspect ratio results in high rigidity of the percolated nanoparticle network and this favors the reinforcing effect on a nanocomposite since a good dispersion is obtained in polymer matrices. In order to have a good transfer of tensions between the matrix and the reinforcement, the minimum value of the aspect ratio must be 10[27], which was obtained in this work however, other aspects, such as the compatibility between the matrix and the dispersed phase, the percolation limit and the dispersion, must be considered[20]. CNCs have a tendency to aggregate, as seen in Figure 5.

As Tg was influenced by the presence of the CNCs, it can be assumed that the nanocomposite can be processed under the same conditions as the neat system and that the thermal behavior under the application conditions would be close. The samples were cured at room temperature. An important point is that, for all samples, the relaxation enthalpy was observed and the cure was not complete. The conversion for each sample was determined following the method of Hardis et al.[29]. The values are shown in Table 1, in which it is observed that the conversion is very similar for all samples, although the presence of (-OH) groups of cellulose may participate in the curing reaction of the epoxy resin and result in increased conversion[29]. It is possible that this effect was diminished by the spatial restriction caused by the CNCs. In order to understand the chemical structure of epoxy resin and especially nanocomposites in more detail, the FTIR spectra were evaluated. Figure 6 shows the spectra of epoxy and nanocomposites with different CNC contents. The epoxy group has three characteristic absorption bands in the infrared spectrum. One is the band appearing in the 1250 cm-1 range, which corresponds to the symmetrical axial deformation of the epoxy ring and the aromatic ether group C-O-C, in which all its bonds expand and contract in phase[30]. Another is the 916 cm-1 band, which is related to the asymmetric axial deformation of the ring, in which the C-C bond increases and the C-O bond contracts. The third band appears in the 840-750 cm-1 range and corresponds to the C-O-C bond[31]. The absorption present in the 30003500 cm-1 range refers to the hydroxyl groups of adsorbed water that were identified in all the nanocomposites. The absorption around 2925 cm-1 is due to the stretching of C-H

3.2 Nanocomposite characterization In order to understand the relationship between processing, morphology, properties and durability of a thermoset material and, consequently, of the composite, it is necessary to understand the cure mechanism of the system. Table 1 summarizes the DSC data. Tg values of the nanocomposites were reduced by 5 °C for the samples with 5.5 and 7.0% of CNCs in relation to pure epoxy. Residues of water molecules on the surfaces of CNC aggregations may have affected the crosslinking of the epoxy matrix[28].

Table 1. Results of thermal analysis: glass transition temperature (Tg), residual enthalpy (∆H R) and conversion degree (X). *V CNC (%)

Tg (°C)

0 4 5.5 7

67 67 63 63

∆H R (mJ/mg) 123 182 249 243

X (%) 97 96 95 95

*CNC volumetric concentration (%).

Figure 5. (a), (b), (c) Transmission electron micrographs of the obtained CNCs. Polímeros, 31(3), e2021034, 2021


Nascimento, N. R., Pinheiro, I. F., Alves, G. F., Mei, L. H. I., Macedo Neto, J. C., & Morales, A. R. bonds. Strong absorptions between 1000 and 1250 cm-1 indicate the C-O-C (ether) or C-O-H groups. The bands in the 2900-2930 cm-1 range indicate stretching of C-H bonds. Monitoring the band at 916 cm-1 allowed us to evaluate the existence of residual epoxy groups, which indicate if the cure was complete with the consumption of most groups. As can be seen in Figure 6, this band almost disappeared in the nanocomposite spectra. This means that there was a small cure residual, as observed in the DSC analysis. All the CNCs/epoxy nanocomposites at different filler loadings displayed similar FTIR spectra. The interfacial interaction of the CNCs and epoxy resin can be evaluated using FTIR. A reduction in bands between 3000-3500 cm-1 was observed for nanocomposites. These bands refer to the hydroxyl groups adsorbed on the surface of the CNCs. This behavior indicates that there may have been a low interfacial interaction between the CNCs and the epoxy in the production of the nanocomposites[17], although the CNCs were expected to be trapped in the cross-linked epoxy resin and completely covered by the polymer chains via hydrogen bonds[18]. The TEM images confirm the low interaction and, as a consequence, the formation of clusters. This low interaction may have affected the mechanical and

thermal properties of the nanocomposites, making them inferior to CNC-free epoxy and this will be discussed later. The thermal stability of the pure epoxy resin and the nanocomposites was evaluated using TG analysis. Figure 7 shows the degradation temperature curves, and the results are presented in Table 2. It is observed that there was a initial mass reduction for temperatures between 50-200 °C, which was due to moisture loss from the dehydration of secondary alcoholic groups and the evaporation of physically weak and loosely bound moisture on the surfaces of the composites[32]. Table 2 shows that the moisture loss was greater for the nanocomposites than for the pure epoxy. This behavior suggests that the CNCs adsorbed water on their surface. It was also observed that the moisture loss values are close to those of the nanocomposites, which suggests that the amount of water did not increase when the CNC fractions increased. The interaction of the epoxy resin matrix acted as a water barrier that prevented the attraction with the hydrophilic hydroxyls on the surface of the CNCs. CNC aggregations in the epoxy matrix may also have prevented access to water on the surface of the CNCs[33]. Thermal decomposition of pure epoxy resin occurs in two stages[34]. The main mass loss occurs in the range of 300 to 500 °C, which involves two stages of degradation. The initial loss peak around 350 °C corresponds to the degradation of the aliphatic amine curing agent due to its low C-N bond breakage energy. The second peak occurs in the temperature range of 405 °C and is attributed to resin decomposition[18]. Table 2. Results of the TG analysis. VCNC (%)**** 0 4 5.5 7

Figure 6. FTIR spectra of pure epoxy resin and nanocomposites.

Tmoisture loss Moisture loss (%) (°C) 58-150 62-205 44-126 65-200

1.9 2.2 2.1 2.0

Tonset* (%) 356 354 344 353

T5%** (°C) 338 336 312 319

Tmax*** (°C) 390 397 382 387

Tonset*: the onset temperature of weight loss; T5%**: the temperature corresponding to 5% of weight loss; Tmax***: the temperature corresponding to the maximum decomposition rate; ****VCNC volumetric concentration (%).

Figure 7. (a) TG curves and (b) Derivative TG analysis of pure epoxy resin and nanocomposites. 6/13

Polímeros, 31(3), e2021034, 2021

Role of cellulose nanocrystals in epoxy-based nanocomposites: mechanical properties, morphology and thermal behavior Comparing the TG analysis curves of the pure epoxy system with the nanocomposites, it can be observed that the thermograms are similar. No specific degradation step of the CNCs is observed in the nanocomposite thermograms, regardless of the volumetric concentration. It was observed for the nanocomposites that as the CNC content increased, all initial degradation temperatures decreased in relation to the pure epoxy. The T5% for the pure epoxy was 338 °C, and this was reduced when CNCs were added for the range of 336 to 319 °C. This behavior has already been described in the literature, in which Tonset is around 320 °C and it decreased when CNCs were added to the epoxy matrix, and is due to the lower Tonset temperature of the CNCs, which is around 230 °C[17,35]. Moreover, for Tmax, there was no relevant difference between the pure epoxy and the nanocomposites. Some of the divergent results due to the addition of CNCs can be explained by a complex energy dissipation mechanism of the interfaces between the nanocrystals and the matrix[19]. The 5.5% nanocomposite, which corresponds to the percolation threshold volume fraction, presented the lowest thermal stability of all samples. The slight decrease in Tmax for the 5.5 and 7.0% nanocomposites may have been due to increased thermal conductivity of the samples with the addition of the CNCs[36,37].

3.3 Mechanical testing There are many factors that directly affect the mechanical properties of nanocomposites, such as the inherent properties of the matrix and fillers, preparation conditions, interfacial adhesion (matrix-filler interaction), nanoparticle dispersion quality, and their concentration, which must be above the percolation threshold. Table 3 shows the results of the following mechanical properties: tensile modulus (ET), flexural modulus (EF), elongation in tensile (ԐT), elongation in flexural (ԐF), tensile strengths (𝛔T), flexural strengths (𝛔F) and impact strength (IS). The values of the modulus, the tensile strength and deformation for the tensile and flexural tests of the neat DGEBA/DEH24 are in accordance with other work in the literature[38], but differ from those shown for the standard material[39], which has a higher flexural modulus of 3.0 GPa and lower tensile strength of 79 GPa. Table 3 also shows that the nanocomposites’ tensile modulus and flexural modulus values were higher than that of the neat epoxy. This behavior indicates that the high surface area of the CNCs and the presence of O-H groups in their surface caused, at a certain level, good interaction of the CNCs with the epoxy resin molecules[17], and includes possible chemical bonding between the CNCs and the epoxy matrix[40]. The values of elongation and tensile and flexural strengths have a tendency to reduction with the increase in the CNC concentration. Two main factors can explain this behavior: 1) the surface characteristics of the CNCs are an important

issue in order to improve the mechanical properties, and in some cases their modification can be suggested[41] and 2) the dispersion must ensure the availability of the surface for an adequate transfer of mechanical stress from the matrix to the filler. CNCs have a rigid nature and a tendency to form agglomerates that act as micro-particles distributed in the epoxy matrix, which caused stress concentration points in the epoxy matrix and made it rigid and hard[14,42]. Then, as the fraction of CNCs increased, the distribution of these agglomerates in the matrix also increased, and this caused the tensile and flexural moduli to increase, although in a lower proportion than expected. Furthermore, the agglomerates act as defects causing a premature fracture that reduces the elongation and strength. The TEM images shown in Figures 8a-c confirm the presence of the agglomerates in the nanocomposites. The impact energy values for the nanocomposites were lower than those of the neat epoxy and the values decreased with the increase in the concentration of CNCs up to the concentration of 5.5% and, for the 7% concentration, there is a certain recovery in this property. This is in accordance with the results for the tensile and flexural values.

3.4 Mechanical model for elastic modulus prediction Mechanical models help to validate experimental data. Most of them assume full or partial bonding between the reinforcement filler and the matrix. Among the best known for predicting the modulus of elasticity are the Mori-Tanaka models (based on the theory of inclusion and Eshelby’s inclusion) and the Halpin-Tsai method (an interpretation of the Hill Potential theory). The latter is the most widely used to predict the modulus of elasticity of composites and

Figure 8. Theoretical and experimental tensile modulus of nanocomposites with different volume fractions.

Table 3. Mechanical properties of neat epoxy and nanocomposites. V CNC (%)* 0 4 5.5 7

ET (GPa)

1.70 ± 0.13 1.77 ± 0.11 1.94 ± 0.14 2.15 ± 0.35

𝛔T (MPa) 96 ± 23 62 ± 12 57 ± 12 51 ± 15

ԐT (%)

2.7 ± 0.9 2.0 ± 0.1 1.5 ± 0.6 1.3 ± 0.6

EF (GPa)

1.60 ± 0.15 2.30 ± 0.01 1.90 ± 0.17 2.10 ± 0.13

𝛔F (MPa) 73.2 ± 1.8 61.6 ± 3.1 50.6 ± 10.8 67.6 ± 8.5

ԐF (%)

3.2 ± 0.8 3.5 ± 0.1 3.1 ± 0.8 3.3 ± 0.2

IS (J/m) 4.8 ± 0.3 3.9 ± 0.1 3.7 ± 0.6 4.3 ± 0.3

*CNC volumetric concentration (%).

Polímeros, 31(3), e2021034, 2021


Nascimento, N. R., Pinheiro, I. F., Alves, G. F., Mei, L. H. I., Macedo Neto, J. C., & Morales, A. R. is a relatively simple universal model that can be applied to various compounds[43]. In addition, it obtains and considers the form and aspect ratio of the reinforcement. The modified Halpin-Tsai model is presented in Equation 2. Coleman et al. [44] states that this model provides better prediction of elastic modulus for low volumetric fractions.


3 8

 lf  1 + 2   n L Vf (2) 5 1 + 2 n T Vf  df  Em + Em 1 − n L Vf 8 1 − n T Vf

In which: Ec: elastic modulus of composites. Em: elastic modulus of epoxy matrix. Ef: elastic modulus of fiber. lf = average fiber length. d f = average fiber diameter, n L = longitudinal efficiency factor and n T = transverse efficiency factor according to Equations 3 and 4, respectively.  Ef   E  − 1  m nL =  Ef   lf   + 2  E  m  df  Ef    −1 E m  nT =   Ef    + 2  Em 

  



This model is useful for predicting the behavior of unidirectionally aligned composite materials and is used to evaluate the reinforcing effect of randomly oriented nanoparticles. It assumes that the dispersion of nanoparticles is uniform. Considering the dimensions obtained for the studied CNCs, they showed the necessary conditions for acting as reinforcement. According to the Halpin-Tsai model, the higher the aspect ratio, the larger the composite modulus. Figure 8 shows the theoretical and experimental results. Although the behavior of the experimental data is similar to the theoretical one, values are lower than those expected for nanocomposites, which shows that the studied system failed in some aspects. To elucidate this behavior, microanalyses are further discussed. In addition to the Halpin-Tsai prediction model being used to calculate the modulus of elasticity, it was also used to estimate the form factor (𝝃). Since the elastic moduli of the components are known (Table 4), it was possible to theoretically evaluate the quality of the dispersion of the CNCs in the epoxy matrix. The form factor (𝝃) was derived from the Halpin-Tsai model in order to be applied to the nanocomposites. This model predicts Ec/Em (where Ec is the composite modulus and Em is the matrix modulus) and the value of 𝝃 for the best fit with the experimental results. The higher the value of 𝝃, the greater the indication of improvement in dispersion quality caused by the tendency of the nanoparticles to agglomerate[46]. Figure 9 and Table 4 show that the dispersion quality is higher for smaller volume fractions. As can be seen, the proximity of the line of the experimental relative modulus is greater for 𝝃 = 10. As the volumetric fractions of the CNCs increase, the quality of the dispersion decreases. The dispersion of nanoparticles is a crucial factor for 8/13

improving the resistance properties. This may explain the mechanical properties. The sample with 4% and 5.5% of CNCs showed better dispersion according to the model, which may justify this sample having greater rupture stress than the samples with larger volumetric fractions of CNCs. Therefore, where there is better dispersion, there is also an increase in rupture stress[45]. Figures 10a-c shows the TEM images of the epoxy/ CNC nanocomposites. The presence of the CNCs can be observed at various scales, as well as their dispersion in the epoxy matrix. The images show aggregations of CNCs in all the nanocomposites, as well as their non-uniform dispersion in the matrix. Ultrasound was applied in order to favor the dispersion of the CNCs in the matrix; however, the dispersion was not uniform and had the formation of aggregations, which may explain the poor performance regarding the mechanical properties in relation to the neat matrix. Uniform and good dispersion of the CNCs in the matrix tends to improve mechanical properties and tends to make the percolation theory applicable[47]. Figure 10d shows an image with a greater magnitude of CNC aggregations in the epoxy matrix, and it is possible to observe its morphology. The lack of a better reinforcing effect could be explained by the hydrophilic nature of CNCs[40], which causes them to Table 4. Values of the tensile modulus of the composites, matrix and CNCs. Vf (4% CNCs)

Vf (5.5% CNCs)

Vf (7% CNCs)

Ec (GPa) 1.94 Ec/Em (ξ = 5) 1.16 Ec/Em (ξ = 10) 1.10 Experimental 1.13

1.77 1.11 1.07 1.03 ECNC (GPa) 50[55]

2.14 1.20 1.14 1.25 Em (GPa) 1.70

Figure 9. Comparison of experimental data with those predicted for the effect of Vf (CNCs) (fraction of volume in percentage) on the relative modulus of epoxy nanocomposites + CNCs. Polímeros, 31(3), e2021034, 2021

Role of cellulose nanocrystals in epoxy-based nanocomposites: mechanical properties, morphology and thermal behavior

Figure 10. TEM images of nanocomposites (a) 4% CNCs, (b) 5.5% CNCs, (c) 7% CNCs, (Magnitude of 40,000 X) and (d) details of aggregation of CNCs (4% CNCs) (Magnitude of 200,000 X).

have a poor interaction with the epoxy matrix[33,42]. Besides the lack of dispersion, the presence of water in the structure of the epoxy matrix can greatly affect the mechanical properties of nanocomposites[48]. The slight reduction in the Tg values of the nanocomposites (5 °C for the 5.5 and 7% CNCs) shows that residues of water molecules on the surfaces of the CNC aggregations may have caused cellulose degradation. The molecules resulting from the degradation increased the free volume of the network structure of the epoxy matrix and, consequently, reduced the mechanical properties of the nanocomposites[28]. The presence of CNCs in the epoxy matrix may also have hindered the crosslinking of the epoxy resin[49], though the conversion was quite similar for all samples, which leads us to disregard this effect. The samples with 7% CNCs showed that, even with the CNC aggregation, it may also have had better dispersion of CNCs and a better interfacial interaction with the matrix, and this led to greater toughness (IS) when compared to Polímeros, 31(3), e2021034, 2021

other nanocomposites. The reduction of impact energy for nanocomposites is explained by the same reason, i.e., the CNC aggregations[42,50]. The SEM images in Figures 11a-h show the effect of the CNCs on the fractured surfaces of the tensile tested samples. Neat epoxy resin (Figure 11a) has a smooth, homogenous and glassy region and semi-elliptical marks and steps[51]. The surface pattern indicates a low resistance to crack propagation that led to brittle failure[52]. The behavior in Figure 11b, marked by a black arrow, revealed its typical brittle nature with poor resistance to cracking or rupturing and its tendency towards propagation[53,54]. Figures 11c-h show the SEM micrographs of the nanocomposites with 4.0, 5.5 and 7.0% CNCs. The incorporation of CNCs caused the roughness in the nanocomposite fractures observed in Figures 11c, e, and g. In the images, the black arrows show that the CNCs increase the roughness, similar to what was observed by Kumar et al.[40]. Based on the images and on 9/13

Nascimento, N. R., Pinheiro, I. F., Alves, G. F., Mei, L. H. I., Macedo Neto, J. C., & Morales, A. R.

Figure 11. SEM micrographs of resin and nanocomposite (a) and (b) neat epoxy; (c) and (d) 4.0% CNCs; (e) and (f) 5.5% CNCs; and (g) and (h) 7.0% CNCs. 10/13

Polímeros, 31(3), e2021034, 2021

Role of cellulose nanocrystals in epoxy-based nanocomposites: mechanical properties, morphology and thermal behavior the mechanical properties, we can state that, during the deformation, the CNCs reinforced the system and increased the elastic modulus, but accelerated the fracture by acting as point defects.

4. Conclusions The CNCs obtained from microcrystalline cellulose using the acid hydrolysis method presented a high degree of crystallinity, with around 82.7% of the majority being cellulose I with dimensions of length and diameter of (142 ± 36) nm and (11 ± 2.5) nm, respectively, which resulted in an aspect ratio of 12.5. This fibrous morphology led to the calculation of the volume fraction percolation threshold of 5.5%. Analysis of the epoxy/CNC nanocomposites showed that the addition of CNCs caused minor changes in the cure conversion and in the thermal properties. The mechanical tensile tests showed an increase in elastic modulus, but the expected abrupt increase over the percolation threshold was not observed. The Halpin-Tsai modulus of elasticity prediction model was applied and showed coherence in the curve tendency, though the experimental results were lower than the prediction, which was attributed to the lack of dispersion of the CNCs and the weak interaction between the phases. Impact resistance and the flexural modulus showed decreases when compared to pure resin. The lack of a better reinforcing effect may be explained by the hydrophilic nature of CNCs, which cause them to have a poor interaction with the epoxy matrix and/or by the lyophilized CNC aggregation, which shows that the applied dispersion process was not capable of uniformly dispersing the CNCs in the matrix.

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ISSN 1678-5169 (Online)


Structural and optical properties o plasma-deposited a-C:H:Si:O:N filmsa Juliana Feletto Silveira Costa Lopes1 , Jean Tardelli1 , Elidiane Cipriano Rangel1  and Steven Frederick Durrant1*  Laboratório de Plasmas Tecnológicos, Instituto de Ciência e Tecnologia de Sorocaba, Universidade Estadual Paulista – UNESP, Sorocaba, SP, Brasil



Abstract Thin a-C:H:Si:O:N films were deposited from plasmas fed hexamethyldisiloxane, oxygen and nitrogen, and characterized as a function of the partial pressure of oxygen in the feed, Rox. Deposition rates varied from 10 to 27 nm min-1. Surface roughness was independent of Rox, being around 10 nm. The films contain C=C and C=O, and also Si-C and Si-O-Si groups. Lower [C] and [N] but greater [O] and [Si] were measured in the films as Rox was increased. Refractive indices of ~ 1.5 and optical energy gaps which fell from ~ 3.3 to ~2.3 eV were observed with increasing Rox. The Urbach energy fell with increasing optical gap, which is characteristic of amorphous materials. Such materials have potential as transparent barrier coatings. Keywords: plasma enhanced chemical vapor deposition; optical band gap; urbach energy. How to cite: Lopes, J. F. S. C., Tardelli, J., Rangel, E. C., & Durrant, S. F. (2021). Structural and optical properties of plasma-deposited a-C:H:Si:O:N films. Polímeros: Ciência e Tecnologia, 31(3), e2021035. https://doi.org/10.1590/01041428.210043.

1. Introduction Thin films of a-C:H:Si:N[1-8], a-C:H:Si:O[9-20], and a-C:H:Si:O:N[21-24], where a designates amorphous, have been produced by plasma deposition from diverse monomers and comonomers. For example, a-C:H:Si:N has been produced, among others, in plasmas fed hexamethyldisilazane and nitrogen[1], diethylsilane and ammonia[3], and methane, silane and nitrogen[5]. Similarly, a-C:H:Si:O has been produced from plasmas of hexamethyldisiloxane (HMDSO)[9], trimethylmethoxysilane and argon[10], and HMDSO and argon[12]. Films of a-C:H:Si:O:N have been deposited from plasmas fed HMDSO, acetylene and nitrogen[22], and from tetramethlydisilazane, oxygen and nitrogen[23]. Applications of a-C:H:Si:N include electrical insulation, protective layers[1], and electronic passivation of n-type c-silicon[25]. Amorphous C:H:Si:O films show promise as hydrophobic and corrosion protection layers[13,18]. In the plasma enhanced chemical vapor deposition (PECVD) of HMDSO/O2 mixtures film precursors of mass 148, such as Me3-Si-O-Si=O-Me, where Me represents methyl, and related Si-O, Si-OH and C=O have been detected[16]. Species present in the plasma phase include H, CH and CO[26]. At high proportions of O2 in the feed, O, CO, OH and H are observed, along with CO+ and CO2+[27]. Another known feature of HMDSO-O2 plasmas is the production

of SiOx powder[28]. There are few studies, however, dealing directly with a-C:H:Si:O:N films. An exception is the study of films produced from HMDSO-N2 mixtures at atmospheric pressure in a dielectric barrier discharge[29]. At high values of [N2]/([Ar + N2]), Si-N groups were detected in the films. Amine and silazane moieties were present. The present study focuses on plasma-deposited a-C:H:Si:O:N produced from HMDSO with O2 and N2. There are few extant studies of such material. The following characterizations of the films allow a composite picture of key properties for possible applications to be mapped. Therefore, morphological, chemical, structural and compositional features, together with optical properties were obtained. Surface roughness and morphology were studied using profilometry and scanning electron microscopy (SEM) respectively. Surface contact angles were measured using goniometry. Chemical structure and composition were probed using infrared reflection-absorption spectroscopy (IRRAS) and energy dispersive X-ray spectroscopy (EDS). The refractive index, optical energy gap, and Urbach energy of the films were calculated from film thickness and ultraviolet-visible near infrared spectroscopy (UVS) data. Subsequently, this suite of properties was used to suggest possible applications.

This paper has been partially presented at the 3rd International Conference on Materials Sciences and Nanomaterials, ICSMN 2019, Oxford University, Oxford, UK, 22nd to 24th July 2019.


Polímeros, 31(3), e2021035, 2021



Lopes J. F. S. C., Tardelli J., Rangel E. C., & Durrant S. F.

2. Materials and Methods The deposition system is shown in Figure 1. Films were deposited in a cylindrical stainless-steel chamber, containing two circular, horizontal, parallel-plate electrodes. Radiofrequency (13.56 MHz) power was fed from a supply (Tokyo Hy-Power, MB-300) at 70 W, via a matching network used to minimize the reflected power, to the lower electrode, while the upper electrode was grounded. Pressure was monitored using a Pirani gauge (Agilent, PCG-750). The base pressure was ~ 0.7 Pa. Nitrogen and oxygen gases (99.95% pure, White Martins, Brazil) were admitted to the chamber from cylinders via precision leak valves (Edwards, LV10-K). For HMDSO, vapor was introduced via a leak valve from a vial of the liquid reagent. For depositions, the chamber was evacuated continuously using a rotary vane pump (Edwards, E2M18). A HMDSO partial pressure of 8.0 Pa was maintained for all depositions. The stoichiometry of the films was altered by varying the partial pressures of oxygen and nitrogen, while maintaining a total pressure of these two gases at 8.0 Pa. All depositions lasted 30 min. Substrates, placed on the lower electrode, were of glass, polished stainless steel, and quartz. The glass substrates were used for films examined by profilometry, and goniometry. Polished stainless steel substrates were used for films examined by IRRAS, SEM and EDS. Quartz substrates were employed for films examined by UVS. Film deposition rates were calculated by dividing the film thickness, h, by the deposition time, t. Film thickness was obtained from a film deposited on a glass slide partially covered by an adhesive tape. Subsequent removal of the tape produced a well-defined film edge, which was measured using profilometry (Veeco DekTak 150). Surface roughness, Ra, was calculated using the arithmetic method from film surface profiles, where Ra is the average of the sum of the absolute values of the deviations from the mean height[30]. Surface morphology was examined using SEM (Jeol JSM-6010LA) and elemental analysis carried out using an EDS accessory (Dry SD Hyper (EX 94410T1L11)). Surface contact angle measurements were made with drops of distilled deionized water using a goniometer (Krüss DSA25E). Ten measurements were made on each of three drops placed at different positions on the film. Infrared reflection absorption spectroscopy (IRRAS) was used to accumulate spectra with a Jasco FTIR-410 instrument over the wavenumber range from 400 to 4000 cm-1. Each spectrum was obtained using 128 scans at a resolution of 4 cm-1. A Perkin Elmer Lambda 750 ultraviolet-visible near infrared spectrometer was used to collect spectra in the wavelength interval of 300 to 2500 nm. For most of the spectra it was possible to determine the refractive index, n, of the film using a method given by Cisneros et al.[31]. When interference maxima and minima are present in a transmittance spectrum, each extremum has an associated integer, m, given by:

m = 4nh / λm (2)

n may be calculated. The absorption coefficient, α(E), can also be calculated as a function of the photon energy (E). When Tauc plots exhibit a linear region, extrapolation to the x-axis (which shows the variable E) reveals the optical energy gap[32], which we designate as the Tauc gap.

3. Results and Discussion Figure 2 shows the film deposition rate as a function of Rox. Oxygen is very reactive and fragmentation of HMDSO increases rapidly as Rox is increased, roughly doubling the deposition rate for intermediate values of Rox. Film deposition, however, is reduced by etching by atomic O. Thus, for Rox of 4.0 Pa the deposition rate is decreased to close to its value in the absence of oxygen in the feed. For greater Rox, etching exceeds film growth and no film is produced. Groups such as CHx (x = 1 to 3) and silicon-containing molecular fragments are film precursors. As proposed by Balu et al.[33] hydrocarbon groups may also be etched via reactions of the form: R + O → R • + OH


R + OH → R • + H 2O (4) R • + O → RO´+CO + CO2 (5)

where R is the polymer backbone and O represents radical or excited oxygen produced in the plasma. The volatile species are lost to the pumping system. In addition, to such species, ions also play a role in film deposition. For example, ions such as O2+ may create surface active sites, remove reaction products held on the film surface, supply energy to drive surface reactions or remove material by direct reactive ion etching[34]. Oxygen ions, such as O2+ may be produced in the plasma by reactions[34] such as those given in Equations 6 to 9.

m ≈ λm −1 / ( λm −1 – λm ) (1)

Employing this relation and 2/8

Figure 1. Schematic of the PECVD system. Polímeros, 31(3), e2021035, 2021

Structural and optical properties o plasma-deposited a-C:H:Si:O:N films e + O2 → O2* + e (6)

2 ( n + 1) C + O2+ + nO2 → 2 ( n + 1) CO (10)

→ 2O + e (7)

2 ( m + 1) C + O2+ + mO2 → ( m + 1) CO2 (11)

→ O* + O + e (8) → O2+ + 2e (9)

Subsequently, carbon may be etched, for example, via reactions[34] indicated in Equations 10 and 11.

Figure 2 also shows the film surface roughness as a function of Rox. While no systematic variation is observed, the absolute values are all small, ~10 nm, and the error bars of the measurements overlap. Thus, under the conditions used, etching does not significantly alter the surface roughness. Scanning electron micrographs of the surface of the films grown at different Rox are shown in Figure 3.

A few, roughly spherical features, of diameter up to ~1 µm are seen in the micrographs of the films deposited at Rox of 0.8 to 2.4 Pa. In cold plasmas containing HMDSO particles often form[26], and when sufficiently large become negatively charged, which keeps them suspended. As the plasma is rich in hydrogen, the particles heat-up owing to surface recombination, which is exothermic. This, in turn, increases the oxidation rate of organic matter. Water contact angle, θ, is shown as a function of Rox in Figure 4. There is a fall in θ from ~90 º for the film deposited without oxygen in the feed to ~73º for the film produced at an oxygen partial pressure of 3.2 Pa. Increased film oxygen content may be responsible for this.

Figure 2. Deposition rate (mean and standard deviation of 6 measurements) and roughness (mean and standard deviation of 6 measurements) as a function of Rox. Films were deposited on glass substrates.

Figure 5 presents IRRAS spectra for the films deposited at different Rox. A peak centered at ~3400 cm-1 shifts to ~3600 cm-1 as Rox increases. This absorption is attributed to hydroxyl groups and, at higher Rox, specifically to OH ν in free SiOH. Absorptions at 2960 and 2900 cm-1 are seen in the spectra of the films deposited at low Rox. These are

Figure 3. Scanning electron micrographs of the surface of the films deposited at (a) Rox = 0 Pa; (b) Rox = 0.8 Pa; (c) Rox = 1.6 Pa; (d) Rox = 2.4 Pa; (e) Rox = 3.2 Pa; (f) Rox = 4.0 Pa. Films were deposited onto glass substrates. Polímeros, 31(3), e2021035, 2021


Lopes J. F. S. C., Tardelli J., Rangel E. C., & Durrant S. F.

Figure 4. Water surface contact angles (mean and standard deviation of 30 measurements) as a function of Rox. Films were deposited on glass substrates.

Figure 5. IRRAS spectra of films deposited at different values of Rox. Film thicknesses were as follows: Rox = 0 Pa (286 nm); Rox = 0.8 Pa (805 nm); Rox = 1.6 Pa (652 nm); Rox = 2.4 Pa (667 nm); Rox = 3.2 Pa (617 nm); Rox = 4.0 Pa (307 nm). Films were deposited onto polished stainless-steel substrates.

attributed to C-H stretching in CH3 and CH2, respectively. Each spectrum exhibits an absorption at ~840 cm-1 caused by –CH3 ρ in Si(CH3)3[19,35,36].

Although oxygen is present in the HMDSO molecule, as residual gas in the deposition chamber, and it is also known that post-deposition reactions may occur between free radicals in plasma polymers and ambient oxygen and water[37,38], no significant absorptions attributable to hydroxyl groups are present in the spectrum of the film deposited at zero Rox. For all non-zero Rox, however, hydroxyl groups are present in the film. This suggests that the incorporation of hydroxyl groups is mainly dependent on oxygen deliberately introduced into the chamber feed. Silicon is not bonded to hydrogen in HMDSO but an absorption at ~2100 cm-1, attributed to SiH, is present in the spectrum of the film grown at zero Rox. Overlapping absorptions occur near 1000 cm-1, are attributed to Si-O-Si 4/8

and Si-O-C. Ethylene and methylene groups are expected to be produced by moderate fragmentation of the HMDSO molecule. The Si-O-Si structure is also found in the monomer. No Si-O-C structures are present in the monomer. A small shoulder is apparent at ~1160 cm-1 in the spectra of films grown at greater Rox, and this is attributed to Si-CH2-CH2Si and is indicative of a loss of hydrogen in relation to the HMDSO molecule. As nitrogen concentrations in the films are only a few at.% (as discussed below in relation to the EDS data) there are no well-defined absorptions caused by nitrogen-containing groups. A small band, whose center lies in the 1630 to 1700 cm-1 region, depending on Rox, may be caused by stretching in C=C and C=O, respectively. Neither of these structures (nor SiH) is present in the monomer molecule, indicating that multiple-step reactions occur in their formation. Figure 6 shows the relative concentration, calculated using the method of Lanford and Rand[39], of Si-O bonded to Si or C, and of OH as a function of Rox. As Si-O-Si is the central structure of the HMDSO molecule, it is retained at high concentrations despite considerable variation in [O] and [C] as discussed below in relation to EDS analyses. The steep rise in [OH] with increasing Rox clearly indicates the strong incorporation of oxygen as a hydroxide. As also shown in Figure 6, [Si(CH)3] and [CH] fall with increasing Rox and rise only for Rox > 3.2 Pa. This reflects the loss of hydrogen and carbon with increasing Rox. Figure 7 shows the concentration (at.%) of the elements C, O, Si and N as determined by EDS as a function of Rox (the hydrogen content cannot be measured and has been ignored). In the HMDSO molecule the number O, Si and C atoms are, 11%, 22% and 67%, respectively. Compared to this, [O] and [Si] increase, while [C] decreases, in the film deposited without oxygen in the feed. The concentration of carbon, [C], falls with increasing Rox, except beyond a Rox of 3.2 Pa, where it rises. Both [O] and [Si] rise with increasing Rox, except beyond a Rox of 3.2 Pa, where the concentrations of both elements fall. These decreases in [O] and [Si] at high Rox may be caused by synergistic effects between O and O2+, producing etching and sputtering. For Rox in the interval from 0.8 to 3.2 Pa, θ falls and [O] rises (Figure 4 and Figure 7). For Rox greater than 3.2 Pa, θ rises and [O] falls. Thus, greater values of [O] are associated with smaller contact angles. Inspection of Figure 7 also shows that [N] falls from ~6 at.% for zero Rox to ~0 at.% when Rox is increased to 4.0 Pa. This reflects the diminishing supply of nitrogen in the plasma feed as Rox is increased. Figure 8 shows the transmittance spectra of the films grown at different Rox in the ultraviolet-visible near infrared region. Multiple interference extrema are observed for wavelengths above about 500 nm. The inset of Figure 8 shows the refractive index of the deposited material as a function of Rox. Typical values of n are ~1.5, lower than those found for films deposited from diethylsilane, ammonia and nitrogen, which are typically 1.7 to 1.8, depending on the deposition temperature[2]. The films of the present study, however, differ in having low nitrogen contents. As the refractive index depends on the effective polarizability, the molar mass and the density[40], the relative constancy of n implies Polímeros, 31(3), e2021035, 2021

Structural and optical properties o plasma-deposited a-C:H:Si:O:N films that the effects in any changes in these parameters tend to cancel out. Values of n close to 1.5 were also found by Mota et al.[19] for a-C:H:Si:O films deposited from HMDSO. Similarly, the values reported here are consistent with those of 1.55 reported by Amri et al.[24] for films deposited from HMDSO and nitrogen at low power (20 W).

Figure 6. Relative concentration of Si-O-Si, Si-O-C, Si-CH3, CH, and OH as a function of Rox. These values were calculated from the area of the relevant infrared absorption divided by the film thickness, and were normalized to the maximum obtained for each species.

Figure 7. Atomic concentration of C, O, Si and N in the films determined by EDS as a function of Rox.

Figure 8. Transmittance spectra for films deposited at different Rox. Films were deposited onto quartz substrates. Inset shows the refractive index of the deposited material as a function of Rox. Polímeros, 31(3), e2021035, 2021

Tauc plots of the transmission spectra of the films, shown in Figure 9, allow the determination of the optical gap. The optical gap, shown as a function of Rox in Figure 10, falls from ~ 3.3 to ~2.3 eV as Rox is increased. This interval overlaps with that of 1.87 and 2.7 eV reported by Swatowska[8] for a-C:H:Si:N films deposited from CH4, SiH4 e NH3. Greater nitrogen or carbon contents tend to increase the gap[8]. Thus, the decline in [C] and [N] with increasing Rox may account, at least in part, for the observed fall in the gap. The decline also depends, however, on the densities and types of dangling, single and multiple bonds. The (unknown) hydrogen content probably accompanies the fall in [C]. Thus, the density, for example, of Si-O probably increases while that of Si-C decreases. As bond lengths of Si-O and Si-C are about 163 pm[41] and 185 pm[42], respectively, such

Figure 9. Tauc plots for films deposited at different Rox. Inset shows the variation of ln(α(ν)) vs. the photon energy.

Figure 10. Optical gap of the deposited material as a function of Rox. Inset shows the Urbach energy as a function of the optical gap. 5/8

Lopes J. F. S. C., Tardelli J., Rangel E. C., & Durrant S. F. changes are expected to modify the optical transmittance and hence the gap. A fall in the band gap of a-C:H:Si:O films at greater applied powers has been attributed to a greater density of dangling bonds, which are responsible for the formation of localized defects and hence of localized states in the band structure[43]. The reduction in the gap with increasing Rox is also consistent with the finding that greater Si-O-Si and Si-O-C concentrations in HMDSO plasma films are associated with lower gaps[19]. The inset of Figure 9 shows plots of ln(α) v E, for each film, from which Urbach energy, EU, was calculated. The resulting EU values are shown as a function of the optical gap in Figure 10 (inset). A linear fall is observed, which is consistent with the behavior of other amorphous materials[44]. Greater values of EU (found at higher Rox) are associated with disordering and defect states and thus with a reduction in the optical gap.

4. Conclusions Smooth, hydrophobic a-C:H:Si:O:N thin films may be produced from cold plasmas fed mixtures of HMDSO-O2-N2. The films have a complex structure, containing a network with Si-O-Si and C-C chains, similar to that reported by Hilbert et al.[45] but with greater Si and O contents. The oxygen content generally rises, and the carbon content falls with increasing Rox. Silicon contents are around 20 at.% and nitrogen contents a few at.%. At low Rox there are relatively high concentrations of Si-CH3 groups, which accounts, at least in part, for the hydrophobicity of these films. As Rox increases, the incorporation of oxygen and complex structural changes, cause a decrease in the optical gap. Lower Urbach energies are associated with greater optical gaps. Owing to the combination of hydrophobicity, smoothness, good adhesion to diverse substrates, and high optical transparency, films grown at low Rox may find application as transparent barrier coatings. Similar films have been investigated as interlayers for SiOx diffusion barrier coatings on polypropylene, but in this case plasma-polymerized HMDSO offered the advantage of greater deposition rates [22]. Films deposited from HMDSO-O2-Ar mixtures may find application as multilayered organosilicon/silica films for the protection of metal surfaces[20]. Although there is a report of plasma-deposited nanocomposite hydrogenated silicon oxycarbonitride films obtained from tetramethlydisilazaneO2-N2 mixtures[23], characterizations beyond chemical composition and structure are required to suggest possible applications of this material. The mechanical properties of our films, such as hardness, Young’s modulus and stiffness, are also relevant to possible applications, but require separate study.

5. Acknowledgements The authors thank Fundação de Amparo à Pesquisa do Estado de São Paulo (2017/15853-0) and Conselho Nacional de Desenvolvimento Científico e Tecnológico for financial support. This study was also financed in part by Coordenação de Aperfeiçoamento de Pessoal de Nível 6/8

Superior – Brasil, CAPES - Finance code 001. We thank Prof. José H. Dias da Silva for helpful discussion of the optical data and Jamille Altheman for technical assistance with the SEM/EDS measurements.

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44. Grenadyorov, A. S., Oskomov, K. V., & Solovyev, A. A. (2018). Effect of deposition conditions on optical properties of a-C:H:SiOx films prepared by plasma-assisted chemical vapor deposition method. Optik (Stuttgart), 172, 107-116. http://dx.doi.org/10.1016/j.ijleo.2018.07.024. 45. Hilbert, J., Mangolini, F., McClimon, J. B., Lukes, J. R., & Carpick, R. W. (2018). Si doping enhances the thermal stability of diamond-like carbon through reductions in carbon-carbon bond length disorder. Carbon, 131, 72-78. http://dx.doi. org/10.1016/j.carbon.2018.01.081. Received: June 01, 2021 Revised: Oct. 06, 2021 Accepted: Nov. 29, 2021

Polímeros, 31(3), e2021035, 2021

ISSN 1678-5169 (Online)


Incorporation of astrocaryum vulgare (tucuma) oil into PCL electrospun fibers Nathan Rampelotto Bressa1 , Vinícius Rodrigues Oviedo2 , Aline Machado Bessow Machado1 , Willians Lopes de Almeida3 , Tiago Moreno Volkmer4 , Luis Alberto Loureiro dos Santos5 , Michele Rorato Sagrillo2  and Luiz Fernando Rodrigues Junior1*  Departamento de Engenharia Biomédica, Universidade Franciscana – UFN, Santa Maria, RS, Brasil 2 Programa de Pós-graduação em Nanociências, Universidade Franciscana – UFN, Santa Maria, RS, Brasil 3 Instituto Federal de Educação, Ciência e Tecnologia do Amapá, Macapá, AP, Brasil 4 Departamento de Engenaharia de Materiais, Universidade Federal de Pelotas – UFPel, Pelotas, RS, Brasil 5 Laboratório de Biomateriais e Cerâmicas Avançadas – LABIOMAT, Departamento de Engenharia de Materiais, Universidade Federal do Rio Grande do Sul – UFRGS, Porto Alegre, RS, Brasil



Abstract The aim of this study was to incorporate tucuma oil (Astrocaryum vulgare) into PolyCaprolactone (PCL) electrospun fibers and evaluate its physicochemical properties and cell viability. FTIR and DRX confirmed that tucuma oil (TO) does not affect the chemical properties of PCL and that the oil was loaded into the PCL microstructure, while TGA analysis showed that the oil increased the thermal stability of the polymeric fibers. SEM showed that the addition of the oil modified fibers structure by reducing the average fiber size from 5.5 μm to 1.7 μm for TO loaded samples. Cell viability assay demonstrated an increment on cell proliferation from 80% of pure PCL to 100% for samples containing TO. Therefore, it can be concluded that tucuma oil can be incorporated into PCL to form fibers by electrospinning, without meaningful changes in its physicochemical properties and increasing its biocompatibility. Keywords: cytotoxicity, vegetal oil, fibers. How to cite: Bressa, N. R., Oviedo, V. R., Machado, A. M. B., Almeida, W. L., Volkmer, T. M., Santos, L. A. L., Sagrillo, M. R., & Rodrigues Junior, L. F. (2021). Incorporation of astrocaryum vulgare (tucuma) oil into PCL electrospun fibers. Polímeros: Ciência e Tecnologia, 31(3), e2021036. https://doi.org/10.1590/0104-1428.20210056

1. Introduction There are several studies focusing on the manufacture of polymeric fibers and one of the most used processes is electrospinning[1-8]. This technique has such as being able to produce scaffolds with a controlled fiber diameter, high surface area, and porous structure[1,9-12]. The ability to reproduce and manipulate the electrospinning process in vitro on a spatiotemporal scale similar to that of native tissue provides a great potential of clinical success[6,13,14]. The wide range of commercially available biomaterials, as well as the strategies adopted in tissue engineering and regenerative medicine, favor the search for new products and methodologies to obtain these[15-17]. In line with this approach, the manufacture of scaffolds from absorbable, hydrolytically degradable polymers belonging to the aliphatic polyester class is being widely investigated for the use in tissue engineering[18-24]. Their inherent biocompatibility properties and the possibility of undergoing hydrolysis in the body make these biomaterials suitable for the tissue reconstruction process. In addition to

Polímeros, 31(3), e2021036, 2021

these, one can list its ease of processing and modulation of degradation rate, mechanical and visco-elastic properties[25]. Electrospun fibers are extremely attractive in the biomaterials field due to their large surface area to volume ratio[26]. Its potential applications include tissue engineering scaffolds, drug delivery media, wound healing, filtration media, composites[8,27,28], among others. It is known that the properties and internal molecular structure of polymers are strongly affected by their processing conditions. Thus, understanding the processing - structure - property relationship is of great importance for the development of polymeric fibers that meet the demands of the desired application[29]. Several studies have been seeking viable alternatives for the use of absorbable polymers containing natural oils as those oils may enhance the material properties without significantly alter the structure of the polymeric matrix. In this idea, the tucuma oil shows great application possibilities, due to the small toxicity presented and efficiency in the controlled drug delivery[30].



Bressa N. R., Oviedo V. R., Machado A M. B., Almeida W. L., Volkmer T. M., Santos L. A. L., Sagrillo M. R. & Rodrigues Junior, L. F. The tucuma palm (Astrocaryum vulgare), which is found in the Amazon Rain forest, is considered a pioneer of expressive growth, fire resistant with ability to sprout after burning and mainly inhabits the poultry and pastures. The kernel of the palm tree is externally covered with an oily orange canopy from which the oil is extracted[31]. Among tropical seeds, tucuma palm is an economical source of vegetal oils[32] and it is also abundant in the northwest, north, and central-west regions of Brazil[33]. Besides that, the simplicity of the extraction process to obtain the oil, which is mainly based on mechanical cold-pressing[34], also turns it into a promising biotechnological resource. In this view, tucuma oil has been already used in the biomedical field due to the high fatty acids content and the fact it has several carotenoids as bioactive compounds[35]. Based on the aforementioned, its incorporation into electrospun biomaterials trends to improve biocompatibility in addition to physical-chemical properties[36-40]. In this work, a study for obtaining PCL fibers with the addition of tucuma oil by the electrospinning process was carried out. The proposal aimed to evaluate whether the tucuma could be electrospun together with the PCL and whether this product would present good biocompatibility for its application as scaffolds or dressings for pressure injuries. Different amounts of tucuma oil were added to the PCL and its impact on morphology, physical and chemical properties, and its cytotoxic effect in cell media were studied.

second. The search was done at the International Center for Diffraction Data (ICDD) Database, file PDF2014 PDF-2 Release 2014 RDB.

2. Materials and Methods

Blood Collection for toxicological tests: Peripheral blood samples were obtained from three discard samples from the Clinical Analysis Laboratory of the Franciscan University, under the approval of the Institution’s Human Ethics Committee (CAAE: 31211214.4.0000.5306) with no identification data. Samples were obtained by venipuncture using Vacutainer®-type heparin tubes, which were used to separate the Peripheral Blood Mononuclear Cell (PBMCs).

2.1 Electrospinning methodology For the preparation of a homogeneous polymer solution, 15% (w/v) PCL was dissolved in acetone at 50 °C using a hot plate and it was magnetic stirred until the polymer was completely solubilized. This solution was transferred to a 5 mL syringe with an 18 gauge needle. After some tests, the experimental parameters were optimized to 10 kV of potential difference and 10 cm distance between the tip of the needle containing the polymeric solution and the target. The tucuma oil was added into polymeric PCL solution before the electrospun process on the concentrations describe in Table 1.

2.2 Characterization 2.2.1 X-ray diffraction (XRD) The characterization of the crystalline planes of the samples was performed using a Bruker D2 PHASER diffractometer. Couple Two Theta / Theta scan with 1482 steps of 1 second, 2θ from 5° to 70° and increment of 0.05° per Table 1. Parameters used in the electrospinning technique. Sample PCL PCLT100 PCLT250 PCLT500


Oil Electric Concentration Distance (cm) potential (kV) (µg/mL) 0 10 10 100 10 10 250 10 10 500 10 10

2.2.2 Fourier-transform infrared spectroscopy (FTIR) The FTIR technique was used to verify the presence of functional groups in the PCL matrix. The assay was performed on a Spectrometer FTIR/NIR, model FRONTIER, brand PERKINELMER, with ATR methodology and resolution of 8 cm-1 and scanning from 4000 to 650 cm-1. 2.2.3 Thermogravimetric analysis (TGA) Thermal behavior of samples was evaluated under N2 atmosphere with flow rate of 20 mL/min and heating rate of 10 °C/min. The testing was performed from 30 to 600 °C using one Perkin Elmer instrument, model TGA 400. 2.2.4 Scanning electron microscopy (SEM) The morphology and size of the fibers were characterized by scanning electron microscopy (SEM), using a Hitachi, model tm3000 (secondary electron – SE). The software ImageJ was used to measure the diameter of the fibers. Fifty measurements were performed per sample 2.2.5 Cell viability (MTT) The MTT assay for cell viability was used to verify the cytotoxicity of the obtained scaffolds by using an experimental protocol similar to that described by Wilms et al.[41].

Treatments: A culture medium containing only the cells was used as a negative control, while the PCL, PCL100, PCL250, and PCL500 samples were incubated with PBMCs in an environment with 5% CO2 at 37 °C for 24h. Moreover, the results were represented as a function of optical density and then, the cell viability was calculated through spectrophotometry at the 570 nm wavelength, according to Equation 1: CellViability ( % ) =

OD570e x100 (1) OD570b

Where: OD570e: mean value for optical density of 100% of the extract of the test samples. OD570b: mean value for optical density of the blanks.

2.3 Statistical analysis For the analysis of the SEM images, the IMAGE J software was used, where 50 measurements were taken from each group. Subsequently, the data were entered at the software OriginPro 8.1. The one-way ANOVA test was used to see if there was a significant difference between the measurements where p would have to be less than 0.05. Polímeros, 31(3), e2021036, 2021

Incorporation of astrocaryum vulgare (tucuma) oil into PCL electrospun fibers Also, for the MTT assay for cell viability, the treatments were compared to the negative control through one-way ANOVA to check for statistical differences between the groups and Dunnett’s post-hoc test to compare their means. The 95% confidence interval was used and p<0.05 were considered significant.

3. Results and Discussions Figure 1 shows the diffractograms selected for PCL and the PCL mixed with tucuma oil at 250 µg/mL and it shows it has the same structural characteristics. When the PCL/tucuma is analyzed at different concentrations, a crystal structure is obtained[30]. Diffractogram peaks present between 2Ɵ ranging from 21° to 24° are typical of PCL. The FTIR-ATR spectrum of tucuma oil, PCL and PCL + tucuma oil are shown in Figure 2 and it contains characteristics

Figure 1. Diffractograms of pure PCL and PCL samples with 250 µg/mL of tucuma oil.

bands of a vegetable oil[42-45]. These bands are described in Table 2, according to the findings of Leonardi et al.[42], Ali et al.[44] and Gomez et al.[45]. Bands in the region between 2980 and 2830 cm-1 are characteristic of the asymmetric and symmetric stretching modes of the C-H methylene group. This molecule is present both in PCL and in vegetable oils, such as tucuma. The carbonyl C=O aliphatic stretching was observed in 1733 cm-1 and 1744 cm-1 to PCL and PCl + tucama oil, respectively. It was observed that the incorporation of tucuma oil in PCL resulted in an increase in the intensity of the bands related to the methylene group, since all samples of PCL + tucuma presented greater intensity in these bands when compared to pure PCL. Nevertheless, it is not possible to see a significant difference in these intensities when the PCL + tucuma oil spectra are compared between them. Thus, the FTIR results suggest that the quantity of 100 µg/mL is sufficient to promote alterations in PCL properties. Figure 3 shows the SEM images from PCL and PCL with tucuma oil. The medium diameter from PCL fibers was the 5.3 ± 3.7 µm and the PCL fibers with tucuma oil were 1.72 ± 0.90 µm, 2.6 ± 1.3 µm, 2.9 ± 1.2 µm, to samples with 100 µg/mL, 250 µg/mL and 500 µg/mL, respectively. The ANOVA test showed a significant difference among PCL fibers diameter and all samples with tucuma oil. The same result was observed between and PCL plus 100 µg/mL and PCL plus 500 µg/mL (p < 0.05), despite of difference between fibers diameter from PCL plus 100 µg/mL and PCL plus 250 µg/mL the ANOVA test did not find a significant difference when comparing them (p = 0.29). These results were different from the work from Felgueiras et al.[46], where there was not statistical change in fiber diameters when an essential oil was added to polymer fibers. However, them are in accord with the showed by Tampau et al.[47] and Hasanpour Ardekani-Zadeh and Hosseini[48], where the addition of oil furthered the diameter fibers reduction, when compared to polymer without oil. This comportment could be explained by change of physical-chemical characteristics

Table 2. ATR-FTIR spectrum analysis for all samples[35,37,38]. Wave number (cm-1) 3006 2922 2853 1743 1710 1650 1464 1417

Figure 2. Infrared spectroscopy of PCL and PCL samples with 100 µg/mL, 250 µg/mL and 500 µg/mL of tucuman oil. Polímeros, 31(3), e2021036, 2021

1377 1239 1161 1117 1095 721

Attributed band


=CH Stretching CH2 Assymetric stretching CH2 Symetric stretching C=O Aliphatic strecthing C=O Stretching C=C Stretching CH2 scissoring deformation In-plane deformation of C-O-H Symetric flexion of CH C-O Stretching C-O (esther) Stretching

Leonardi et al.[42] Leonardi et al.[42] Leonardi et al. [42] Leonardi et al.[42] Leonardi et al.[42] Gomez et al.[45] Leonardi et al.[42] Leonardi et al.[42]

CH2 Rocking deformation

Leonardi et al.[42] Leonardi et al.[42] Leonardi et al. [42], Ali et al.[44] Leonardi et al. [42], Ali et al.[44]


Bressa N. R., Oviedo V. R., Machado A M. B., Almeida W. L., Volkmer T. M., Santos L. A. L., Sagrillo M. R. & Rodrigues Junior, L. F.

Figure 3. SEM of samples with different concentrations of oil and pure PCL, (A) 100 µg/mL, (B) 250 µg/mL, (C) 500 µg/mL and (D) PCL.

of the polymer solution by tucuma oil that modifying the chain entanglements, responsible for fiber formation[49]. Figure 4 shows the SEM image to PCL plus 500 µg/mL tucuma oil where can be see a region in the sample that polymeric fibers collapsed. This could be justified by the insufficient evaporation of solvent due the large quantity of tucuma oil[49]. This result could explain the increased of diameter fibers from samples loaded with more tucuma oil, where minus tucuma oil has finner fibers and great quantity of tucuma oil has thicker fibers (1.72 ± 0.90 µm to PCL plus 100 µg/mL and 2.9 ± 1.2 µm to PCL plus 500 µg/mL). Thermogravimetric analysis (Figure 5) showed that the tucuma oil incorporation (100-500 µg/mL) in PCL leads to a higher thermal stability. Moreover, the events in the region around 330 °C in Figure 5B, Figure 5C and Figure 5D may be occurring together with another thermal event, resulting in a wide endothermic peak around 423 °C. To verify this, the DTG curve for sample PC500 (Figure 6) was analyzed by least squares fit of gaussian functions and deconvolution revealed two peaks (392 °C and 424 °C). This result is consistent with the other samples (PCL100 and PCL250), however, it is noted that the endothermic peaks were moved to higher temperatures. This is due to the greater amount of tucuma oil added to the PCL. 4/8

‘Figure 4. SEM of PCL500 sample.

The in vitro effect of tucuma-loaded fiber scaffolds on PBMCs donated by volunteers was investigated by MTT (Figure 7). Test results showed no cytotoxicity behavior in any of the evaluated samples. Pure PCL nanofibers showed a cell viability around 95.7% against PBMC cells, however, Polímeros, 31(3), e2021036, 2021

Incorporation of astrocaryum vulgare (tucuma) oil into PCL electrospun fibers

Figure 5. Termogravimetric analysis (A) PCL, (B) PCLT100, (C) PCLT250 and (D) PCLT500.

Figure 6. DTG curve for PCLT500 sample analyzed by least squares fit of gaussian functions. Polímeros, 31(3), e2021036, 2021

Figure 7. Cell viability assay of Peripheral Blood Monocytes (PBMC) Lymphocytes on PCL, 100, 250 and 500 μg/mL Tucuma-loaded PCL nanofibers after 24 hours (p≤0.05). Values are expressed as mean ± S.D. of three parallel measurements. (* - significant difference; NS - nonsignificant difference). 5/8

Bressa N. R., Oviedo V. R., Machado A M. B., Almeida W. L., Volkmer T. M., Santos L. A. L., Sagrillo M. R. & Rodrigues Junior, L. F. this result is significantly smaller when compared to negative control and to all tucuma-loaded PCL samples (p≤0.05). On the other hand, there was no significant difference in cell viability with all tucuma-loaded samples and negative control (p≤0.05). The results exhibited that tucuma oil could induce cell proliferation on PCL fiber mats. In agreement with the results obtained on this paper, Ongaratto et al.[50] studied the cellular viability of different tucuma extract concentrations (5, 10, 50, 100 e 500 μg/mL) on PBMCs. The viability of the cells seeded on the pure PCL fibers mats was significantly lower than that of cells cultured on the control well, this result can be credited to the PCL fibers hydrophobic nature[51]. The incorporation of tucuman oil into PCL fiber mats leaded to an improvement on cell viability, which suggests that it possess in its chemical composition molecules such as the β-carotene[52] that can regulate the expression of genes responsible for cell proliferation and differentiation by controlling ROS production and lipid peroxidation[53-55].

4. Conclusions PCL fibers incorporated with tucuma oil were successfully electrospun and were evaluated with respect to their chemical, physical and morphological properties as well as cytotoxicity. The incorporation of tucuma oil did not affected the PCL chemical structure, which was confirmed by XRD and FTIR. Also, the SEM analysis confirmed the fibrous network of PCL and the addition of tucuma oil into this microstructure. In respect to the morphologic properties, it was noted that the mean diameter of PCL fibers decreased with the addition of tucuma oil. Besides that, the incorporation of tucuma oil into the PCL matrix led to a higher thermal stability compared to pristine PCL. Regarding cytotoxicity, the incorporation of tucuma oil showed enhanced biocompatibility, once it increased the cell viability evaluated by an MTT assay and did not present cytotoxicity against PBMCs. Thus, PCL can be electrospun with tucuma oil to achieve a fibrous biomaterial with increased biocompatibility and interesting physical and morphological properties, without altering the PCL chemical structure.

5. Acknowledgements Universidade Franciscana for providing the main materials and infrastructure required to develop this research. To the Instituto Federal de Educação, Ciência e Tecnologia do Amapá, the Universidade Federal de Pelotas, and the Universidade Federal do Rio Grande do Sul. This study was financed in part by the Coordenação de Aperfeiçoamento de Pessoal de Nível Superior - Brasil (CAPES) and Conselho Nacional de Desenvolvimento Científico e Tecnológico (CNPq).

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Polímeros, 31(3), e2021036, 2021

DESCUBRA o conjunto de instrumentos que conduzem a percepções mais profundas sobre as PROPRIEDADES e ESTRUTURA DO POLÍMERO em cada etapa



Polímeros VOLUME XXXI - Issue III - July/Sep., 2021

30 years of Polímeros Three Decades Sharing Polymer Science

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