Polímeros: Ciência e Tecnologia (Polimeros) 1st. issue, vol. 29, 2019

Polímeros VOLUME XXIX - Issue I - Jan./Mar., 2019

São Paulo 994 St. São Carlos, SP, Brazil, 13560-340 Phone: +55 16 3374-3949 Email: abpol@abpol.com.br 2019

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

P o l í m e r o s - I ss u e I - V o l u m e X X I X - 2 0 1 9 I n d e x e d i n : “ C h e m ic a l A b s t r a c t s ” — “ RA P RA A b s t r a c t s ” — “A l l - R u s s i a n I n s t i t u t e o f S ci e n c e ­T e c h n ic a l I n f o r m a t i o n ” — “ R e d d e R e v i s t a s C i e n t i f ic a s d e A m e r ic a L a t i n a y e l C a r i b e ” — “ L a t i n d e x ” — “ W e b o f S ci e n c e ”

and

Polímeros E d i t o r i a l C o u nci l Antonio Aprigio S. Curvelo (USP/IQSC) - President

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

Members Adhemar C. Ruvolo Filho (UFSCar/DQ) Ailton S. Gomes (UFRJ/IMA) Alain Dufresne (Grenoble INP/Pagora) Antonio Aprigio S. Curvelo (USP/IQSC) Bluma G. Soares (UFRJ/IMA) César Liberato Petzhold (UFRGS/IQ) Cristina T. Andrade (UFRJ/IMA) Edson R. Simielli (Simielli - Soluções em Polímeros) Edvani Curti Muniz (UEM/DQI) Elias Hage Jr. (UFSCar/DEMa) Eloisa B. Mano (UFRJ/IMA) João B. P. Soares (UAlberta/DCME) José Alexandrino de Sousa (UFSCar/DEMa) José António C. Gomes Covas (UMinho/IPC) José Carlos C. S. Pinto (UFRJ/COPPE) Júlio Harada (Harada Hajime Machado Consutoria Ltda) Laura H. de Carvalho (UFCG/DEMa) Luiz Antonio Pessan (UFSCar/DEMa) Luiz Henrique C. Mattoso (EMBRAPA) Marco-Aurelio De Paoli (UNICAMP/IQ) Osvaldo N. Oliveira Jr. (USP/IFSC) Paula Moldenaers (KU Leuven/CIT) Raquel S. Mauler (UFRGS/IQ) Regina Célia R. Nunes (UFRJ/IMA) Richard G. Weiss (GU/DeptChemistry) Rodrigo Lambert Oréfice (UFMG/DEMET) Sadhan C. Jana (UAKRON/DPE) Sebastião V. Canevarolo Jr. (UFSCar/DEMa) Silvio Manrich (UFSCar/DEMa)

A ss o ci at e E d i t o r s Adhemar C. Ruvolo Filho Alain Dufresne Bluma G. Soares César Liberato Petzhold José António C. Gomes Covas José Carlos C. S. Pinto Paula Moldenaers Richard G. Weiss Rodrigo Lambert Oréfice

Sadhan C. Jana

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

www.editoracubo.com.br

“Polímeros” is a publication of the Associação Brasileira de Polímeros São Paulo 994 St. São Carlos, SP, Brazil, 13560-340 Phone: +55 16 3374-3949 emails: abpol@abpol.org.br / revista@abpol.org.br http://www.abpol.org.br Date of publication: March 2019

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Polímeros / Associação Brasileira de Polímeros. vol. 1, nº 1 (1991) -.- São Carlos: ABPol, 1991Quarterly v. 29, nº 1 (Jan./Feb./Mar. 2019) ISSN 0104-1428 ISSN 1678-5169 (electronic version)

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

1. Polímeros. l. Associação Brasileira de Polímeros.

Polímeros, 29(1), 2019

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

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

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

O r i g in a l A r t ic l e Synthesis and characterization of the photoswitchable poly(methyl methacrylate-random-methacrylate spirooxazine) Tam Huu Nguyen, Phuc Huynh Tran, Linh Duy Thai, Thuy Thu Truong, Le-Thu T. Nguyen and Ha Tran Nguyen ....................................... 1-8

Natural ageing of polyaramide fiber from ballistic armor Vitor Hugo Cordeiro Konarzewski , Fernando Ludgero Spiekemann and Ruth Marlene Campomanes Santana............................................ 1-7

Application of ashes as filling in reprocessed polypropylene: thermomechanical properties of composites Lina Marcela Crespo and Carolina Caicedo ................................................................................................................................................... 1-7

Coating of urea granules by in situ polymerization in fluidized bed reactors Bruno Souza Fernandes, José Carlos Pinto, Elaine Christine de Magalhães Cabral-Albuquerque and Rosana Lopes Lima Fialho ......... 1-10

Morphological, thermal and bioactivity evaluation of electrospun PCL/β-TCP fibers for tissue regeneration

Lilian de Siqueira , Fábio Roberto Passador, Anderson Oliveira Lobo and Eliandra de Sousa Trichês......................................................... 1-6

Non-isothermal melt crystallization kinetics of poly(3‑hydroxybutyrate), poly(butylene adipate-co-terephthalate) and its mixture Anna Raffaela Matos Costa , Edson Noryuki Ito, Laura Hecker Cavalho and Eduardo Luís Canedo.......................................................... 1-16

Rheological properties of low-density polyethylene filled with hydrophobic Co(Ni)-Al layered double hydroxides Silvia Jaerger, Andreas Leuteritz, Rilton Alves de Freitas and Fernando Wypych ........................................................................................ 1-13

Alternative use of oily fractions of olive oil Melina Bagni , Dolly Granados and María Reboredo..................................................................................................................................... 1-5

Sulfonated poly(ether ether ketone)/hydroxyapatite membrane as biomaterials: process evaluation Cristiane Agra Pimentel , José William de Lima Souza, Flávia Suzany Ferreira dos Santos, Mayelli Dantas de Sá, Valéria Pereira Ferreira, Gislaine Bezerra de Carvalho Barreto, José Filipe Bacalhau Rodrigues, Wladymyr Jefferson Bacalhau de Sousa, Cláudio Orestes Britto Filho, Francisco Kegenaldo Alves de Sousa and Marcus Vinicius Lia Fook............................................................... 1-8

Predicting LDPE/HDPE blend composition by CARS-PLS regression and confocal Raman spectroscopy Daniel José da Silva and Hélio Wiebeck.......................................................................................................................................................... 1-7

Extraction and characterization of cellulose microfibers from Retama raetam stems Abdelkader Khenblouche , Djamel Bechki, Messaoud Gouamid, Khaled Charradi, Ladjel Segni, Mohamed Hadjadj and Slimane Boughali.............................................................................................................................................................................................. 1-8

Evaluation of the dispersion of carbon nanotubes in an elastomeric polyurethane and fatigue test Felipe Luiz Queiroz Ferreira, Magnovaldo Carvalho Lopes, Ana Paula Mendes Lopes, Rodrigo Lassarote Lavall and Glaura Goulart Silva ....................................................................................................................................................................................... 1-7

Effects of mercerization in the chemical and morphological properties of amazon piassava Viviane Rebelo, Yuri da Silva, Saulo Ferreira, Romildo Toledo Filho and Virginia Giacon ............................................................................ 1-6

Gamma irradiation effects on polycaprolactone/zinc oxide nanocomposite films Marcos Paula, Ivo Diego, Ronaldo Dionisio, Glória Vinhas and Severino Alves ........................................................................................... 1-7

Synthesis and characterization of isoprene oligomers to compare different production chemical processes Renata Vieira Pires, Larissa Mota Barros Pessoa, Monica de Almeida de Sant’Anna, Alexander Fainleib, Rita de Cassia Pessanha Nunes and Elizabete Fernandes Lucas .................................................................................................................... 1-9

Cover: TEM image of Carbon Nanotube (CN).

Arts by Editora Cubo.

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Polímeros, 29(1), 2019

Braskem Picks Polypropylene Distributor for U.S. Market Brazilian-owned petrochemical firm Braskem USA formed a distribution partnership with North American polymer resin distributor and recycler PolyQuest Inc. for its U.S. polypropylene (PP) business, a recent company press release announced.

The biggest distributor of PET in the U.S., Wilmington, N.C.-based PolyQuest serves a variety of customers, including the sheet, film, textiles, strapping, and BCF markets. Braskem said the new relationship allows the company to reach new customers and markets and will complement its efforts to launch a new line of polypropylene. “PolyQuest shares a focus on advancing the future of the circular economy through a commitment to sustainable development. We look forward to leveraging our expertise in polypropylene recycling and closed-loop technologies with PolyQuest’s established U.S. recycling facilities,” said Mark Nikolich, chief executive officer of Braskem America, in a statement. “Our leverage across these activities will only grow in importance as Braskem remains on track to launch Delta, our new U.S. 450 kt polypropylene production line in early 2020.”

Construction is currently underway on the new $675 million Delta PP production line in La Porte, TX. When the project reaches completion in early 2020, the line will have a capacity of 450 kt/yr, or about 1 billion lb/yr, for homopolymers, random copolymers, impact copolymers, and reactor TPOs. PolyQuest operates two recycling facilities in New York and South Carolina that produce post-consumer polypropylene products. The company is working to launch several solutions produced from recycled propylene this year. “There is no better partner for PolyQuest in the polypropylene market than Braskem and we look forward to expanding our relationship responsibly and rapidly,” said Tod Durst, executive vice president of PolyQuest, in the release.

“The expansion of our resin offerings in the U.S. to include polypropylene was a logical one for us as it further enhances our position as a value-added supplier to an important and growing segment of the industry.”

Source: Powder Bulk Solids - www.powderbulksolids.com

Thermoplastic composite tape features in revolutionary change in vehicle panel production Sabic announced a new, cutting-edge technology for producing lightweight, cost-effective and recyclable vehicle panels using its Udmax tape, a unidirectional, fiber-reinforced thermoplastic composite at the recent JEC World 2019 event. This innovative technology, which is designed to replace traditional panels made of metal and thermoset materials for interior and exterior automotive applications, will soon be commercialized in the bulkhead of a light commercial vehicle (LCV) produced in large scale for the global automotive market.

Produced using a highly efficient, one-shot process of lamination and low-pressure molding, the bulkhead featuring

Udmax tape is 35% lighter and complies with ISO 27956 standard for securing cargo in delivery vehicles. The bulkhead was developed through international collaboration among Sabic; RLE International, an engineering services provider headquartered in the United Kingdom; AMA Composites, an Italian toolmaker; and Setex Textil GmbH, a weaver based in Germany.

Vehicle panels made with Udmax tape are said to combine strength and impact resistance with lightweight, which can result in mass reduction of interior panels of up to 35 percent in comparison to metal parts. In the case of exterior panels, the composite material can help reduce mass up to 50 percent.

They are produced using a highly efficient, one-shot process of lamination and low-pressure molding. “Our Udmax tapes offer the automotive industry a powerful solution to the ongoing challenges of reducing weight, lowering costs and improving sustainability,” said Hans Warmerdam, CEO and Chief Sales & Marketing Officer, Sabic FRT – a Sabic affiliate. “We’re confident that the light commercial vehicle’s bulkhead is the first of many structural applications where our innovative materials, combined with this novel processing approach, can help solve our customers’ challenges of achieving lighter weight without compromising safety, durability, and fuel or energy use. Through continued collaboration among this unique team of engineers, designers and technical experts in materials and in conversion processes, we intend to explore more ways to expand the adoption of our thermoplastic composite technology.”

Compared to metallic or injection molded part of a conventional, multi-piece bulkhead, the new method – designed, developed and engineered by RLE International – can reduce tool costs by up to 80 percent compared to injection molding tools. This saving is due to the ability to replace an expensive, high- pressure tool with a lower-cost, low-pressure tool. Overall, the supplied cost of the LCV’s bulkhead can be 10 percent lower than the conventional metallic bulkhead that it replaces.

This technology also represents a revolutionary change in vehicle panel production by increasing efficiency and reducing complexity. With molding cycle times under two minutes, this streamlined process avoids sourcing and assembling multiple components, traditionally at different supplier locations, as well as secondary painting and trim operations. A proprietary lamination featuring a core of extra-wide Udmax tape woven by Setex incorporates aesthetic finish a one-shot compression step.

The process also allows varying the thickness of the panel in order to improve noise, vibration, and harshness (NHV) levels, helping to reduce noise in the vehicle. AMA Composites created the tool and molded the concept parts. “Our new technology for producing panels using Udmax thermoplastic composite tape offers tremendous opportunities to the automotive industry,” said Mark Grix, head of Interior & Exterior Engineering for RLE International. “One example is the electric vehicle sector, where lower panel weight can extend driving range and lower-cost tooling can reduce capital investments for start-up companies.

RLE International stands ready to assist automotive tiers in mastering this new process so they can leverage its advantages on behalf of their OEM customers.”

Source: Plastics Today - www.plasticstoday.com

Polímeros, 29(1), 2019 E3

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

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

July

October

6th Global Conference on Polymer and Composite Materials Date: July 8-11, 2019 Location: Bangkok, Thailand Website: www.cpcmconf.org 3rd International conference on Polymer Materials and Nanocomposites Date: July 17-19, 2019 Location: Aveiro, Portugal Website: www.advanced-nanomaterials-conference.com Polymer Composites and High Performance Materials Date: July 22-25, 2019 Location: Sonoma, United States Website: polyacs.net/Workshops/19Composites/home.html 5th Edition of International Conference on Polymer Science and Technology Date: July 30-31, 2019 Location: Amsterdam Netherlands Website: polymerscience.annualcongress.com

Conference Arab Rubber Expo 2019 Date: October 9-10, 2019 Location: Sharjah, United Arab Emirates Website: http://www.rubber-expo.com 3rd International Conference on Metal Organic Frameworks and Porous Polymers Date: October, 27-30, 2019 Location: Paris, France Website: euromof2019.sciencesconf.org 15th Congresso Brasileiro de Polímeros (15th CBPol) Date: October, 27-31, 2019 Location: Bento Gonçalves, Rio Grande do Sul, Brazil Website: www.cbpol.com.br Polymers + 3D Date: October 31 - November 01, 2019 Location: Houston, United States Website: www.poly3d.org

August

Conductive Plastics 2019 Date: November 14-17, 2019 Location: Jakarta, Indonesia Website: www.plasticsandrubberindonesia.com 31st International Plastics & Rubber Machinery, Processing & Materials Exhibition Date: November 14-17, 2019 Location: Jakarta, Indonesia Website: www.plasticsandrubberindonesia.com 10th International Conference on Biopolymers and Polymer Sciences Date: November 18-19, 2019 Location: Helsinki, Finland Website: biopolymers.materialsconferences.com PPS Europe-Africa 2019 Regional Conference Date: November 18-21, 2019 Location: Pretoria, South Africa Website: http://www.pps2019.com Plastics & Rubber Vietnam Date: November 27-29, 2019 Location: Hanoi, Vietnam Website: plasticsvietnam.com

6th International Conference and Exhibition on Polymer Chemistry Date: August 02-03, 2019 Location: Chicago, United States Website: polymer.conferenceseries.com

September 33rd Polymer Degradation Discussion Group Conference Date: September 1-5, 2019 Location: St. Julian, Malta Website: pddg.org Performance Polyamides 2019 Date: September 4-5, 2019 Location: Cologne, Germany Website: www.ami.international/events/event?Code=C0993 13th International Symposium on Ionic Polymerization (IP 2019) Date: September 8-13, 2019 Location: Beijing, China Website: iupac.org/event/international-symposium-on-ionicpolymerization-ip-19 Brightlands Rolduc Polymer Conference (BRPC 2019) Date: September 9-11, 2019 Location: Kerkrade, Netherlands Website: www.klinkhamergroup.com/polymer2019 International Rubber Conference (IRC 2019) Date: September 10-12, 2019 Location: London, United Kingdom Website: www.iom3.org/rubber-engineering-group/event/ international-rubber-conference-irc-2019 International Conference on Materials Science and Engineering Date: September 16-18, 2019 Location: Melbourne, Australia Website: www.materialsconferenceaustralia.com Polymer Testing & Analysis – 2019 Date: September 18-19, 2019 Location: Düsseldorf, Germany Website: www.ami.international/events/event?Code=C0990 Advances in Polyolefins (APO-2019) Date: September 22-25, 2019 Location: California, United States Website: www.polyacs.net/19apo 10th European Symposium on Biopolymers Date: September 25-27, 2019 Location: Straubing, Germany Website: www.esbp2019.com

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November

December Polymer Chemistry Date: December 02-03, 2019 Location: Tokyo, Japan Website: polymer.chemistryconferences.org Conference Latex Expo 2019 Date: December 4-5, 2019 Location: Chennai, India Website: www.latex-expo.com 16th Pacific Polymer Conference (PPC16) Date: December 8-12, 2019 Location: Singapore City, Singapore Website: www.pacificpolymer.org 6th World Congress on Smart Materials and Polymer Technology Date: December 16-17, 2019 Location: Dubai, United Arab Emirates Website: smart.materialsconferences.com

Polímeros, 29(1), 2019

ABPol Associates Sponsoring Partners

Collective Members Master Polymers Ltda.

Nexo International Ltda.

Nitriflex S/A Ind. e Com. Radici Plastics Ltda.

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

PolĂ­meros, 29(1), 2019

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ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.01118

Synthesis and characterization of the photoswitchable poly(methyl methacrylate-random-methacrylate spirooxazine) Tam Huu Nguyen1, Phuc Huynh Tran2, Linh Duy Thai2, Thuy Thu Truong1, Le-Thu T. Nguyen1 and Ha Tran Nguyen1,3*  Faculty of Materials Technology, Ho Chi Minh City University of Technology – HCMUT, Vietnam National University, Ho Chi Minh City, Vietnam 2 Faculty of Chemical Technology, Ho Chi Minh City University of Technology – HCMUT, Vietnam National University, Ho Chi Minh City, Vietnam 3 National Key Laboratory of Polymer and Composite Materials, Vietnam National University, Ho Chi Minh City, Vietnam 1

*nguyentranha@hcmut.edu.vn

Abstract The photoswitchable poly(methyl methacrylate-random-methacrylate spirooxazine) was synthesized via atom transfer radical polymerization with the feed mole ratio of MMA/MSp comonomer of about 5.5/1. Well-defined poly(methyl methacrylate-random-methacrylate spirooxazine) have been obtained with the average molecular weight (Mn) of 6500 g/mol and polydispersity of 1.21. The structure and properties of the resulting copolymers were characterized by proton nuclear magnetic resonance (1H NMR), gel permeation chromatography, Fourier Transform infrared, UV-visible spectroscopy, and differential scanning calorimetry. Moreover, the copolymer exhibited the erasable and rewritable photoimaging on the solid state film which could to be as potential candidate for optical data storage materials.

Keywords: spirooxazine, controlled radical polymerization, photoswitching polymers.

How to cite: Nguyen, T. H., Tran, P. H., Thai, L. D., Truong, T. T., Nguyen, L.-T. T., & Nguyen, H. T. (2019). Synthesis and characterization of the photoswitchable poly(methyl methacrylate-random-methacrylate spirooxazine). Polímeros: Ciência e Tecnologia, 29(1), e2019001. https://doi.org/10.1590/0104-1428.01118

1. Introduction The creation of functional optical materials having photoactive properties has become one of the most promising objects in materials science. These materials are used in the fabrication of several optoelectronic devices such as optical memories, switches, and holograms. These integrated systems are basically formed by two components: the support media and the photoactive material, most of them are polymers functioned with photoactive molecules[1]. Photochromism has attracted much attention recently from the viewpoint of optical applications because of interest in refractive index or absorbance changes through optical extraction.

Since the last decade, the development of such optoelectronic devices has included the photochromic compounds as the active ingredient[1]. The photochromism of spiropyrans (SP) was reported by Fischer and Hirshberg[2] in 1952 and since these organic compounds have been extensively studied due to their possible application in many fields. Following, spirooxazine compounds synthesized belonged to the spiroindolinonaphthoxazine ring system.

Spirooxazine being a class of photochromic compounds are potentially applicable as chemical-UV-dosimeters for personal protection[3,4] or they can be incorporated into the materials used for packaging applications in the case

Polímeros, 29(1), e2019001, 2019

of UV sensitive products like food products. Generally, there is a great interest for the properties of photochromic compounds inserted into polymer films due to their possible use in practical applications. Spirooxazine are composed of an imide and a chromene moieties that are linked by a spirocarbon atom. Irradiation of spirooxazine with UV light induces heterolytic cleavage of the spiro-carboneoxygen bond, thus, producing the ring opened form, the intensively coloured merocyanine (MC). Merocyanine returns to the initial spirooxazine form in the dark or by visible light irradiation[5-7]. The time to resume the initial colour depends on temperature and on the nature of the compound. In recent years, photochromic and thermochromic spiropyrans and spiroxazines have been receiving considerable attention, due to their potential application in many new technologies, such as data recording and storage, optical switching, displays, and non-linear optics[8-12]. Recently, Ventura et. al have reported the synthesis of new and well-defined poly(6-benzospiropyran hexylmethacrylate)s bearing a BSP moiety on the side chain of each unit of the polymer with linear, star-like and molecular brush architectures and narrow molecular weight distributions were successfully synthesized combining ATRP and click chemistry[13]. More

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Nguyen, T. H., Tran, P. H., Thai, L. D., Truong, T. T., Nguyen, L.-T. T., & Nguyen, H. T. recently, Dübner et. al have published the copolymer brushes, composed of glycidyl methacrylate and a furanprotected maleimide-containing monomer, after postpolymerization modification, they have functioned microperoxidase-11 and photochromic spiropyran moieties, the polymer brushes catalyzed the oxidation of 3,3′5,5′- tetramethylbenzidine.

Their obtained brushed copolymers exhibited the light‑induced spiropyran-merocyanine transition under UV or visible-light, turnover by more than 1 order of magnitude[14].

Here we reported the synthesis of poly(methyl methacrylate)‑r-poly(methacrylate spirooxazine) (PMMA‑rPMSp) via atom transfer radical polymerization (ATRP).

The well- defined copolymers were characterized via 1H NMR, GPC, FTIR, DSC, UV-vis spectrum. Especially, the synthesized copolymers exhibited the spiroozazine‑merocyanine transition under UV irradiation for both solution and solid film state.

2. Materials and Methods 2.1 Materials 2,7-Dihydroxynaphthalene (97%), sodium nitrite (NaNO2), sodium hydroxide (NaOH), sulfuric acid (H2SO4), hydrochloric acid (HCl), potassium carbonate (K2CO3), 1,3,3-trimethyl-2-methyleneindoline (97%), triethylamine (99%), methacryloyl chloride (97%), Copper(I) bromide (CuBr, 98%), N,N,N′,N″,N″-pentamethyldiethylenetriamine (PMDETA, 99%), ethyl α-bromoisobutyrate (Et2BriB, 98%), 4-(Dimethylamino)pyridine (99%) were purchased from Aldrich. Methyl methacrylate (MMA, 98%) was purchased from Sigma-Aldrich passed through a plug of basic alumina before use in order to remove the hydroquinone inhibitors and stored under nitrogen atmosphere. Dichloromethane (99.8%) and tetrahydrofuran (THF, 99%) were purchased from Fisher/Acros and dried using molecular sieves under N2. Chloroform (CHCl3, 99.5%), hexane (99%), methanol (99.8%), absolute ethanol (99%) and ethyl acetate (99%) were purchased from Fisher/Acros and used as received.

2.2 Characterization H NMR spectra were recorded in deuterated chloroform (CDCl3) with TMS (δ 0.00 ppm). The following abbreviations are used to describe the NMR signals: s (singlet), d (doublet), t (triplet), q (quartet), and br (broad). FT-IR spectra, collected as the average of 264 scans with a resolution of 4 cm−1, were recorded from KBr disk on the FT-IR Bruker Tensor 27.

Size exclusion chromatography (SEC) measurements were performed on a Polymer PL-GPC 50 gel permeation chromatograph system equipped with an RI detector, with THF as the eluent at a flow rate of 1.0 mL/min. Molecular weight and molecular weight distribution were calculated with reference to polystyrene standards. UV-vis absorption spectra of polymers in solution and polymer thin films were recorded on a Shimadzu UV-2450 spectrometer over a wavelength range of 250-800 nm.

1

2.3 Synthesis of 1-nitrosonaphthalene-2,7-diol After dissolving NaOH (2.5 g, 62.4 mmol) in 100 mL of H2O, 2,7-dihydroxynaphthalene (1) (10 g, 62.4 mmol) and Na2NO2 (4.3 g, 62.4 mmol) were added to the solution.

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The reaction solution was heated for 1 h at 60 °C and then cooled to 0 °C. The mixture of 8 mL of concentrated H2SO4 and 15 mL of distilled water was added dropwise to the reaction solution with the temperature remained at 0 °C. The reaction continued for 1 h. After the reaction, the precipitate was isolated by vacuum filtration and washed with 0.1 M aqueous HCl followed by cold water to obtain compound 2 as a dark brown powder solid. Yield: 11.22 g, 95%.

1 H NMR, (500 MHz, methanol-d4), δ (ppm): 7.42 (d, 1H), 7.59 (d, 1H), 7.56 (d, 1H), 6.8 (d, 1H), 6.18 (s, 2H). FT-IR (cm−1): 3143 (O–H), 1301 (N=O).

2.4 Synthesis of 1,3,3-trimethylspiro[indoline-2,3’naphtho[2,1-b][1,4]oxazin]-9’-ol (spirooxazine-hydroxyl) To a suspension of 2,7-dihydroxy-1-nitrosonaphthaline (compound 1) (2.84 g, 15 mmol) in absolute ethanol (50 mL) was added dropwise, under refluxing a solution of 1,3,3-trimethyl-2-methyleneindoline (2.59 g, 2.65 mL, 15 mmol) in absolute ethanol (20 mL). After continuous refluxing under a N2-stream, the obtained brown solution was purified over silica column with ethyl acetate/hexane (2:1) to obtain the crude product. Then, solvents were evaporated under vacuum to give a black powder. The black powder was washed with distilled water and extracted with CHCl3.

Finally, the product was crystallized in hexane to obtain the pure white powder of SP. The product (2.74 g, yield: 53%) was dried in a vacuum oven at RT and was characterized by 1H NMR and FT-IR. 1 H NMR, (500 MHz, CDCl3), δ (ppm): 1.35 (s, 6H), 2.77 (s, 3H), 6.58 (t, 1H), 6.84 (d, 1H), 6.9 (d, 1H), 7.02 (d, 1H), 7.09 (d, 1H), 7.23 (t, 1H), 7.58 (d, 1H), 7.65 (d, 1H), 7.71 (s, 1H), 7.88 (s, 1H), 10.04 (s, 1H, OH).

FT-IR (cm−1): 3313 (O–H), 3065 (=C–H), 1627 (C=N), 1301 (N=O).

2.5 Synthesis of the methacrylate spirooxazine monomer (MSp) Spirooxazine-hydroxyl (1.03 g, 3.00 mmol) was added to 25 mL of anhydrous dichloromethane in a 50 mL round bottomed flask. After cooling the solution to 0 °C, triethylamine (0.46 g, 0.627 mL, 4.5 mmol) was added and the reaction was stirred for an hour. Then, methacryloyl chloride (0.47 g, 0.44 mL, 4.5 mmol) was dissolved in 5 mL of anhydrous dichloromethane and added dropwise to the reaction mixture under N2 atmosphere, cooled to 0 °C. The reaction was continuously stirred for 24 h at room temperature. The solvent was rotary evaporated and the crude product was recrystallized in hexane. This crude product was purified over column silica hexane/ethyl acetate (v:v = 30/1) to obtain the pure white powder of Msp. The product was dried in a vacuum oven at RT and was characterized by 1H NMR and FT-IR. Yield: 75%.

1 H NMR, (500 MHz, CDCl3), δ (ppm): 1.35 (s, 6H), 2.1 (s, 3H), 2.77 (s, 3H), 5.6 (s, 1H), 6.3 (s, 1H), 6.58 (dd, 1H), 6.84 (dd, 1H), 6.9 (s, 1H), 7.02 (d, 1H), 7.09 (d, 1H), 7.23 (d, 1H), 7.66 (d, 1H), 7.72 (s, 1H), 7.78 (d, 1H), 8.28 (d, 1H). FT-IR (cm-1): 2971 (=C–H), 1728 (C=O), 1627 (C=N), 1121, 740.

Polímeros, 29(1), e2019001, 2019

Synthesis and characterization of the photoswitchable poly(methyl methacrylate-random-methacrylate spirooxazine) 2.6 Synthesis of Poly(methyl methacrylate-comethacrylate spirooxazine) (Poly(MMA-co-MSp)) Copolymer Poly(MMA-co-MSp) were synthesized using a ratio of [MMA]:[MSp]:[Et2BriB]:[CuBr]:[PMDE TA] = 55:10:1:1:2. Into a dried glass tube with a magnetic bar, CuBr (2.15 mg g, 0.015 mmol) and ligand PMDETA (26 mg, 0.15 mmol) were added, then monomer MMA (82.60 mg, 0.75 mmol), MSp (61.87 mg, 0.15 mmol), and anhydrous THF (2 mL) were added under nitrogen. The mixture was degassed by three freeze-pump-thaw cycles and purged with nitrogen. Et2BriB (2.93 mg, 0.015 mmol) was added before the polymerizations were carried out at 65 °C oil bath for 24 h. The flask was then opened and exposed to air.

The copolymer was recovered by removing the solvent under reduced pressure, re-dissolving the polymer in chloroform, passing the polymer solution through an alumina column to remove excess copper catalyst and was then concentrated by using a rotary evaporator. The concentrated solution was added in drops into an approximately 80 mL of cold methanol. The precipitation product was filtered through a medium frit funnel. Finally, the polymer was dried under vacuum at 60 °C for 24 h until a weight constant. Yield 75%, Mn = 8000 g/mol, Mw/Mn = 1.34, as obtained by gel permeation chromatography.

1 H NMR (500 MHz, CDCl3) δ (ppm): 0.8-1.3 (s, 15H), 1.35 (s, 6H), 1.6-2.4 (s, 4H), 2.76 (s, 3H), 3.60 (s, 3H), 4.09 (m, 2H), 6.58 (dd, 1H), 6.84 (dd, 1H), 6.9 (s, 1H), 7.02 (d, 1H), 7.09 (d, 1H), 7.23 (d, 1H), 7.66 (d, 1H), 7.72 (s, 1H), 7.78 (d, 1H), 8.24 (d, 1H).

3. Results and Discussions 3.1 Monomer synthesis First, 2,7-Dihydroxynaphthalene (compound 1) was nitrosated by NaNO2 in presence of NaOH and H2SO4 to obtain 1-nitrosonaphthalene-2,7-diol (compound 2, Scheme 1). Then, 1-nitrosonaphthalene-2,7-diol reacted with 1,3,3-trimethyl2-methyleneindoline to form spirooxazine-hydroxyl (compound 3, Scheme 1), which subsequently reacted with methacryloyl chloride to give methacrylate-spirooxazine

(MSp) (compound 4, Scheme 1). Finally, the synthesis of poly(MMA-co-MSp) (coumpound 5, Scheme 1) was performed via atom transfer radical polymerization using CuBr and PMDETA as catalyst system. The synthesis of poly(MMA-co-MSp) was described as Scheme 1.

The FTIR was used to characterize the spirooxazine hydroxyl (compound 3) and methacrylate spirooxazine (MSp) (compound 4) to determine the chemical functional groups in their structure. Figure 1 showed the FTIR of spirooxazine hydroxyl and methacrylate-spirooxazine, the vibration peaks at 1630 cm-1 corresponding to N=C linkage and the broad peaks at 3300 cm-1 assigned to OH group in spirooxazine hydroxyl (green line c, Figure 1). Following, Spirooxazine-hydroxyl was reacted with methacryloyl chloride to give methacrylate-spirooxazine (MSp) in the presence of triethylamine as catalytic. The reaction was performed in 24 hours. The obtained product was purified over column using heptane/ethyl acetate as eluent, and the purified product was recrystallized in methanol in the yield of 75%. The obtained methacrylate spirooxazine exhibited the vibration peaks at 1750 cm-1 corresponding to C=O linkage, and the peaks of OH group at 3300 cm-1 disappears completely that confirmed that the esterification of spirooxazine hydroxyl with methacryloyl chloride was taken place successfully (cyan line d, Figure 1).

3.2 Polymer synthesis The 1H NMR spectrum of spirooxazine hydroxyl showed the proton resonance of imine linkage at 7.65 ppm. Moreover, the signals at 2.65 ppm and 1.42 ppm exhibited for the methyl groups of 1,3,3- trimethyl-2-methyleneindoline in sprirooxazine hydroxyl moieties. The 1H NMR spectrum of MSp (Figure 2) showed the proton resonance of methylene linkage of methacrylate at 5.70 ppm and 6.30 ppm. The 1H NMR spectrum of MSp also showed the imine linkage in methacrylate - spirooxazine ring at 8.4 ppm which was shifted from 7.65 ppm in spirooxazine-hydroxyl. All other proton resonances appear in the reasonable intensities and are correlated to spirooxazine structure, thus confirming the expected molecular structure of MSp. The attribution

Scheme 1. Synthesis of poly(methyl methacrylate-co-methacrylate spirooxazine).

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Nguyen, T. H., Tran, P. H., Thai, L. D., Truong, T. T., Nguyen, L.-T. T., & Nguyen, H. T.

Figure 1. FTIR of monomer compounds: (a) FTIR of 2,7-dihydroxynaphthalene compound; (b) FTIR of 1-nitrosonaphthalene-2,7-diol compound; (c) FTIR of spirooxazine-hydroxyl; (d) FTIR of methacrylate spirooxazine monomer.

Figure 2. The 1H NMR of spirooxazine hydroxyl and methacrylate spirooxazine (MSp): (A) 1H NMR of spirooxazine hydroxyl; (B) 1H NMR of methacrylate spirooxazine).

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Polímeros, 29(1), e2019001, 2019

Synthesis and characterization of the photoswitchable poly(methyl methacrylate-random-methacrylate spirooxazine) of the 1H NMR signals of spirooxazine hydroxyl and MSp is presented in Figure 2[15].

In final step, the MSp and MMA comonomers were polumerized via ATRP using ethyl α-bromoisobutyrate in the presence of CuBr and PMDETA as catalyst and ligand, respectively. The feed ratio of MMA/MSp comonomer of about 5.5/1 was established for achieving a good control over ATRP ([MMA]/[MSp]/[Initiator]/[CuBr]/[PMEDETA] = 55:5:1:1:2). The polymerization was performed in THF at 60 °C for 24 h under nitrogen atmosphere. The polymerization was stopped by cooling the reaction mixture followed by dilution with extra volume of THF, then the mixture was purified over aluminium column to remove CuBr catalyst.

The copolymers were obtained by precipitation in cold n-heptane, the copolymers were filtered and dried under vacuum until constant mass.

The FT-IR spectra was used to characterize the copolymers Poly(MMA-r-MSp) as Figure 3. The appearance of the high intensity signal observed at 1728 cm-1 that attributed to carbonyl vibrational (C=O) of Poly(MMA-r-MSp).

In addition, the N=C linkage of MSp in Poly(MMA-r-MSp) was also exhibited at 1651 cm-1.

The polymer structure of Poly(MMA-r-MSp) was also confirmed by 1H NMR. The Figure 4 exhibited all characteristic peaks of copolymers Poly(MMA-r-MSp).

The polymerization degree of the Poly(MMA-r-MSp) was calculated from recorded 1H NMR spectrum by comparing the relative signal intensities of the imine proton of the MSp and methylene protons of the MMA residue at δ = 8.24 ppm (peak “s”, Figure 4) and δ = 3.6 ppm (peak “f”, Figure 4), respectively, with that of the methylene protons of initiator at 4.09 ppm (peak “b”, Figure 4). The molecular weight of the copolymers Poly(MMA-r-MSp) was determined to be 6500 g/mol and to compise 52 and 3 of MMA and MSp units, respectively. These results give the weight compositions of 80% and 20% of MMA and MSp, respectively in copolymers Poly(MMA-r-MSp).

As seen from Table 1, the Poly(MMA-r-MSp) was obtained with a relatively good approximation between theoretical and experiment molar masses that approved for an initiation efficiency close to 1. A narrow molecular weight distribution of the copolymers Poly(MMA-r-MSp) was recorded by GPC, with ĐM = 1.21.

3.3 Photoisomerization properties of polymers

Figure 3. The FT-IR of copolymers Poly(MMA-r-MSp).

Photoisomerization property of poly(MMA-r-MSp) was studied in THF. Figure 5A and Figure 5B show the color changing of polymer solution (0.1 g/L; pH 9.0) measured at 25 oC under UV irradication. Without UV irradiation, the

Figure 4. The 1H NMR spectrum of copolymers Poly(MMA-r-MSp).

Polímeros, 29(1), e2019001, 2019

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Nguyen, T. H., Tran, P. H., Thai, L. D., Truong, T. T., Nguyen, L.-T. T., & Nguyen, H. T. Table 1. Macromolecular characteristics of Poly(MMA-r-MSp) synthesized by ATRP using α-bromoisobutyrate initiator and CuBr/ PMDETA ([CuBr]/[PMDETA] = 1/2) as the catalytic complex.

Diblock copolymers

T °C

1 2

60 60

Conversion (%)(a) 92 87

MMA (b)

M n the

(g/mol) 5060 4785

MSp (c)

M n exp

(g/mol) 5200 5100

(b)

M n the

(g/mol) 2150 2150

Poly(MMA-r-MSp) (c)

M n exp

(g/mol) 1300 1300

f(d) 1 1

M n MMR (g/mol) 6500 6400

Ð(e) 1.21 1.25

Conversion as determined after precipitation in cold n-heptane: Conv.= (m-mI-mCu-mL)/mM where m denotes the weight of product, and mI, mCu, mL, mM the weights of the initiator, copper catalyst, ligand (PMDETA), and monomers, respectively; (b)MMA and MSp theoretical number‑average‑molar mass as calculated by [MMA] or [MSp)]0/[Initiator]0 x Conv(%) x Mw MMA(or MSp) assuming a living process; (c)MMA (or MSp) (a)

experimental number-average molar mass as determined by 1H NMR spectroscopy (see Figure 3): M n exp = DPEXP x MW MMA (OR MSP) where DPEXP is the experimental degree of polymerization, as calculated from the relative intensities of methylene protons of MMA (δ = 3.6 ppm), imine proton of MSp (δ = 8.24 ppm) and methylene protons of initiator (4.09 ppm); (d) Initiation efficiency as calculated from M n the of P(MMA-R-MSp) / M n exp of P(MMA-R-MSp); (e) Dispersity index as determined by GPC in THF at 35 °C.

Figure 5. Poly(MMA-r-MSp) in THF solutionwith MSp in closed form (A), Poly(MMA-r-MSp) in THF solution with MSp in opened form after UV irradication (B), UV-Vis spectra of poly(MMA-r-MSp) under UV-irradiation (C).

polymer shows almost no absorption in the visible region (blue line, Figure 5C) indicating that the spirooxazine units exist as a closed form, as shown in Figure 5C. However, UV irradiation of the solution (334 nm), creates a distinctive absorption band centered at 540 nm, assigned to the generation of opened form of spirooxazine.

In addition, we measured the absorption spectra changes of poly(MMA-r-MSp) to obtain an insight into their photochromic properties in the solid state film which shows similarly photochromic performance in solution.

The copolymers poly(MMA-r-MSp) solution was coated on transparent glass substrate via spin coating method, then 6/8

the obtained polymer film was dried in oven at 60 °C in 2 hours . Irradiation a colorless poly(MMA-r-MSp) film, the colorless poly(MMA-r-MSp) turned to blue and a new absorption band appears at around 550 nm and gradually increased and reached a photostationary state.

The inherent characteristics of poly(MMA-r-MSp) film make it possible to use such materials for data recording.

A possible procedure for data recording and erasing is presented in Figure 6. Upon UV light irradiation through the mask, the optical data were recorded on poly(MMA-r-MSp) film irradiation region, when irradiation with visible light on irradiation region, the optical data were erased. In solid Polímeros, 29(1), e2019001, 2019

Synthesis and characterization of the photoswitchable poly(methyl methacrylate-random-methacrylate spirooxazine)

Figure 6. Principle scheme of the optical data recording on poly(MMA-r-MSp) film.

copolymer poly(MMA-r-MSp) has the average molecular weight of 6500 g/mol with polydispersity index of around 1.2-1.3. The copolymer exhibited excellent photochromic behavior in solid film under UV irradiation. Erasable and rewritable photoimaging on the solid film was successfully demonstrated. This optical data storage materials based on photochromic poly(MMA-r-MSp) could to be as candidate for fundamental studies and eventual technical application for all-photo mode – high density optical data.

Figure 7. The DSC diagram of poly(MMA-r-MSp).

state, the practical capability of rewritable photoimaging on solid state investigated by patterned illumination through photomasks. The word ‘ĐHBK’ (Ho Chi Minh City University of Technology, Vietnamese abbreviated) was recorded as a first image (Figure 6), which was subsequently erased and followed by the recording of a second image. The cycles of writing and erasing was repeated more than 30 times.

Last but not least, the thermal properties of poly(MMA‑r‑MSp) (Mn = 6500 g/mol, Đ = 1.21) was characterized to determine the glass transition of copolymers

by DSC . The DSC second-heating traces in the range from 0 to 150 °C of the poly(MMA-r-MSp) are shown in Figure 7. For such studied molecular weight, the glass transition temperature (Tg) value of poly(MMA-r-MSp) is around 66 °C. Otherwise, the melting point (Tm) of poly(MMA‑r–MSp) is observed at ca. 108 °C which is similar to the DSC results of pure homopolymer PMMA in the same of average molecular weight.

4. Conclusions In conclusion, we have successfully designed and synthesized a novel photoswitching poly(MMA-r-MSp) via controlled polymerization method. The obtained Polímeros, 29(1), e2019001, 2019

5. Acknowledgements The study was supported by Science and Technology Incubator Youth Program, managed by the Center for Science and Technology Development, Ho Chi Minh Communist Youth Union, the contract number is “15/2017/HĐ-KHCN-VƯ “, and project “C2017-20-31” from Ho Chi Minh City University of Technology - Vietnam National University – Ho Chi Minh City, 268 Ly Thuong Kiet, District 10, Ho Chi Minh City, Viet Nam.

6. References 1. Mizokuro, T., Mochizuki, H., Kobayashi, A., Horiuchi, S., Yamamoto, N., Tanigaki, N., & Hiraga, T. (2004). Selective doping of photochromic dye into nanostructures of diblock copolymer films by vaporization in a vacuum. Chemistry of Materials, 16(18), 3469-3475. http://dx.doi.org/10.1021/ cm049557u.

2. Fischer, E., & Hirshberg, Y. (1954). Photochromism and reversible multiple internal transitions in some spiropyrans at low temperatures. Journal of the Chemical Society, 297-303. http://dx.doi.org/10.1039/JR9540000297.

3. Lokshin, V. A., Samat, A., & Metelitsa, A. V. (2002).

Spirooxazines: synthesis, structure, spectral and photochromic properties. Russian Chemical Reviews, 71(11), 893-916. http:// dx.doi.org/10.1070/RC2002v071n11ABEH000763.

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Nguyen, T. H., Tran, P. H., Thai, L. D., Truong, T. T., Nguyen, L.-T. T., & Nguyen, H. T. 4. Bouas-Laurent, H., & Dürr, H. (2001). Organic photochromism.

Pure and Applied Chemistry, 73(4), 639-665. http://dx.doi. org/10.1351/pac200173040639.

5. Berkovic, G., Krongauz, V., & Weiss, V. (2000). Spiropyrans and spirooxazines for memories and switches. Chemical Reviews, 100(5), 1741-1754. http://dx.doi.org/10.1021/cr9800715.

PMid:11777418.

6. Lin, J. S. (2003). Interaction between dispersed photochromic compound and polymer matrix. European Polymer Journal, 39(8), 1693-1700. http://dx.doi.org/10.1016/S0014-3057(03)00058-2.

7. Ock, K., Jo, N., Kim, J., Kim, S., & Koh, K. (2001). Thin film optical waveguide type UV sensor using a photochromic molecular device, spirooxazine. Synthetic Metals, 117(1-3), 131-133. http://dx.doi.org/10.1016/S0379-6779(00)00553-1.

8. Myles, A. J., Wigglesworth, T. J., & Branda, N. R. (2003). A multi-addressable photochromic 1,2-dithienylcyclopentenephenoxynaphthacenequinone hybrid. Advanced Materials, 15(9), 745-748. http://dx.doi.org/10.1002/adma.200304917.

9. Corredor, C. C., Huang, Z. L., & Belfield, K. D. (2006). Two‐ photon 3D optical data storage via fluorescence modulation of an efficient fluorene dye by a photochromic diarylethene.

Advanced Materials, 18(21), 2910-2914. http://dx.doi. org/10.1002/adma.200600826.

10. Lim, S. J., Seo, J., & Park, S. Y. (2006). Photochromic switching of excited-state intramolecular proton-transfer (ESIPT) fluorescence: a unique route to high-contrast memory switching and nondestructive readout. Journal of the American Chemical Society, 128(45), 14542-14547. http://dx.doi.org/10.1021/ ja0637604. PMid:17090038.

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11. Lim, S. J., An, B. K., & Park, S. Y. (2005). Bistable photoswitching in the film of fluorescent photochromic polymer: enhanced fluorescence emission and its high contrast switching.

Macromolecules, 38(15), 6236-6239. http://dx.doi.org/10.1021/ ma0504163.

12. Jiang, G., Wang, S., Yuan, W., Jiang, L., Song, Y., Tian, H., & Zhu, D. (2006). Highly fluorescent contrast for rewritable optical storage based on photochromic bisthienylethene-bridged naphthalimide dimer. Chemistry of Materials, 18(2), 235-237. http://dx.doi.org/10.1021/cm052251i.

13. Ventura, C., Thornton, P., Giordani, S., & Heise, A. (2014).

Synthesis and photochemical properties of spiropyran graft and star polymers obtained by ‘click’ chemistry. Polymer Chemistry, 5(21), 6318-6324. http://dx.doi.org/10.1039/C4PY00778F. 14. Qu, W.-J., Li, W.-T., Zhang, H.-L., Wei, T.-B., Lin, Q., Yao, H., & Zhang, Y.-M. (2017). A rational designed fluorescent and colorimetric dual-channel sensor for cyanide anion based on the PET effect in aqueous medium. Sensors and Actuators.

B, Chemical, 241, 430-437. http://dx.doi.org/10.1016/j. snb.2016.10.100.

15. Kim, S. H., Hwang, I. J., Gwon, S. Y., & Son, Y. A. (2010).

Photoregulated optical switching of poly(N-isopropylacrylamide) hydrogel in aqueous solution with covalently attached spironaphthoxazine and D-π-A type pyran-based fluorescent dye. Dyes and Pigments, 87(2), 158-163. http://dx.doi. org/10.1016/j.dyepig.2010.03.014.

Received: Feb. 10, 2018 Revised: Apr. 16, 2018 Accepted: Apr. 28, 2018

Polímeros, 29(1), e2019001, 2019

ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.05617

Natural ageing of polyaramide fiber from ballistic armor Vitor Hugo Cordeiro Konarzewski1* , Fernando Ludgero Spiekemann1 and Ruth Marlene Campomanes Santana1 Laboratório de Polímeros – LAPOL, Universidade Federal do Rio Grande do Sul – UFRGS, Porto Alegre, RS, Brasil

1

*vitorhk@sinos.net

Abstract Ballistic armor has been manufactured primarily based on polyaramide (Kevlar and Twaron) or Dyneema but the lifespan warranty in Brazil is only 5 years and after this time period they are incinerated or comminuted and ground up. This study aims to evaluate the changes on the physical, mechanical and morphological properties of polyaramide fibers of ballistic armor after natural aging. These samples with different fabrication (2005 and 2010) and usage time were exposed to natural weathering in the city of Porto Alegre, southern Brazil, during the period of one year. Morphology fiber results surfaced after ageing, it showed fiber swelling, stress cracking and defibrillation, and the results of the mechanical tensile testing of the polyaramide fibers showed a pronounced decrease (80%) in tensile strength. It can be concluded that the weight, the dtex of the fiber and the kind of fabric can influence the degradation degree under natural exposure.

Keywords: ballistic armor, aramid fibers, polymer degradation.

How to cite: Konarzewski, V. H. C., Spiekemann, F. L., & Santana, R. M. C. (2019). Natural ageing of polyaramide fiber from ballistic armor. Polímeros: Ciência e Tecnologia, 29(1), e2019002. https://doi.org/10.1590/0104-1428.05617

1. Introduction The ballistic armor market in Brazil is approximately one million users, who act directly in police actions and other inherent activities[1,2]. With an expiration date of five years from the date of manufacture, the legislation in Brazil determines the destruction of the material by shredding or incineration after the theoretical expiry date, without taking into account the commercial value of the polyaramide or the stability of the raw material[3-5].

The use of lightweight ballistic vests was increased after the 1960s with the development of high-performance aramid fibers, or poly (p-phenylene terephthalamide), referred to as Kevlar by DuPont, registered in 1971[6].

Another registered polyaramide trademark of Teijin, Twaron, whose properties of heat resistance, flexibility, dimensional stability and high mechanical resistance have made the polyaramide fibers quickly gain space in the most demanding industrial applications and also in the security area, being subject to constant study and improvement for use in ballistic armor[7-9]. There are usually two basic characteristics that respond to the trust placed in the aramid shield until now: the chemical composition of the fiber and the way the yarns are woven and interlaced in multilayers. Under the impact of the projectile, the energy is dissipated through the panel, by friction, absorption and dissipation of the shock waves avoiding perforation[10-17].

The aromatic polyamides were obtained by reactions that lead to the formation of amide bonds between aromatic rings of high thermal stability and high strength[18-21].

The exceptionally rigid molecular chain structure, good orientation and organization of the crystalline structure provides high strength and low elongation of the orientation

Polímeros, 29(1), e2019002, 2019

of the molecular chains, offering high tensile strength, impact and with differentiated thermal stability for various temperature ranges for an extended time[22-25].

The susceptibility of aramid fibers exposed to light and moisture is confirmed by several researchers, such as Bittencourt[12], Wang et al.[26] and De Paoli[27], among other authors. These works describe the fiber surface and the photo degradation process, which gradually reduces its molar mass, mainly due to oxidation and the splitting of the amide group. Another factor that may influence the degradation is the presence of impurities in the fiber due to the weakening of the structure by the formation of free radicals and causing the polymer chains to break up, which characterizes the hydrolytic degradation of the polyaramide fibers[28-33]. The absorption of humidity together with the presence of oxygen and moisture can produce molecular reorganization, from the breakdown of a molecular chain to its oxidation[9,34]. Another factor that may influence the polymer life cycle is related to the effect of stress cracking, which may influence the macroscopic properties of the tissues, such as the weave and warp of the fabric[35,36].

The degradation process may also cause changes in the properties of the polyaramide, such as loss of brightness, changes in color, cracking and the decrease of mechanical properties[26,37]. The blades of armor are not exposed directly to solar radiation, nor direct exposure to weathering, since the blades are arranged in layers, coated by at least two additional layers of different impermeable fabrics[1,38].

In this sense, the objective of this study is to evaluate the loss of the mechanical properties of polyaramide fibers

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

Konarzewski, V. H. C., Spiekemann, F. L., & Santana, R. M. C. used as raw material for the manufacture of ballistic armor with different dates of manufacture and use after exposure to natural weathering.

2. Materials and Methods 2.1 Material The polyaramide fibers came from tissues samples of ballistic armor that was more than 5 years old, that had been disposed of by the logistic center of the Brigada Militar of the State of Rio Grande do Sul. The fiber samples had a dtex of 1670 (2010 - unused vest) and 933 (2005, for a vest that was normally used for five years), referred to as K10 and K05 respectively. The fabrics were trimmed into squares sized 100 mm × 100 mm, and were exposed to natural weathering following ASTM standard[39-41].

The main characteristics of the fibers used for the tensile test are summarized in Table 1 and Figure 1 showing the difference of the weave of the polyaramide fabric samples from 2005 (left) and 2010 (right) fabrics.

2.2 Exposure to natural ageing The specimens tested (100 mm × 100 mm) were exposed to weathering for 12 months (from February 2014 to February 2015), under real climatic conditions in accordance with ASTM 1435-13. Specimens tests were exposed to natural sunlight outdoors in northern direction at 45° inclination as can be seen in Figure 1.

The geographical location of the samples was latitude (30° 05’ South), longitude (51° 11’ West), and altitude (174 m) in the city of Porto Alegre, in the south of Brazil.

This allowed for a normal incidence of solar radiation over the entire surface of the samples. During the natural weathering test it was monitored the average UV radiation index, temperature, and the rainfall that was provided by the CPTED-INPE (Center for weather and climate studies - National Institute for Space Research - Brazil). Figure 2 shows the climatic conditions (maximum Precipitation and Table 1. Characteristics of the polyaramide fibers.

Sample Denier Dtex Warrantee*

K05 840 933 2010

K10 1500 1670 2015

maximum UV index (UVI), moisture and average total monthly rainfall) during the total period studied.

It was observed that the highest temperatures occurred during the months of November 2014 to March 2015, but the maximum temperature in January 2015, which reached over 2014 to January 2015, and between September 2014 to January 2015 the frequency of rainfall was higher, reaching 778 mm values. As a result, the relative humidity was high during this period, between 70 to 90%. Regarding the UVI, the values were high during the spring and summer, reaching values over 12, which is classified as extreme.

2.3 Colorimetric and optic analysis The surface color of the polyaramide samples (every two months) were measured using a Spectro-Guide spectrophotometer, BYK, in accordance with ASTM D2244.

Spectrometer’s Spectro-Guide software transforms spectral data into CIELAB color coordinates (L*, a*, and b*) based on a D65 light source. Lightness (L*) and two chromaticity coordinates (a* and b*) were measured at three different positions on each sample. In addition, the gloss (G) of the samples was measured. An increase in L* value means the color of the sample becomes lighter. A positive a* signifies a color shift towards red, and a negative a* signifies a color shift towards green. A positive b* signifies a shift towards yellow, and a negative b* signifies a shift towards blue.

2.4 Tensile test The fiber tensile test was performed on the INSTRON 3382 universal testing machine based on ASTM D7269M-11[42].

By limiting the length of the fibers to 100 mm, it was opted for the D3822M-14 standard[43]. As can be observed in Figure 3, the fibers were glued on to a cardboard base, with a grammage of 180 g. m-2 and dimensions of 2 × 5 mm, being cast to the center with a diameter of 9 mm. Tensile tests of polyaramide fibers were carried out according to ASTM D3822M-14 in a universal test machine using specimens tested with a length of 100 mm and a load cell of 1000 N, an extension speed of 5 mm min-1, and a gauge length of 25 mm. Each tensile value reported is the average of a 10 specimen test.

*Warrantee given by the armor manufacturers.

Figure 1. Samples from armor fabrics used in this research.

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Figure 2. Climatic precipitation and Ultraviolet Index (UVI) in the city of Porto Alegre (2014-2015).

Polímeros, 29(1), e2019002, 2019

Natural ageing of polyaramide fiber from ballistic armor 2.5 Morphologic analysis, SEM For the scanning electron microscopy (SEM) analysis, the samples were deposited in a carbon type stuck to stub, metalized with gold. For image acquisition, a SEM model JOEL 6060 was used with 2 kV and magnification of 1000 x[44].

2.6 Diameter calculation of the polyaramide fiber For the calculation of the polyaramide fiber diameters a 50 micron scale was used and based on the image it was divided into an area of nine equal divisions containing a significant amount of fibers for analysis, the average was

obtained. Within each area, the amount of fibers “j” were counted, obtaining the same diameter as “fj”. With both of the samples, the average of the area “Zi” was obtained.

The procedure was repeated in the delimited areas, and the mean diameter was calculated. The advantage of this method was the correct division of the zones with a low error in the diameter estimation, with two weights, considerably reducing the error.

3. Results and Discussions 3.1 Colorimetric and optical proprieties Figure 4 shows the results of the colorimetric parameters of the polyaramide K05 and K10 samples before and after exposure to natural weathering. The most significant change in color with a tendency to darken the samples from yellow to a darker shade occurred in the first two months of exposure, with a reddish tendency due to the increase of parameter “+a” to positive values, changes that can be better visualized in Figure 5, with woven samples before and after exposure.

In addition, a reduction of parameter “b*” lower yellowness was observed, because of the increase the redness (+a) of the sample. These color changes indicate chain cleavage due to the hydrolytic degradation of the polyamides.

Figure 3. Specimen test preparation for tensile test.

Such photo degradation features are reported by Bittencourt[12] and Billingham[34], describing the degradation processes of the polyaramide fibers. The change of coloration, as well as the loss of brightness, represented by the decrease of Luminosity (c) (the fiber became more opaque) and increase of fiber roughness. The decrease of the Gloss (d), can be attributed to hydrolysis too, in addition to this, there was environmental stress cracking too, due to the climatic

Figure 4. Colorimetric and optical results of the samples before and after natural ageing: (a) parameter “a”; (b) parameter “b”; (c) Luminosity; and (d) gloss.

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Konarzewski, V. H. C., Spiekemann, F. L., & Santana, R. M. C. conditions to which they were exposed, greater periods of precipitation and temperature variations, which depleted the fibers, being faster in samples k10. The signs of degradation reached the superficial layers of the tissues, causing an initial discoloration, facts reported by Wang et al.[26] and De Paoli[27]. The initial colorimetric results of the degradation of the samples were very similar for the K05 and K10 samples.

However, due to the difference in the fabric weaving and the yarn, after two months of continuous exposure to the

Figure 5. Color changes of the samples before and after natural ageing.

weather, it was observed that the brightness indicators related to the roughness were not uniform. Such signs may indicate a fiber degradation process, which is directly related to the exposure area, as well as the interlacing of the threads.

3.2 Morphological properties, SEM Figure 6 shows the micrographs of the polyaramide fibers of samples K05 and K10 before and after exposure to natural weathering. It was possible to verify that both of them presented a clean and regular surface before exposure, but after four months it was observed that the fibers K 05 presented less aggression on their surface when compared to the fibers K10, this was probably influenced by the type of weave, weight and lower dtex as shown in Table 1.

The degradation of the samples after exposure to natural weathering is related to UV radiation, acid rain and impurities, which under temperatures ranging between 3 to 40 °C altered the fibers over a year. The increase of striations and surface deformations can be attributed to hydrolysis, stress cracking as well as UV radiation, with morphological changes being evident visually similar to the two aramid samples[26,27].

For the initial sample, without exposure to the k05 and K10, no signs of premature wear of the samples were observed.

Such morphology denotes the stability of the fiber, since it is not exposed to light effects[33,36]. These effects were observed with the progressive appearance of striations, roughness, surface defects and incrustations, indicating the progressive degradation of the fiber[36,37]. After four months the morphological changes became more visible, with the appearance of cracks showing a few obvious signs of fiber

Figure 6. MEV from samples K05 and K10 before and after natural ageing.

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Polímeros, 29(1), e2019002, 2019

Natural ageing of polyaramide fiber from ballistic armor degradation, especially in the K10 sample, which despite having a larger dtex, had a more open tissue weave.

Photo degradation was more severe with the more open frame, exposing the samples to light and humidity, increasing photochemical degradation, hydrolysis and stress-cracking[27,30,31]. The variation of fiber diameters during one year, as shown in Figure 7, caused a light increase of the diameter of the aramid fibers of sample K10, when compared to sample k05.

Comparing the sample K10 with K05, from a ballistic vest used over 5 years, it was observed that the diameter of the fibers is very similar in the initial sample “0 months”, not presenting high degradation or absorption of significant humidity over five years of sample difference[12,39]. With the exposure time, the K 05 sample presented a slightly lower humidity absorption over the period of one year, however an increase of the sample diameter occurred, caused by hydrolysis degradation. The initial diameter of the fibers corresponds to that reported by other researchers when characterizing the diameter of the polyaramide fibers 49, around 12 μm, for new fibers[6,20,22]. These results reinforce the hypothesis of fiber stability, even after five years of use (before natural ageing), maintaining similar morphological characteristics as when it was new, thanks to the protective cover of the ballistic panels, avoiding direct exposure to the sun or hydrolysis of the fiber.

Figure 7. Diameter changes of polyaramide fiber after natural ageing.

3.3 Mechanical properties Table 2 shows the average results of the maximum tensile stress of fibers K05 and K10 before and after exposure to natural weathering. Comparing the initial results, the polyaramide fiber presented good stability, even considering the time of use, storage and the fact that the samples k05 and K10 were removed from fabrics, theoretically subjected to mechanical degradation in function of the weaving[11,23,30]. Table 2 also summarizes the mechanical properties of maximum stress until fiber rupture and elongation, in mN/tex, recorded during the tensile test. Considering the difference of origin of the fibers, it was decided to compare the data in mN/tex, and it was possible to verify the reduction of strength force and elasticity over a year[16,26,30]. The hydrolysis, stress-cracking environment and photo degradation mechanisms present on the fibers degradation process were more accentuated in the first 60 days, remaining stable for two months with resistance decline and reduced after 6 months.

As can be seen in Figure 8, the major change occurred after two months, when K10 fiber presented only 52% of the maximum stress until rupture and the sample K05 34%, proving the loss of properties of the polyaramide when submitted to natural ageing.

The initial reduction of resistance of samples k10 and k05 can be attributed to the amount of fibers and the stress‑cracking environment, hydrolysis and photo degradation, since the sample k 10 had a higher amount of fibers when compared to sample K05, which when exposed to the weather were affected from the outside to the inside[31]. From the fourth month of exposure until the sixth month, the sample k 10 had a more pronounced loss of strength, after that the samples started to have a similar resistance decrease, reaching 18% of the original resistance of the yarns in 12 Months of exposure[22,26,38].

Figure 8. Maximum tension from polyaramide by natural ageing (mN/dtex).

Table 2. Mechanical proprieties of polyaramide fibers.

Exposition time (months) 0 2 4 6 8 10 12

Breaking strength (mN/tex) K 2005 K 2010 1898 ± 159.4 1870 ± 95.2 649 ± 134.3 964 ± 180.5 650 ± 57.2 968 ± 118.4 637 ± 151.4 667 ± 85.6 483 ± 115.7 526 ± 31.7 445 ± 88.4 422 ± 54.7 382 ± 63.3 339 ± 23.3

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Elongation at break (%) K 2005 K 2010 6.0 7.0 5.0 5.0 4.0 5.0 3.8 4.0 3.0 3.3 2.0 2.7 3.0 2.3

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Konarzewski, V. H. C., Spiekemann, F. L., & Santana, R. M. C.

4. Conclusions Results of this study showed that new or used polyaramide blankets (before exposure to natural weathering) did not present colorimetric, optical or mechanical changes. However, after exposure to natural weathering for a total period of one year, in the city of Porto Alegre, RS, Brazil, they showed a pronounced color change and a considerable loss in the mechanical performance of the fibers during the first 2 months. Highlighting the largest changes in K10 samples, although this has never been used, but with different weft and dtex from the sample used for 5 years K05a. K05 samples showed higher mechanical performance and lower swelling and less aggression from environmental stress-cracking in the first 2 months. Finally, it is concluded that all samples analyzed after exposure to natural weathering presented hydrolytic degradation, environmental stress-cracking and photochemistry.

5. Acknowledgements The authors would like to thank LAPOL of the Federal University of Rio Grande do Sul and the Brigada Militar of the State of Rio Grande do Sul for the materials and support provided for the development of this study.

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21. Bencomo-Cisneros, J. A., Tejeda-Ochoa, A., García-Estrada, J. A., Herrera-Ramírez, C. A., Hurtado-Macías, A., MartínezSánchez, R., & Herrera-Ramírez, J. M. (2012). Characterization of Kevlar-29 fibers by tensile tests and nanoindentation. Journal of Alloys and Compounds, 536(Suppl 1), S456-S459. http:// dx.doi.org/10.1016/j.jallcom.2011.11.031.

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26. Wang, H., Xie, H., Hu, Z., Wu, D., & Chen, P. (2012). The influence of UV radiation and moisture on the mechanical properties and micro-structure of single Kevlar fibre using optical methods. Polymer Degradation & Stability, 97(9), 1755-1761. http://dx.doi.org/10.1016/j.polymdegradstab.2012.06.010.

27. De Paoli, M. A. (2008). Degradação e estabilização de polímeros. São Paulo: Artliber.

28. Santos, A. S. F., Agnelli, J. A. M., Trevisan, D. W., & Manrich, S. (2002). Degradation and stabilization of polyolefin from municipal plastics waste during multiple extrusions under different eprocessing conditions. Polymer Degradation & Stability, 77(3), 441-447. http://dx.doi.org/10.1016/S01413910(02)00101-5.

29. Bertin, D. M., Larissa, S., Catto, A. L., Camargo, M. M. F., Chiellini, E., Corti, A., Morelli, A., & Campomanes, R. M. S. (2010). Polypropylene degradation: theoretical and experimental investigation. Polymer Degradation & Stability, 95(5), 186-192. http://dx.doi.org/10.1016/j.polymdegradstab.2010.02.006.

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carbon/Kevlar/epoxy hybrid composite for orthopaedic usage and its influence on mechanical performance. Composites.

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Received: Sept. 08, 2017 Revised: May 24, 2018 Accepted: May 25, 2018

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ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.02018

Application of ashes as filling in reprocessed polypropylene: thermomechanical properties of composites Lina Marcela Crespo1 and Carolina Caicedo1,2*  Grupo de Investigación en Desarrollo de Materiales y Productos – GIDEMP, Centro Nacional de Asistencia Técnica a la Industria – ASTIN, Servicio Nacional de Aprendizaje – SENA, Cali, Valle del Cauca, Colombia 2 Consejo Nacional de Ciencia y Tecnología – CONACYT, Consorcio de Investigación y de Innovación del Estado de Tlaxcala, Centro de Investigación en Química Aplicada, Saltillo, Coahuila, México 1

*carolina.caicedo@conacyt.mx

Abstract The life cycle of a product depends to a great extent on its reuse and ease of recycling. This work had developed of composite materials of reprocessed polypropylene composites with rice husk ash (RHA) and sugarcane bagasse ash (SBA) through the coextrusion and injection processes as main purpose. The polymeric matrix was reprocesed until six generations by the injection technique. The reprocessed PP was mixed in 80:20 proportions with respect to filler mineral, using maleic anhydride as coupling agent in a coextrusion machine. The new series of composite materials were analyzed thermal, mechanical, rheological and morphologically. The incorporation of ashes in the PP matrix achieved characteristics of improved tensile strength, conserving the thermal properties. For this reason, this work presents an alternative for the manufacture of composite materials from post-industrial waste.

Keywords: rheological analysis, rice husk ash, sugarcane bagasse ash, mechanical properties, thermal properties.

How to cite: Crespo, L. M., & Caicedo, C. (2019). Application of ashes as filling in reprocessed polypropylene: thermomechanical properties of composites. Polímeros: Ciência e Tecnologia, 29(1), e2019003. https://doi. org/10.1590/0104-1428.02018

1. Introduction In the last decade, researchers have been interested in reuse of organic and inorganic materials to be incorporated as filler in polymer matrices controlling aspects of shape, concentration and interface which get to contribute efficiently to the improvement of physicmechanical properties of the resulting composite material[1,2]. This practice is industrially known because of the reduction of cost in production.

In general, the addition of filler mineral on the polymers decreases the impact energy, however, a weak interaction contributes to the fragmentation of it, due to the irregular flow of stresses which it is produced[3]. In fact, most studies on modification of semicrystalline polymers with rigid particles indicate a significant loss of tenacity in comparison with pure polymer[4]. In the development of composite materials, it has been found that isotactic polypropylene (PP) has been used in investigations, despite the inconvenience related to low impact resistance[5-7]. In the case of PP copolymers, the formation of covalent bonds between the PP matrix and the elastomer phase (rich in ethylene) avoiding the separation of the macroscopic phases. This elastomeric microdispersion phase can modify the mechanisms of plastic deformation of the PP matrix in such a way as to obtain high impact resistance[8].

Rigid particles must be separated and create a free volume in the blend in a submicronic size level, which

Polímeros, 29(1), e2019003, 2019

it is explained with the mechanism of cavitation in the hardened rubber systems[9]. Through time, the fillers most commonly used in polypropylene have been calcium carbonate and talc[10-12], however, currently it continues being studied alternatives for the use of inorganic fillers from physic‑chemical, thermal or biological degradation of agricultural residues as the case of ashes[13]. Rice husk is one of the main products of agricultural waste, according to statistics of FAO, rice global production was approximately of 756.7 millions tons in 2017, moreover, a production of 989.96 tons were reported by DANE in Colombia in the first half of 2017[14]. Therefore, the rice husk is a derivative of great interest for researchers and producers due to the challenge it means the offering alternative materials which generate added value to this residue.

The literature reports a series of studies about polymer composites and rice husk[15,16], one of the favorite fillings after palm oil and rubber, as well as; the product of incineration: Rice husk ash (RHA) compound mainly by silica between 87-89%[1,17,18]. Some results show an 18% increase in mechanical strength and elongation to break of high density polyetilene (HDPE) composites with RHA incorporated at 1.5% in weight in presence of maleic anhydride as coupling agent[19]. In other research, it was obtained the improvement of the stress modulus in composites of white rice husk ash

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

Crespo, L. M., & Caicedo, C. (WRHA) with polypropylene (PP)/natural rubber, however, tensile strength, elongation to break decreased with increasing the filler of WRHA, while the incorporation of the coupling agent (Silane, 2-aminopropyl triethoxysilane), reduced the amount of absorbed water by the composite material[20].

On the other hand, some proposals have been presented to redirect the use of sugar cane industry waste through the implementation of husk ash as reinforcement in rubber, and the results are promising, an increase of 300% in the stress modulus was presented[21]. Therefore, in this work we proposed obtaining two series of composite materials from recovered PP copolymer random with fillers mineral of Rice Husk Ash (RHA) and Sugarcane Bagasse Ash (SBA), due to the above, to know the influence that presented by incorporation of particulate inorganic material in the different generations of the PP on the physic properties of the PPnRHA y PPnSBA composite material.

2. Materials and Methods 2.1 Materials Polypropylene copolymer random was used (PP) 02R01CA-1 produced by Propilco with a melt flow index (MFI) equal to 1.6 dg/10 min (230 °C, 2.16 kg), the coupling agent used was modified polypropylene with licocene PP‑g-MA 7452 from Clariant, presenting an addition of 7% maleic anhydride, 156 °C melting point, 0.91 g/cm3 density and high crystallinity. The ashes of rice husk and sugarcane bagasse ash were from a rice mill and a sugar mill in Valle del Cauca, respectively. The granulometric characteristics are presented in Table 1.

2.2 Preparation of composites Multiple injection (six successive) were carried out either on DEMAG injection machine 150 tons model 1991, hydraulic and insertable valves. The temperature profile Table 1. Granulometric characteristics of rice husk ash (RHA) and sugarcane bagasse ash (SBA).

Types of Ashes #Sieve 50 100 140

RHA Retained (%) 0.00 0.10 1.20

SBA Retained (%) 0.29 9.87 36.99

200 325 400 Fondo Average particle size (µm)

11.43 53.06 18.86 15.35 63±15

40.38 12.09 0.07 0.29 106±37.5

was 190 °C - 195 °C - 195 °C and 200 °C for the nozzle.

The injection pressure was kept constant at 75.6 b, The mold was kept at 45 °C by a water cooling system and a constant injection speed of 45 cm3.s-1 was applied. The injection parameters were kept constant in each cycle. The specimens obtained by injection were type 1b “dogbone” according to ISO 527-2 (step 2 or 3). Then, a part of the test pieces was passed through a Rotrogram Mold-tek blade mill at 1745 rpm, where granules with a diameter of ~8 mm were obtained. The processing methodology is summarized in Figure 1 and the successive stages are detailed.

The composite materials were mixed in a double-screw extruder model Haake Rheomex OS PTW in co-rotating parallel configuration. The temperature profile with gradual increase of 5 °C from 155 °C to 200 °C, at the end of the screws, divided into 10 cylinder zones, each zone covering a length of 4D. Finally, the composites were injected and molded under the same conditions as the material reprocessing. The samples were differentiated for the next symbol PPnRHA and PPnSBA, where PP indicates the used matrix of polypropylene, RHA and SBA the incorporated mineral fillers (RHA: rice husk ash, SBA: sugarcane bagasse ash), n subscript corresponds to reprocessing cycle number.

2.3 Thermogravimetric Analysis (TGA) and Differential Scanning Calorimetry (DSC) The thermal properties were determined in a thermogravimetric TGA/DSC 2 STAR System, Mettler Toledo. The samples (10±0.5 mg) were put in alumina crucibles in a temperature range between 25 °C to ~500 °C under nitrogen atmosphere (50 cm3/min) with a heating rate of 10 °C/min. It was worked out according to the ASTM E1131-98 and ASTM D3418-12 standards, respectively.

2.4 Tensile and flexural properties The mechanical property measures were done in a universal Goodbrand machine according with the ASTM D638 standard, with a test speed of 50 mm/min and a load cell of 500 kgf. The enlargement values were determined (by extension of the jaws) to the tearing and tensions resistance.

The environmental chamber Dies, a Baker caliper gauge and an Oakton thermohygrometer. The flexural tests were performed in Goodbrand machine according to the ASTM D790-10 standard, with a test speed of 5 mm/min and a 50 kgf cell. Modules, resistances and elongation percentages were obtained. Five samples were analyzed for every composites mix and the average values are presented.

Figure 1. Stages for process of obtaining composites and composite material development.

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Polímeros, 29(1), e2019003, 2019

Application of ashes as filling in reprocessed polypropylene: thermomechanical properties of composites 2.5 Microstructural analysis The development of compounds were analyzed through scanning electron microscopy (SEM). Usign an SEM JEOL de Mesa JCM 50000. For the filler particles analysis within the composites, cross-section samples of composites were prepared through the same methodology. They were obtained from the middle part of the dog-bone samples. Then, they were submitted to a gold coverage using PVD, in a high vacuum and a 5 kV voltage was used. The magnifications were x44 to determine the dispersion, x200 to asses the general morphology of the particles and in order to observe the ash/matrix interface. It is worth to mention, that for every analysis the sample were conditioned to 25 °C in an environment of relative humidity of 50±5% during 48 h.

2.6 Rheological analysis The rheological measurements were performed at 190 °C and the shear rate was in the range of 0.001 to 300 s-1.

The dynamo-mechanical tests settings were: the percentage of deformation between 1 and 10%, with a frequency range of 0.1 to 628.10 rad/s, to determine the storage modulus and the loss modulus.

The results of the tensile test for the sample of PPnRHA and PPnSBA composite samples are presented in Table 4, where the conservation of the properties is appreciated, which is attributed to the good dispersion of the particles within the matrix by the presence of the coupling agent (PP‑g-MA) which in turn it facilitates the transfer of properties between both materials, allowing the ash particles which present high stiffness to generate difficulty in the deformation of composites, making them less ductile[24].

This way, the addition of ashes to PP lead to the increses of stiffness with values over to 50% in composite materials with respect to the reprocessed starting materials (see Table 3)[25].

The stiffness in the PPnSBA was more notorious in a 25% than the PPnRHA. Table 5 shows an increase (~10%) for the values of maximum stress obtained by the flexural test for PPnRHA and PPnSBA in relation to the PP resin. On the other hand, Table 2. Result of TGA and DSC obtained for the processed polypropylene and composed materials of RHA and SBA.

3. Results and Discussions 3.1 Thermal analysis The results obtained from the curves of thermogravimetric analysis (TGA) and differencial scanning calorimetry (DSC) are shown in the Table 2, which indicates in every cases that initial degradation for generations 1, 3 and 6 of PP and developed composite materials (PPnRHA and PPnSBA) are fewer than required temperature to transform them under the process of injection and extrusion, without exceeding 200 °C, as a result, the integrity of the reference reprocessed materials and the blend of composite material is guaranteed. In addition, the temperatures and fusion enthalpy were evaluated.

In the Table 2 it is shown that fusion temperatures of composite materials PPnRHA y PPnSBA remained in values of 147 °C; this in comparison with the two departure reprocessed materials of PP. The temperatures of initial degradation presente a more pronounced increase of PPnSBA, which it evinces the influence of ashes in absortion of the heat generated in the degradation processes, producing a deceleration of itself[22,23]. This behavior is also attributed to a set of characteristics produced by the presence of ashes in composite materials of the polymeric matrix, within them it is possible to find the effect that inorganic ashes have as nucleating agent, causing a foretaste in the beginning of crystallization of PP. Besides, the limitation in mobility of the polymeric chains which get to delay the volatilization of the generated products in the temperature in which occurs the break of the carbon-carbon junctions[18,22].

Ti

Temperatures Tmax

Tm

PP1

239.4

436.2

149.1

PP3

231.2

428.9

147.5

PP6

229.0

429.8

146.5

PP1RHA

219.5

423.1

145.4

PP3RHA

229.1

390.3

145.8

PP6RHA

229.9

415.5

146.7

PP1SBA

233.2

413.1

143.7

PP3SBA

241.4

428.9

146.1

PP6SBA

249.0

456.5

146.7

Sample

Ti = Temperature of initial degradation; T max = Temperature of maximum degradation; Tm = Melting temperature.

Table 3. Characterization of the mechanical properties of the matrix of reprocessed polypropylene.

Sample PP

Stress Fracture stress (MPa)

Ultimate stress (MPa)

PP1

25.46±0.21 16.59±0.23 85.04±3.58

PP3

23.95±0.51 15.95±0.23 132.75±1.34 0.97±0.04 2.45±0.21

PP6

26.74±0.38 17.47±0.24 130.12±0.41 1.05±0.01 2.73±0.33

Table 4. Results of the tests of stress of the composite materials with ashes.

Ultimate stress (MPa)

PP1RHA PP3RHA

3.2.1 Flexural and stress resistance In Table 3 can be found the results of maximum resistance obtained for the reprocessed materials until the sixth cycle, which refer to a stable behavior with an average value of 25.38 ± 1.40 MPa. Polímeros, 29(1), e2019003, 2019

0.99±0.04 2.16±0.12

MPa = Megapascals.

Sample

3.2 Mechanical properties

Strain (%)

Flexural Ultimate Fracture stress stress (MPa) (MPa)

Fracture

Strain (%)

22.66±0.71

Stress (MPa) 15.88±0.50

46.86±3.99

24.07±0.10

15.90±0.44

48.82±1.24

PP6RHA

24.91±0.04

16.23±0.33

49.97±2.45

PP1SBA

25.28±0.71

21.39±1.00

19.29±0.75

PP3SBA

25.03±0.34

23.25±0.76

18.20±0.84

PP6SBA

26.342±0.09

23.16±0.30

18.47±0.67

MPa = Megapascals.

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Crespo, L. M., & Caicedo, C. the modulus of elasticity for the compounds with the highest number of reprocesses decreases by 50%. These results present a relation with the stress results previously explained.

3.3 Microstructural analysis The SEM micrographs of the fracture surfaces of the polypropylene compounds with 20% by weight of the ashes are shown in Figure 2. It is clearly seen in these images that the ash particles were detached from the matrix (PP) in Table 5. Results of the flexural tests of the composite materials with ashes.

PP1RHA

Ultimate stress (MPa) 1.15±0.04

Fracture Stress (MPa) 2.45±0.21

PP3RHA

1.10±0.02

1.91±0.25

PP6RHA

1.15±0.03

1.70±0.22

PP1SBA

1.21±0.02

1.71±0.20

PP3SBA

1.24±0.04

1.40±0.18

PP6SBA

1.22±0.03

1.27±0.15

Sample

MPa = Megapascals.

the material which underwent ductile tearing around these particles. When comparing the images of Figure 2b and 2c, it can be seen that the PPnRHA composites exhibited a lower degree of ductile tearing of the filler present in the matrix, due to the presence of PP fibrils. The above is in agreement with the significant increase in the maximum effort found for PPnSBA compounds. It should be mentioned that Figure 2c (PP6SBA) shows particles with a cenosphere morphology of fly ash[26].

The Figure 3 shows a view furthest from the compounds that allows observing an increase in the concentration of fillers is greater in the center of the specimen, which can be evidenced by the presence of pores and agglomerates that with the increase of generations is less noticeable due to the easy dispersion of filler at low viscosity. These are responsible for facilitating the rupture at the time of stressing the material. The holes with measured lengths of 0.65 to 0.85 mm denote the presence of agglomerates.

These are observed more exposed in the PPnSBA resulting in rougher surfaces; while for PPnRHA a behavior similar to the matrix is shown when generating short fibers.

Figure 2. Micrographs obtained by SEM of PPnRHA and PPnSBA compounds at 200x magnification.

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Polímeros, 29(1), e2019003, 2019

Application of ashes as filling in reprocessed polypropylene: thermomechanical properties of composites

Figure 3. Micrographs obtained by SEM of PPnRHA and PPnSBA compounds at 44x magnification.

3.4 Rheological analysis The complex viscosity values found for the compounds decrease as the matrix increases the number of reprocesses (Figure 4), this behavior is due to the degradation suffered by the polymer due to the temperature and mechanical stress demonstrated in a previous study[27]. In general, the fillers generate a strong shearing thinning effect, this dependence against frequency is less noticeable for the pure polymer.

It is interesting to observe the influence of the type of charge in the viscosity curves at low frequencies, in this case, rice husk ash has a lower dispersion compared to sugarcane bagasse ash, keeping the trend of the matrix which in turn reflects a less stable behavior.

On the other hand, the curves corresponding to the storage and loss modules were analyzed as a function of the angular frequency for the PP and the composites with ashes, which are presented in Figure 5. These exhibit a directly proportional increase of the modules with respect to the angular frequency, similar behavior among the samples developed.

The Figure 5a shows in the frequency range studied, that the material between 10 and 51.05 rad/s, has a viscous behavior as liquids, but at 51.05 rad/s, the viscous to elastic Polímeros, 29(1), e2019003, 2019

Figure 4. Complex viscosity of PP, PPnRHA and PPnSBA.

transition occurs, since G’ > G” (behaves like a solid), this behavior is typical of thermoplastic materials however for this material it is mainly elastic. In the Figure 5b, the PPRHA sample has the highest elastic behavior, since the transition 5/7

Crespo, L. M., & Caicedo, C.

Figure 5. Storage and loss modules of (a) PP; (b) PPnRHA; and (c) PPnSBA.

from viscous to elastic (G’ > G”) occurs at 21.24 rad/s, the sample PP3RHA, the sample PP3RHA, follows in this order, since the viscous to elastic transition occurs at 25.51 rad/s.

The sample PP6RHA has a higher viscous behavior, since the angular frequency at which G’ > G” has a value of 69.37 rad/s. These results denote a viscoelastic behavior.

Finally, the PP1SBA sample of Figure 5c showed the highest elastic behavior, since the viscous to elastic transition was given at an angular frequency of 22.57 rad/s, in the case of the samples PP3SBA and PP6SBA the angular frequency values​​ a which occurred the transition were 37.53 and 25.08 rad/s, respectively. This result indicates that the PP3SBA sample has the highest viscous behavior. In addition, the reprocessing of the material increases its fluidity, which may be due to a better dispersion or orientation of the reinforcer. An important aspect is that the sample reprocessed six times (PP6SBA), showed lower viscous behavior than the one processed 3 times (PP3SBA), due to a greater interaction of the reinforcing material with the PP than the PP3SBA sample.

In addition, according to the results, no direct relationship of the viscoelastic behavior with the reprocessing number is observed. On the other hand, the polydispersity indices for the compounds were estimated with respect to the PP, which presents a value of 3.1 and values between 3.9 and 5.1 were found for PP6RHA and PP3SBA, respectively. The above is indicative of the loss of crystallinity in these compounds.

4. Conclusions The results obtained show that the use of post-industrial ashes from the production of rice and cane are promising for the development of composite materials, as well as, the use of recovered polymeric material (PP). The developed composite materials showed a stable behavior for the melting temperatures. The degradation temperatures showed more pronounced increases for the PPnSBA series. In the same way, in the mechanical results, the rigidity of the PPnSBA composite materials was favored about 25% more than the PPnRHA series. However, the filler of ash had an effect on the stiffness of the composite materials, this occurs as a consequence of the suitable dispersion generated by the coupling agent (PP-g-MA), which facilitated the transfer of the matrix to the fillers.

The images captured by SEM showed important characteristics that complemented the mechanical analysis, 6/7

with this it is possible to conclude that the rice husk ash favor the modulus of elasticity with respect to those obtained with sugarcane bagasse ash with the presence of fibrils. The rupture caused in the tension test was favored by agglomerates of the filler and cenosphere particles for the case of the PPnSBA compounds. The process used was adequate to disperse the reinforcing material, in the enlarged images the presence of pores and particles was observed homogeneously. In the rheological analysis, the angular frequency range that corresponds to the viscous to elastic behavior studied, it was found that the PP object of study is mainly elastic, while PPnRHA and PPnSBA are viscoelastic.

The reprocessing of the materials produced a greater fluidity of the mixtures, because the elastic behavior of the solids decrease with the reprocessing of the material and incorporation of the ashes, this possibly is due to a greater dispersion of the reinforcement or a smaller interaction of this with the matrix. In general, the storage module exceeded the loss module at lower frequencies for the compounds (between 20 and 30 rad/s) compared to the PP (55 rad/s), which stand out the high viscosity of these compounds. The PP6RHA sample, despite revealing elasticity conditions higher than 70 rad/s, presents modules 25% lower than the PP.

5. Acknowledgements The autors gratefully acknowledges the ASTIN-SENA, Research Group in Development of Materials and Products (GIDEMP by its Spanish acronym), Bio- y Nano-technology line of Tecnoparque in 2017. Also, we thank the SENNOVA (projects SGPS-2195-2017) for financial support. C. C. acknowledges the economic support of the Program of Cátedras CONACYT from CONACyT-CIQA (2018-2028).

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Received: Mar. 14, 2018 Revised: May 23, 2018 Accepted: May 28, 2018 7/7

ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.06617

Coating of urea granules by in situ polymerization in fluidized bed reactors Bruno Souza Fernandes1, José Carlos Pinto2, Elaine Christine de Magalhães Cabral-Albuquerque3 and Rosana Lopes Lima Fialho3*  Centro de Ciência e Tecnologia em Energia e Sustentabilidade, Universidade Federal do Recôncavo da Bahia – UFRB, Feira de Santana, BA, Brasil 2 Laboratório de Engenharia de Polimerização, Programa de Engenharia Química, Instituto Alberto Luiz Coimbra de Pós-graduação e Pesquisa de Engenharia – COPPE, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brasil 3 Laboratório de Polímeros e Bioprocessos, Programa de Pós-graduação em Engenharia Industrial, Universidade Federal da Bahia – UFBA, Salvador, BA, Brasil 1

*rosanafialho@ufba.br

Abstract The main objective of the present work is to produce and characterize urea granules coated with polymers prepared with aqueous solutions of acrylic acid and glycerol. Both coating and drying of urea granules were performed in a fluidized bed reactor. Fourier transform infrared spectroscopy analyses indicated the presence of poly(acrylic acid) and acrylic acid / glycerol copolymers on the granule coating and the formation of chemical bonds between urea and the polymer coating. Scanning electron microscopy images showed that the original and coated urea granules presented different characteristics, reinforcing the idea that coating occurs in the fluidized bed. Finally, rates of urea release showed that the coated granules presented slightly slower rates of urea dissolution in water due to the presence of the coating layer.

Therefore, it is shown that it is possible to produce coated urea granules through in-situ polymerization onto the granule surface using a fluidized bed.

Keywords: acrylic acid, coating, fluidized bed, glycerol, urea.

How to cite: Fernandes, B. S., Pinto, J. C., Cabral-Albuquerque, E. C. M., & Fialho, R. L. L. (2019). Coating of urea granules by in situ polymerization in fluidized bed reactors. Polímeros: Ciência e Tecnologia, 29(1), e2019004. https:// doi.org/10.1590/0104-1428.06617

1. Introduction Urea is the fertilizer used most often in agriculture; for this reason, the worldwide production and consumption of urea has steadily increased in recent years. Urea is characterized by its high nitrogen content (46 wt%), low production cost and high water solubility. In addition, urea is noncorrosive and can be easily mixed with other compounds[1-3]. The main problem associated with the use of urea as a fertilizer is the high rate of loss to the environment through leaching and volatilization[3-5]. Losses can reach 50 wt% of the applied urea fertilizer, depending on the climate, soil conditions and application technologies, causing environmental pollution and increasing the costs of crop production[6-10].

A possible alternative to reduce nutrient losses is the development of slow-release fertilizers by coating urea granules with materials that feature lower water solubility[5,9,10] or by using materials that allow for the slow release of urea[8,11-14].

Particularly, the production of slow-release urea products with help of polymer coatings constitutes a promising technological solution for many applications[7,9,15].

The use of poly(acrylic acid), PAA, obtained through free-radical polymerization of acrylic acid in aqueous medium

Polímeros, 29(1), e2019004, 2019

presents a number of comparative advantages in various applications, including the low cost, biodegradability and good biocompatibility. PAA-based resins can also exhibit high capacity of water absorption and retention and can be possibly used for production of superabsorbent coatings and manufacture of slow-release fertilizers[16-22]. It is also important to observe that glycerol can act as a chain transfer agent, esterification and/or a crosslinking agent in some free‑radical and functional polymerizations, and particularly in aqueous free-radical acrylic acid polymerizations, leading to formation of chain branches and modifying the molecular weight distributions of the final products[23,24]. For this reason, glycerol has been used frequently as a comonomer in aqueous free-radical acrylic acid polymerizations.

Based on the previous remarks, the main objective of the present study was to produce coated urea granules in a fluidized bed reactor, using polymer materials produced in‑situ and ex-situ through aqueous free-radical polymerizations of acrylic acid and glycerol. The coated and uncoated urea granules were characterized by Fourier transform infrared spectroscopy (FTIR), scanning electron microscopy (SEM)

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

Fernandes, B. S., Pinto, J. C., Cabral-Albuquerque, E. C. M., & Fialho, R. L. L. and rates of urea release in distilled water. The obtained results indicated that polymer coating was performed successfully in all analyzed cases, confirming that in-situ polymerization of the monomer solution on the urea granules is possible in fluidized bed reactors. Particularly, FTIR analyses indicated the presence of poly(acrylic acid) and acrylic acid / glycerol copolymers on the granule coating and the formation of chemical bonds between urea and the polymer coating.

SEM images showed that the original and coated urea granules presented different characteristics, reinforcing the idea that coating occurs in the fluidized bed. Finally, rates of urea release showed that the coated granules presented slightly slower rates of urea dissolution in water due to the presence of the coating layer.

2. Materials and Methods 2.1 Materials Urea (CO(NH2)2, with minimum purity of 99.5 wt%, containing 46.4 wt% of nitrogen, with Mw of 60.07 g.mol-1) was provided in the form of granules (average particle size of 1.84 mm) by Petrobras (Camaçari, Bahia, Brazil).

Acrylic acid (C3H4O2, with minimum purity of 99.5 wt% and Mw of 72.02 g.mol-1), glycerol (C3H8O3, with minimum purity of 99.5 wt% and Mw of 92.1 g.mol-1) and potassium persulfate (K2S2O8, with minimum purity of 99.0 wt% and Mw of 270.32 g.mol-1) were acquired from Vetec (Duque de Caxias, Brazil). The enzyme urease and other reagents used to hydrolyze urea and determine the urea content of analyzed samples were provided by Doles (Goiânia, Brazil) as the enzymatic kit Urea 500. All reagents and solvents used for polymer characterization were purchased at analytical grades from Vetec (Duque de Caxias, Brazil). All chemicals were used as received.

2.2 Coating of urea granules The coating of urea granules was conducted in two ways. First, a copolymer was produced prior to coating in the fluidized bed. In the second, copolymerization and coating occurred simultaneously in the fluidized bed.

Detailing the first way, the copolymerization reactions were performed in an Atlas Sodium reactor (Syrris, United Kingdom). Copolymerization reactions were conducted in distilled water by mixing 100 g of the aqueous initiator solution and 200 g of the aqueous monomer solution.

The final reacting mixture contained 3.00 g of potassium persulfate, 39.80 g of urea, 23.86 g of acrylic acid and 3.00 g of glycerol. Reactions were performed in a 500-mL glass flask, with mechanical stirring of 300 rpm, temperature of 80 °C and time of 2 h, in accordance with previous experimental studies[18,25-28]. As the final copolymer medium was biphasic (a lighter and less viscous liquid phase and a heavier viscous liquid phase, which contained the copolymer), continuous stirring of the copolymer feed was necessary to prevent heterogeneous and varying feed conditions.

Detailing the second way, the monomer solution used to perform the coating of urea granules through in-situ copolymerization inside the fluidized bed reactor contained 3.00 g of potassium persulfate, 63.66 g of acrylic acid, 3.00 g of glycerol and 230.34 g of distilled water.

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Coating and drying of the urea granules were performed in a Midi Model Fluidized Bed Reactor (Glatt, Germany) with height of 900 mm, width of 700 mm and length of 740 mm. The coating process in the fluidized bed reactor was conducted with 300 g of urea granules and 100 mL of the respective coating solutions. In all cases, the fluidization pressure was kept constant and equal to 0.8 bar. The atomization pressure was equal to 0.5 bar and 1.0 bar, when the monomer and polymer solutions were used for coating, respectively.

The feed flow rate of the coating solution was constant and equal to 5 mL/min. The temperature of the fluidizing air was equal to 80 °C. The blow down rate was 1 every 4 s.

The height of the cylindrical partition, which was parallel with the largest mouth of the bed, was 6.5 cm.

The coating process was conducted with intermittent interruption of the feed flow rate of the coating solution in order to prevent particle agglomeration, as described elsewhere[29-31]. Usually, feed flow rates were kept constant for periods of 20 s and stopped for periods of 1 min and 40 s.

Coating operations were performed as presented in Table 1.

2.3 Fourier Transform Infrared Spectroscopy (FTIR) Coated and uncoated samples of urea granules were analyzed by FTIR using a Vertex 70 Spectrophotometer (Bruker, United States), equipped with Attenuated Total Reflectance (ATR) probe and operating in the range of 4000-600 cm-1. Spectral data were reported as averages of 50 readings obtained with resolution of 2 cm-1 at 16 ± 1 °C. The backgrounds were collected using uncoated urea granules in order to reduce the effect of the urea signal.

2.4 Scanning Electron Microscopy (SEM) Coated and uncoated samples of urea granules were also evaluated by SEM. Granules were placed onto double-sided carbon tapes and affixed to gold-coated aluminum carriers.

A Quanta 400 model SEM Microscope (FEI Company, United States), with maximum operation voltage of 30 kV and nominal resolution of 1.2 nm in high vacuum and equipped with a SE (secondary electron) detector, was used.

In all cases, the voltage was set at 20 kV and the images were acquired with the SE detector.

2.5 Urea release in water Analyses of urea release profiles were performed in distilled water in order to compare the relative performances of coated and uncoated urea granules. Coated and uncoated urea samples were weighed to provide 0.80 mg of urea per mL of water. Tests were performed in 25 mL of distilled water at 30 °C. The aqueous solution containing the sample was kept under continuous magnetic stirring of 100 rpm.

10-μL samples of the aqueous solutions were withdrawn for analyses at time intervals of 1, 5, 10, 20, 30, 60 and 90 min.

Table 1. Coatings performed in the fluidized bed reactor.

Coating 1 2 3 4

Solution Monomer Monomer Polymer Polymer

Pauses in solution feed No Yes No Yes

Polímeros, 29(1), e2019004, 2019

Coating of urea granules by in situ polymerization in fluidized bed reactors The total volume of the system was kept constant through addition of distilled water immediately after sampling.

The urea concentration was determined enzymatically with help of the Urea 500 kit. Ammonia concentrations (and therefore, concentrations of urea) were obtained at room temperature with help of a calibration model, using a UV-visible spectrophotometer (Lambda 35 model, Perkin Elmer, United States) operating in the wavelength range of 570 to 720 nm. The wavelength considered for calculation and determination of the released urea content was equal to 600 nm.

3. Results and Discussions 3.1 Fourier Transform Infrared Spectroscopy (FTIR) Figure 1 shows the FTIR spectra of coated and uncoated urea granules. One can observe the characteristic urea peaks at 3430-3336 cm-1, which correspond to the symmetrical stretching of NH. The peak placed at 1675 cm-1 corresponds to the stretching of the double bond C=O, while the peak positioned at 1590 cm-1 is related to the stretching of bonds in NH or NH2. The peaks placed at 1460 cm-1 and 1003 cm-1 correspond to the shortening of the CN bond, while the peak

positioned at 1149 cm-1 is due to the symmetrical stretching of NH. The peak placed at 787 cm-1 corresponds to the out‑of-phase bending of OCNN, while the peak placed at 714 cm-1 corresponds to the bending of the NH bond. These results were reported in previous studies[32-34].

With regard to the coated urea granules, it is worth noting that all spectra showed the characteristic NH peaks of urea (near 3430-3336 cm-1), although with lower intensity.

This indicates the existence of additional materials, changes of the chemical structure of the urea and new bonds of the NH group of urea with acrylic acid. The appearance of the peaks at 3472 cm-1 and 3225 cm-1 and the analysis of some peaks in the range of 1700-1100 cm-1 justify this statement.

Coating 1 presented a peak of high intensity at 1614 cm-1, revealing the presence of the bond C=C of acrylic acid and the appearance of low-intensity peaks at 1403 cm-1, 1337 cm-1, 1260 cm-1 and 1112 cm-1. These peaks indicate the presence of acrylic acid and formation of PAA and acrylic acid / glycerol copolymers, indicating the in-situ polymerization of the monomer solution and the physical interaction of polymer materials with the urea granules.

Coating 2 presented two new peaks in the characteristic NH spectral region of urea. In addition to the peaks placed at 3421 cm-1, 3329 cm-1 and 3252 cm-1, peaks positioned at 3472 cm-1 and 3225 cm-1 showed that the NH group of urea interacted with the polymer coating. The peak placed at 1693 cm-1 showed the formation of the group CONHR between the urea granules and the carboxylic group of the monomer. The peak placed at 1159 cm-1 indicated the formation of the COC group, through reaction of acrylic acid and glycerol. Therefore, it seems clear that both in‑situ polymerization and chemical interactions between the polymer coating and the urea granules took place in the fluidized bed.

The FTIR spectrum of Coating 3 was similar to the FTIR spectrum of Coating 1. As the coating operation was performed with the polymer solution in this case supports the idea that in-situ polymerizations did occur during preparation of Coatings 1 and 2. The peak placed at 1160 cm-1, which was also present in the first two coatings, indicates the existence of COC groups due to reaction between acrylic acid and glycerol. Furthermore, the peak positioned at 1620 cm-1 indicates the formation of NH3+ and HCOO- groups, formed through reaction of carboxylic groups of acrylic acid and amino groups of urea during the polymerization reaction or coating.

Coating 4 also showed the characteristic peaks that reveal the presence of the copolymer on the urea granules (1612 cm-1, 1335 cm-1, 1232 cm-1). In particular, the peak placed at 1684 cm-1 may indicate the formation of a CONHR group between the urea and the copolymer, due to reaction between the carboxylic groups of acrylic acid and amino groups of urea during the polymerization reaction or coating.

3.2 Scanning Electron Microscopy (SEM)

Figure 1. FTIR spectra of the uncoated and coated urea granules.

Polímeros, 29(1), e2019004, 2019

Figures 2, 3, 4, 5 and 6 show SEM images of uncoated and coated urea granules with various levels of magnification in order to provide better visualization and allow for proper comparison of the images.

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Fernandes, B. S., Pinto, J. C., Cabral-Albuquerque, E. C. M., & Fialho, R. L. L. As one can observe in Figure 2, urea granules present spherical geometry, crystalline structure and irregular surfaces, with large pores. When the magnification reaches 5,000× and 10,000×, the presence of canes can be observed in several directions on the surface of the urea granules.

Costa et al.[35] described the surface of the urea as having a uniform and wrinkled appearance.

Samples prepared through Coating 1 featured fairly regular surfaces with small pores. First, this indicates the presence of the coating on the surfaces of the urea granules.

Figure 2. Images of uncoated urea obtained by SEM at magnifications of (a) 50×; (b) 100×; (c) 500×; (d) 1,000×; (e) 5,000×; and (f) 10,000×.

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Coating of urea granules by in situ polymerization in fluidized bed reactors The characteristic canes of urea granules disappeared, indicating good level of particle finishing. Furthermore, the high regularity and low roughness of the surface possibly indicate that the monomers were converted to copolymers during coating.

Samples prepared through Coating 3, which used the polymer solution, presented a rather uneven and rougher surface and also exhibited a darker region, where coating could not be applied or was applied to a lesser extent.

These characteristics of the coating were probably due to

Figure 3. Images of coating 1 (monomeric solution; without pauses) obtained by SEM at magnifications of (a) 50×; (b) 100×; (c) 500×; (d) 1,000×; (e) 5,000×; and (f) 10,000×.

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Fernandes, B. S., Pinto, J. C., Cabral-Albuquerque, E. C. M., & Fialho, R. L. L. the higher viscosity of the copolymer solution, which can prevent the uniform distribution of polymer material on the granule surface, leading to formation of a less regular and rougher surface during drying. In addition, it can be seen that the geometry of the urea granules was less spherical,

as there was larger concentration of coating in some parts of the urea granules than in others.

Samples prepared through Coating 4 were similar to samples prepared through Coating 3, with uneven covering of the granule surfaces and formation of rough coating

Figure 4. Images of coating 2 (monomeric solution; with pauses) obtained by SEM at magnifications of (a) 50×; (b) 100×; (c) 500×; (d) 1,000×; (e) 5,000×; and (f) 10,000×.

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Coating of urea granules by in situ polymerization in fluidized bed reactors layers. The dark regions of the surface presented lower degree of coating and granules lost the characteristic spherical geometry of urea particles, due to accumulation of coating on some parts of the urea granule. However, in this case, irregularity and roughness of the outer layer

seemed to be smaller to some extent, indicating that the pauses in the flow rate during coating and the more regular drying operation allowed for better uniformity and distribution of the copolymer over the entire surface of the urea granules.

Figure 5. Images of coating 3 (polymeric solution; without pauses) obtained by SEM at magnifications of (a) 50×; (b) 100×; (c) 500×; (d) 1,000×; (e) 5,000×; and (f) 10,000×.

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Fernandes, B. S., Pinto, J. C., Cabral-Albuquerque, E. C. M., & Fialho, R. L. L. Samples prepared through Coating 2 exhibited intermediate characteristics between samples prepared through Coating 1 (using the monomer solution) and Coatings 3 and 4 (using the polymer solution). In this case, samples presented

spherical and regular geometry (although slightly rough).

This indicated that interruption of the flow rate allowed for higher conversion of the monomers into the copolymer, as also observed by FTIR analyses.

Figure 6. Images of coating 4 (polymeric solution; with pauses) obtained by SEM at magnifications of (a) 50×; (b) 100×; (c) 500×; (d) 1,000×; (e) 5,000×; and (f) 10,000×.

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Coating of urea granules by in situ polymerization in fluidized bed reactors Therefore, it was shown that it is possible to produce coated urea granules through in-situ polymerization onto the granule surface using a fluidized bed.

5. Acknowledgements The authors thank CAPES (Coordenação de Aperfeiçoamento de Pessoal de Nível Superior), CNPq (Conselho Nacional de Pesquisa e Desenvolvimento Tecnológico) and FAPERJ (Fundação Carlos Chagas Filho de Apoio à Pesquisa do Estado do Rio de Janeiro) for their financial support and scholarships. The authors also thank Núcleo de Catálise (NUCAT) of COPPE/UFRJ for the SEM analyses.

Figure 7. Release profiles of coatings 1, 2, 3, 4 and uncoated urea granules.

3.3 Urea release profile in distilled water Figure 7 shows results of urea release tests in distilled water. It can be observed that the rate of dissolution of coated products was slower than observed for uncoated urea granules, indicating the existence of a polymer coating layer on urea granules after processing in the fluidized bed reactor.

Coatings 1 and 2, which used monomer solutions, apparently allowed for more efficient coating than Coatings 3 and 4, which were performed with the polymer solution, as also revealed by the slower urea release curves. This result is in accordance with the FTIR analyses, which revealed the formation of bonds between urea and the coating material, and the SEM analyses, which showed more regular coating in the first two cases.

It is important to note that these urea release tests were led out in an environment of extreme concentration, in pure water. In addition, the results obtained in this work seem promising for some agricultural applications, such as the prevention of ammonia evaporation.

4. Conclusions Coated urea granules were produced in a fluidized bed reactor through in-situ copolymerization of acrylic acid and glycerol and addition of acrylic acid / glycerol copolymers in water. Coated and uncoated urea granules were characterized by Fourier transform infrared spectroscopy, scanning electron microscopy and rates of urea release in water. The obtained results indicated that it is possible to coat urea granules with polymer materials using the two proposed solutions. In particular, FTIR analyses indicated the formation of polymer layers and chemical interaction between urea and the polymer material. Besides, coating of the urea granules significantly reduced the number of pores on the surface and allowed for production of more regular and smoother surfaces, particularly when the in‑situ polymerization scheme was applied. The rates of urea release in water showed that the coated granules exhibited slightly slower rates of dissolution due to the presence of the coating layer and reduced porosity of the granule surfaces.

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28. Fernandes, B. S., Carlos Pinto, J., Cabral‐Albuquerque, E., & Fialho, R. L. (2015). Free‐radical polymerization of urea, acrylic acid, and glycerol in aqueous solutions. Polymer Engineering and Science, 55(6), 1219-1229. http://dx.doi. org/10.1002/pen.24081.

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35. Costa, M. M., Cabral-Albuquerque, E. C., Alves, T. L., Pinto, J. C., & Fialho, R. L. (2013). Use of polyhydroxybutyrate and ethyl cellulose for coating of urea granules. Journal of Agricultural and Food Chemistry, 61(42), 9984-9991. http:// dx.doi.org/10.1021/jf401185y. PMid:24059839.

Received: July 27, 2017 Revised: Feb. 20, 2018 Accepted: June 01, 2018

Polímeros, 29(1), e2019004, 2019

ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.02118

Morphological, thermal and bioactivity evaluation of electrospun PCL/β-TCP fibers for tissue regeneration Lilian de Siqueira1* , Fábio Roberto Passador2, Anderson Oliveira Lobo3 and Eliandra de Sousa Trichês1 Laboratório de Biocerâmica – BIOCERAM, Instituto de Ciência e Tecnologia – ICT, Universidade Federal de São Paulo – UNIFESP, São José dos Campos, SP, Brasil 2 Laboratório de Tecnologia de Polímeros e Biopolímeros – TecPBio, Instituto de Ciência e Tecnologia – ICT, Universidade Federal de São Paulo – UNIFESP, São José dos Campos, SP, Brasil 3 Laboratório Interdisciplinar de Materiais Avançados – LIMAV, Departamento de Engenharia de Materiais, Universidade Federal do Piauí – UFPI, Teresina, PI, Brasil 1

*lilian.siqueira@unifesp.br

Abstract Electrospinning is a simple and low-cost way to fabricate fibers. Among the various polymers used in electrospinning, polycaprolactone (PCL) stands out due to its excellent biodegradability and biocompatibility. However, PCL has some limitations such as low bioactivity, hydrophobic surface, and long in vivo degradation. Calcium phosphate ceramics have been recognized as an attractive biomaterial. They are bioactive and osteoinductive, and some are even quite biodegradable. Different contents of particles of beta-tricalcium phosphate (β-TCP) were incorporated in polymer matrix to form fibers of PCL/β-TCP composites by electrospinning for possible application in tissue regeneration. The presence of β-TCP particles promoted some changes in the thermal properties of the fibers. The immersion of PCL/β-TCP 8 wt-% fibers in simulated body fluid (SBF) caused the formation of a denser and homogeneous apatite layer on its surface.

Keywords: electrospinning, fibers, polycaprolactone, scaffolds, tricalcium phosphate.

How to cite: Siqueira, L., Passador, F. R., Lobo, A. O., & Trichês, E. S. (2019). Morphological, thermal and bioactivity evaluation of electrospun PCL/β-TCP fibers for tissue regeneration. Polímeros: Ciência e Tecnologia, 29(1), e2019005. https://doi.org/10.1590/0104-1428.02118.

1. Introduction The use of electrospun fibers in biomedical applications as scaffolds has increased in the last years because these fibers offer a range of attractive features such as high surface area, high porosity, and ease of incorporation of functional components (bioactive nanoparticles, drug, gene, enzyme, etc.)[1]. Electrospinning is regarded as a simple and versatile top‑down approach for fabricating uniform ultra-fine fibers in a continuous process and at long length scales[2] by applying an electrostatic field to a polymer solution driven by high voltage supply between a needle tip and the collector[3].

Among the biodegradable polymers used in the electrospinning, the synthetic aliphatic polyesters, such as poly (lactic acid) (PLA), poly (glycolic acid) (PGA), and polycaprolactone (PCL), stand out[4,5]. However, these polymers have some limitations such as low bioactivity, hydrophobic surface, and long in vivo degradation[6].

The incorporation of inorganic particles into the polymer matrix or coatings of polymer matrix are alternatives to improve these limitations[7].

Calcium phosphate bioceramics are recognized as an attractive biomaterial because their chemical composition is similar to the mineral component of bone. Moreover, they are bioactive and biodegradable, with show osteoinductive

Polímeros, 29(1), e2019005, 2019

properties. Among the ceramics of calcium phosphate, beta-tricalcium phosphate (β-TCP) [Ca3(PO4)]2 stands out for its osteoconductive activity[8]. Fibers of PCL and its composites have been obtained successfully as shown in the literature by the works of Lu et al.[9], Ribeiro et al.[10], and Hassan et al.[3,11]. Recently, Park et al.[12] prepared and characterized composite fibers of lactic acid and PCL (LA/PCL) containing β-TCP nanoparticles (1-2 wt-%) by electrospinning. The incorporation of β-TCP in LA/PCL fibers could change the microstructure and could lead to a better degradation and biocompatibility of the composite mat. In another study, Kim and Kim[13] generated highly porous electrospun 3D polycaprolactone/β-TCP biocomposites (with 5, 10 and 15 wt-% of β-TCP particles) for tissue regeneration using modified wet-electrospinning supplemented with a femtosecond laser. The fabricated scaffolds demonstrated improved mechanical properties and relatively high cellular activities compared to other methods like rapid-prototyped.

In the literature there are reports of tissue engineering applications of electrospun fibers of PCL and PCL/β-TCP. However, most of these works have focused on the biological tests. Therefore, we fabricated and characterized electrospun fibers of PCL and PCL/β-TCP composites with different

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

Siqueira, L., Passador, F. R., Lobo, A. O., & Trichês, E. S. contents of β-TCP that exhibited interesting structural properties. In this work, we investigated how the addition of β-TCP particles influenced the thermal behavior of the polymer matrix; for example, the crystallinity change as related to the degradation time. The relationships between structure (e.g., morphology and fiber diameter) and property (e.g., thermal and bioactive properties) were established and elucidated in detail.

2. Materials and Methods 2.1 Materials The polymer used in this work was PCL supplied

by Sigma-Aldrich Company Ltd. (USA) with 80 kDa of molecular mass. The solvents used were chloroform [CHCl3, 99%] and methanol [CH3OH, 99%], both supplied

by Synth Acessórios e Equipamentos para Laboratório (Brazil). The β-TCP was produced in the Laboratory of Bioceramics (BIOCERAM) of the Federal University of Sao Paulo using the following reagents: calcium carbonate [CaCO3, 99%] and bibasic anhydrous phosphate [CaHPO4, 99%], both supplied by Synth Acessórios e Equipamentos para Laboratório (Brazil).

2.2 Synthesis and characterization of β-TCP particles The β-TCP powder was synthesized by a solid-state reaction, as described elsewhere[14]. Briefly, a 2:1 molar ratio mixture of CaHPO4 and CaCO3 was calcined at 1050 ºC (Oven, Inox Line/3000 3P, EDG Equipamentos e Controles Ltda., Brazil) for 6 h followed by milling in a horizontal ball mill (Ball mill, MA500, Marconi Equipamentos para Laboratório Ltda., Brazil) during 48 h (alumina milling media of 6 mm of diameter and ball/powder weight ratio of 10:1) and then ground in a high energy ball mill (Planetary mill, Pulverisette 5, Fritsch Company Ltd., Germany) for 4 h (rotation 250 rpm, alumina balls with 2 mm of diameter).

The resulting powder was analyzed by laser light diffraction (Laser Particle Size Analyzer, CILAS 1190L, Cilas Company Ltd., France) and presented a mean particle size of 1.13μm (D50 value) and a particle size distribution between 0.04μm (D10 value) and 4.0μm (D90 value), where D = diameter.

2.3 Preparation of PCL/β-TCP composites by electrospinning The PCL was first dissolved in chloroform under ultrasound dispersion for 30 min; after its complete dissolution, methanol was added, forming a 75/25 v/v the solvent solution[10]. The PCL final concentration in this solution was 1.2 g/ml. To produce the fibers of PCL/β-TCP composites, the particles of β-TCP (D50 = 1.13 μm) were first mixed with chloroform under ultrasound dispersion for 30 min. The PCL solution was then added to the β-TCP suspension to form a 75/25 ratio solvent and after complete dissolution, methanol solvent was added. The same ratio of solvents was also used to prepare solutions containing just the β-TCP particles (75/25 v/v). The final suspensions were maintained under agitation for at least 12 h until complete homogenization. The final concentrations of β-TCP in the suspensions were 0.01, 0.05, and 0.08 g/ml.

The samples were named PCL/β-TCP 1 wt-%, PCL/β-TCP 2/6

5 wt-% and PCL/β‑TCP 8 wt-%, according to the β-TCP content. To assemble the electrospinning apparatus we used: a high voltage source (High Voltage Power Supply, Series 230R, Bertan Ind. e Com. de Máquinas, Brazil), a static collector, attached to an insulating rod claw holding a syringe containing the polymer solution, the needle (Inbras® 0.8 mm diameter and 25 mm length), and a glass syringe with a volume of 20 ml with the plunger. The working distance was maintained at 10 cm from the collector, and the applied voltage was 12 kV, with the fibers collected as randomly oriented.

2.4 Fibers characterization The morphology of the fibers of neat PCL and the PCL/β-TCP composites with different contents of β-TCP was evaluated by scanning electron microscopy (SEM, EVO MA10, ZEISS Company Ltd., Germany). The mean diameter of the fibers was measured using the ImageJ 6.0 software.

The data were obtained from 30 fibers and expressed as mean ± standard deviation.

The chemical composition of the samples was investigated using attenuated total reflection (ATR) with an FT-IR Spectrometer (FTIR, Nicolet iS5, Thermo Fisher Scientific Inc., USA) with a scanning range of 400 to 4000 cm-1 and resolution of 4 cm-1.

The thermal behavior of the samples was analyzed using differential scanning calorimetry (DSC, 204 F1 – Phoenix, NETZSCH Group, Germany). A first heating scan was done from 25 to 120 °C, at 10 °C/min to obtain the glass transition and crystalline melting temperatures of the samples.

Degree of crystallinity (Xc) for the first heating was calculated according to Equation 1.

Xc ( % ) =

∆H m x100 ∆H m∞ . φ PCL

(1)

Where: ∆HM is the melting enthalpy of the sample, ∆HM∞ is the melting enthalpy of a 100% crystalline sample and ϕPCL is the mass fraction of PCL in the fiber. For the PCL samples, ∆HM∞ is 136 J/g[15].

In vitro bioactivity of the samples was evaluated by immersing them in 15 ml of the Simulated Body Fluid solution[16] (1.5 SBF, pH 7.4) at 37 ºC. After given times of 7, 14 and 21 days of immersion, the samples were extracted from the SBF and analyzed by SEM (EVO MA10, ZEISS Company Ltd., Germany) and X-Ray diffraction (XRD, X’pert Powder, PANalytical Co., Ltd., Netherlands) operating at 45 kV/40 mA (CuKα, λ = 0.154 nm); the samples were scanned with steps of 0.02º between 2θ = 20 and 50º.

3. Results and Discussion Figure 1 shows SEM micrographs of the fibers. The fibers of PCL were smooth, homogeneous, and without defects (Figure 1a and b). When 1 wt-% of β-TCP particles was added, some clusters appeared within the polymeric fibers indicating that these particles were incorporated into the polymer matrix (Figures 1c and d). The fiber diameters of Polímeros, 29(1), e2019005, 2019

Morphological, thermal and bioactivity evaluation of electrospun PCL/β-TCP fibers for tissue regeneration

Figure 1. SEM micrographs of the fibers: (a, b) PCL; (c, d) PCL/β-TCP 1 wt-%; (e, f) PCL/β-TCP 5 wt-%; and (g, h) PCL/β-TCP 8 wt-%.

PCL/β-TCP decreased with addition of 1 wt-% of β-TCP (640 ± 20 nm) when compared to PCL fibers (774 ± 49 nm).

This physical phenomenon may be associated with the difference in the density between the filler and the polymer Polímeros, 29(1), e2019005, 2019

matrix. In addition, with increasing content of β-TCP, the presence of a larger number of agglomerates within the fibers (Figures 1e and f) was observed. However, with the addition of 8 wt-% of β-TCP, the presence of agglomerates 3/6

Siqueira, L., Passador, F. R., Lobo, A. O., & Trichês, E. S. of larger size within the fibers (Figure 1g) and changes in the roughness of the fibers (Figure 1h) can be seen. PCL fibers with the addition of 5 and 8 wt-% of β-TCP showed an increase in the average diameter with values of 867 ± 40 nm and 726 ± 110 nm, respectively.

related to PO43-, which is an indication of incorporation of the β-TCP in the polymeric fibers, showing that a physical interaction occurred between them. When β-TCP particles were present at 8 wt-%, an increase in the intensity of the β-TCP absorption bands was observed.

FTIR analyses of the functional organic groups in the polymers and in the β-TCP are shown in Figure 2. In the case of the PCL, the main absorbance band at 1723 cm-1 corresponds to carbonyl stretching[17] while the bands at 1294, 1240.2, and 1162.2 cm-1 correspond to the stretching vibrations of the C-O-C groups. In all spectra, except for the PCL and PCL/β-TCP 1 wt-% fibers, the bands at 557 and 615 cm-1 correspond to the absorption bands

Figure 3 shows the DSC thermograms of the fibers and Table 1 shows the thermal parameters calculated from these thermograms. The crystalline melting peak occurred in the first heating at 61.2 °C, 62.2 °C, 62.7 °C, and 61.8 °C for the samples of PCL, PCL/β-TCP 1 wt-%, PCL/β-TCP 5 wt-%, and PCL/β-TCP 8 wt-%, respectively The glass transition temperature values of the different samples remained virtually unchanged.

The degree of crystallinity was modified by the addition of β-TCP particles. Increasing the content of β-TCP decreased the degree of crystallinity. This result can be considered representative and can be associated with particle behavior, which acted as obstacles for the diffusion of the macromolecules toward the growing crystal surface (Table 1). It is a positive result because it favors the degradation process, remembering that the PCL scaffold or implant takes many months or even years to degrade in vitro or in vivo[6,18,19].

Bioactivity, one of the most desirable properties of the materials for bone tissue regeneration, can be inferred by the formation of bone-like apatite on the surface of the materials in contact with the SBF solution in vitro. Figure 4a,b,c,d shows micrographs of the fibers after the bioactivity assay.

The presence of the typical HA globular morphology on the surface of all of the electrospun fibers can clearly be seen.

However, when 8 wt-% of β-TCP was incorpored to the polymeric matrix, a more homogeneous layer of HA can be observed on the surface of the fibers.

Figure 2. FTIR spectra of fibers of PCL and PCL/β-TCP composites with different contents of β-TCP.

The XRD pattern of the fibers before and after the biomineralization process can be observed at right in Figure 4. For all fibers, two distinct diffraction peaks were observable, at 2θ = 21.5° and 2θ = 23.9°, indexed to the (110) and (200) planes respectively of the orthorhombic crystal structure of PCL[20]. The peaks at 31.2° and 34.5° correspond to the diffraction peaks of the β-TCP (JCPDS 009-0169) while the peaks at 26, 32, and 45.6° were related to HA (JCPDS 74-0565).

The obtained results demonstrated the fibers were very promising for a possible application in tissue engineering.

Among all the compositions studied it is suggest that the fibers of PCL/β-TCP with 8 wt-% of β-TCP are the most suitable due to improvement in its morphological and bioactive properties.

Figure 3. DSC thermograms of the fibers.

Table 1. Thermal parameters of the fibers of PCL and PCL/β-TCP composites with different contents of β-TCP. Material PCL PCL/β-TCP 1 wt-% PCL/β-TCP 5 wt-% PCL/β-TCP 8 wt-%

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Tg (°C) -63.2 -63.4 -61.4 -61.4

Tm (°C) 61.2 62.2 62.7 61.8

∆Hm (J/g) 71.3 67.7 51.0 40.1

Xc (%) 52.4 50.3 39.5 32.0

Polímeros, 29(1), e2019005, 2019

Morphological, thermal and bioactivity evaluation of electrospun PCL/β-TCP fibers for tissue regeneration

Figure 4. Micrographs of the fibers: (a) PCL; (b) PCL/β-TCP 1 wt-%; (c) PCL/β-TCP 5 wt-%; and (d) PCL/β-TCP 8 wt-% after biomineralization for 21 days. Right: XRD patterns.

Polímeros, 29(1), e2019005, 2019

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Siqueira, L., Passador, F. R., Lobo, A. O., & Trichês, E. S.

4. Conclusions Fibers of PCL and PCL/β-TCP with 1, 5 and 8 wt-% of β-TCP were successfully produced by electrospinning.

The best fibers obtained for future applications in tissue engineering were the PCL/β-TCP 8 wt-%, due to its lower degree of crystallinity and better bioactive properties.

5. Acknowledgements The authors would like to thank the São Paulo Research Foundation - FAPESP (grant numbers 2011/17877-7 for A.O.L. and 2010/00863-0), National Counsel of Technological and Scientific Development - CNPq (grant number 303752/2017-3 for A.O.L) for the financial support. We also thank João Paulo Barros Machado, PhD from the Laboratory of Sensors and Materials (LAS) of the National Institute of Space Research (INPE) for permitting the XRD analyse and Ana Paula Lemes, PhD from Federal University of São Paulo (UNIFESP) for thermal analyses.

6. References 1. Zhou, H., Lawrence, J. G., & Bhaduri, S. B. (2012). Fabrication aspects of PLA-CaP/PLGA-CaP composites for orthopedic applications: a review. Acta Biomaterialia, 8(6), 1999-2016. http://dx.doi.org/10.1016/j.actbio.2012.01.031. PMid:22342596.

2. Sun, B., Long, Y. Z., Zhang, H. D., Li, M. M., Duvail, J. L., Jiang, X. Y., & Yin, H. L. (2014). Advances in three- dimensional nanofibrous macrostructures via electrospinning. Progress in Polymer Science, 39(5), 862-890. http://dx.doi.org/10.1016/j. progpolymsci.2013.06.002.

3. Hassan, M. I., Sun, T., & Sultana, N. J. (2014). Fabrication of Nanohydroxyapatite/Poly(caprolactone) composite microfibers using electrospinning technique for tissue engineering applications. Journal of Nanomaterials, 2014(65), 1-7. http:// dx.doi.org/10.1155/2014/209049.

4. Garkhal, K., Verma, S., Jonnalagadda, S., & Kumar, N. (2007). Fast degradable poly (L -lactide- co - e -caprolactone) microspheres for tissue engineering : synthesis, characterization, and degradation behavior. Journal of Polymer Science. Part A, Polymer Chemistry, 45(13), 2755-2764. http://dx.doi. org/10.1002/pola.22031.

5. Zhang, Y., Lim, C. T., Ramakrishna, S., & Huang, Z. M. (2005).

Recent development of polymer nanofibers for biomedical and biotechnological applications. Journal of Materials Science. Materials in Medicine, 16(10), 933-946. http://dx.doi. org/10.1007/s10856-005-4428-x. PMid:16167102.

6. Ma, P. X. (2004). Scaffolds for tissue fabrication. Materials Today, 7(5), 30-40. http://dx.doi.org/10.1016/S1369-7021(04)00233-0.

7. Holzapfel, B. M., Reichert, J. C., Schantz, J. T., Gbureck, U., Rackwitz, L., Nöth, U., Jakob, F., Rudert, M., Groll, J., & Hutmacher, D. W. (2012). How smart do biomaterials need to be? A translational science and clinical point of view.

Advanced Drug, 65(4), 581-603. http://dx.doi.org/10.1016/j. addr.2012.07.009. PMid:22820527.

8. Kawachi, E. Y., Bertran, C. A., Reis, R. R., & Alves, O. L. (2000). Biocerâmicas: tendências e perspectivas de uma área interdisciplinar. Quimica Nova, 23(4), 518-522. http://dx.doi. org/10.1590/S0100-40422000000400015.

9. Lu, L., Zhang, Q., Wootton, D., Chiou, R., Li, D., Lu, B., Lelkes, P., & Zhou, J. (2012). Biocompatibility and biodegradation

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studies of PCL/β-TCP bone tissue scaffold fabricated by structural porogen method. Journal of Materials Science.

Materials in Medicine, 23(9), 2217-2226. http://dx.doi. org/10.1007/s10856-012-4695-2. PMid:22669285.

10. Ribeiro, W. A. R., No., Pereira, I. H. L., Ayres, E., De Paula, A. C. C., Averous, L., Góes, A. M., Oréfice, R. L., & Bretas, R. E. S. (2012). Influence of the microstructure and mechanical strength of nanofibers of biodegradable polymers with hydroxyapatite in stem cells growth. Electrospinning, characterization and cell viability. Polymer Degradation & Stability, 97(10), 2037-2051. http://dx.doi.org/10.1016/j.polymdegradstab.2012.03.048.

11. Hassan, M. I., Sultana, N., & Hamdan, S. (2014). Bioactivity assessment of poly(ɛ-caprolactone)/hydroxyapatite electrospun fibers for bone tissue engineering application. Journal of Nanomaterials, 2014, 1-6. http://dx.doi.org/10.1155/2014/573238.

12. Park, C. H., Kim, E. K., Tijing, L. D., Amarjargal, A., Pant, H. R., Kim, C. S., & Shon, H. K. (2014). Preparation and characterization of LA/PCL composite fibers containing beta tricalcium phosphate (β-TCP) particles. Ceramics International, 40(3), 5049-5054. http://dx.doi.org/10.1016/j. ceramint.2013.10.016.

13. Kim, M. S., & Kim, G. H. (2014). Highly porous electrospun 3D polycaprolactone/β-TCP biocomposites for tissue regeneration.

Materials Letters, 120, 246-250. http://dx.doi.org/10.1016/j. matlet.2014.01.083.

14. Siqueira, L., Passador, F. R., Costa, M. M., Lobo, A. O., & Sousa, E. (2015). Influence of the addition of β-TCP on themorphology, thermal properties and cell viability of poly (lactic acid) fibers obtained by electrospinning. Materials Science and Engineering C, 52, 135-143. http://dx.doi.org/10.1016/j. msec.2015.03.055. PMid:25953550.

15. Pereira, R. B., & Morales, A. R. (2014). Estudo do comportamento térmico e mecânico do PLA modificado com aditivo nucleante e modificador de impacto. Polímeros: Ciência e Tecnologia, 24(2), 198-202. http://dx.doi.org/10.4322/polimeros.2014.042.

16. Kokubo, T., Kushitani, H., Sakka, S., Kitsugi, T., & Yamamuro, T. (1990). Solutions able to reproduce in vivo surface-structure changes in bioactive glass-ceramic A-W. Journal of Biomedical Materials Research, 24(6), 721-734. http://dx.doi.org/10.1002/ jbm.820240607. PMid:2361964.

17. Sasmazel, H. T. (2011). Novel hybrid scaffolds for the cultivation of osteoblast cells. International Journal of Biological Macromolecules, 49(4), 838-846. http://dx.doi.org/10.1016/j. ijbiomac.2011.07.022. PMid:21839769.

18. Vert, M., Li, S. M., Spenlehauer, G., & Guerin, P. (1992).

Bioresorbability and biocompatibility of aliphatic polyesters.

Journal of Materials Science. Materials in Medicine, 3(6), 432-446. http://dx.doi.org/10.1007/BF00701240.

19. Pereira, C. S., Gomes, M. E., Reis, R. L., & Cunha, A. (1999).

Hard cellular materials in the human body: properties and production of foamed polymers for bone replacement. In J. F. Sadoc , & N. Rivier (Eds.), Foams and emulsion (pp.

354, 193-206). USA: Springer Netherlands. http://dx.doi. org/10.1007/978-94-015-9157-7_12.

20. Baji, A., Wong, S. C., Liu, T., Li, T., & Srivatsan, T. S. (2007).

Morphological and x-ray diffraction studies of crystalline hydroxyapatite-reinforced polycaprolactone. Journal of Biomedical Materials Research. Part B, Applied Biomaterials, 81B(2), 343-350. http://dx.doi.org/10.1002/jbm.b.30671.

PMid:17022054.

Received: Apr. 10, 2018 Accepted: June 02, 2018

Polímeros, 29(1), e2019005, 2019

ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.11217

Non-isothermal melt crystallization kinetics of poly(3‑hydroxybutyrate), poly(butylene adipate-co-terephthalate) and its mixture Anna Raffaela Matos Costa1* , Edson Noryuki Ito1, Laura Hecker Cavalho2 and Eduardo Luís Canedo2 Programa de Pós-graduação em Ciência e Engenharia de Materiais, Departamento de Engenharia de Materiais, Centro de Tecnologia – CT, Universidade Federal de Rio Grande do Norte – UFRN, Natal, RN, Brasil 2 Unidade Acadêmica de Engenharia de Materiais – UAEMa, Universidade Federal de Campina Grande – UFCG, Campina Grande, PB, Brasil

1

*raffaela_matos@yahoo.com.br

Abstract Nonisothermal crystallization and melting of the biodegradable thermoplastics poly(3-hydroxybutyrate) (PHB), poly(butylene adipate-co-terephthalate) (PBAT), and a 1:1 PHB/PBAT blend were investigated by differential scanning calorimetry (DSC) over an extensive range of heating/cooling rates (2 to 64°C/min). The different phase transition behavior of the neat components was reflected in the mixture and suggest an immiscible blend. Pseudo-Avrami, Ozawa and Mo classical macrokinetic models were used to describe the evolution of the melt crystallization process. Results suggest that none of these models could be used to predict the experimental results of crystallization kinetics of the blend with sufficient precision for polymer processing applications. However, some methods may be of used for the neat resins over restricted ranges of cooling rate, temperature or conversion (e.g., Ozawa for PHB at low cooling rate, Mo for PBAT).

Keywords: PHB, PBAT, blends, crystallization kinetics.

How to cite: Costa, A. R. M., Ito, E. N., Carvalho, L. H., & Canedo, E. L. (2019). Non-isothermal melt crystallization kinetics of poly(3-hydroxybutyrate), poly(butylene adipate-co-terephthalate) and its mixture. Polímeros: Ciência Tecnologia, 29(1), e2019006. https://doi.org/10.1590/0104-1428.11217

1. Introduction One the most challenging problems of materials technology today is the negative environmental impact of plastic materials. The search of alternative materials and technologies to solve the problem of waste disposal of plastics of petrochemical origin is driving the development of biodegradable polymers[1].

Poly(3-hydroxybutyrate) (PHB) is a biodegradable, biocompatible thermoplastic obtained from renewable natural sources that can be processed with conventional techniques and equipment. However, many applications require an improvement of its properties[2,3]. Moreover, widespread commercial use of PHB is limited by its thermal instability during processing, excessive crystallinity and high cost.

Poly(butylene adipate-co-terephthalate) (PBAT) is a synthetic random aliphatic-aromatic copolyester with good physical properties which is completely biodegradable in municipal landfills. Moreover, PBAT is a flexible, low‑crystallinity thermoplastic, with thermal and mechanical properties similar to some polyethylenes, appropriate for the production of films for the packaging industry. Compared with low density polyethylene (LDPE) films, PBAT films are less permeable to oxygen (50%) and much more permeable to water vapor (80×), which recommend it for food packaging[3,4]. PBAT is relatively stable under processing[5,6]. Blending polymers

Polímeros, 29(1), e2019006, 2019

to improve properties and open new fields of application is a common procedure. Thus the development of a fully biodegradable blend PHB/PBAT could be an interesting alternative.

Most industrial processes are conducted under nonisothermal conditions and successful process development involving semicrystalline polymers requires knowledge of the nonisothermal crystallization and meting processes and its kinetics. Although several studies on thermal properties and crystalline structures of PHB and PBAT are available in the technical literature[7-11], very few – if any – considers the kinetics of crystallization PHB/PBAT blends.

Kinetic models of nonisothermal crystallization could be used to predict the evolution of crystallinity as function of time and temperature. The predictive behavior of an analytical model is useful in that the model may be inserted into processing and manufacturing protocols to develop and control processes in which crystallization in one among several concurrent phenomena, e.g., molding, blowing, spinning, etc.[12].

It must be emphasized that the nonisothermal kinetic models available in the literature are all empirical correlation procedures. Model parameters have no physical meaning

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

Costa, A. R. M., Ito, E. N., Cavalho, L. H., & Canedo, E. L. and are not amenable to microstructural interpretation.

Nevertheless, they are valuable tools in the hands of process engineers, insofar as they can predict the kinetics of crystallization of real systems under processing conditions The present contribution is concerned with a thorough investigation of the nonisothermal crystallization and melting behavior of PHB, PBAT, and a blend PHB/PBAT (1:1) by differential scanning calorimetry (DSC), under a wide range of cooling/heating rates (2 to 64°C/min). Classical kinetic models, associated to the names of Avrami, Ozawa, and Mo were used to correlate the experimental data of melt crystallization. The discrepancy between model predictions and experimental data, and range of applicability of the models were discussed in detail.

Conversions and rates may be expressed as functions of temperature T during each non-isothermal event conducted at constant cooling/heating rate Ď• = |dT/dt|: T= T1 Âą φ(t − t1 )

where t1 and T1 are the time and temperature at the onset of the event. Maximum and mean temperatures and rates were estimated from such plots.

The specific heat of phase change ΔH was computed from the total energy transferred E0 (from the sample to the neighborhood during the exothermic crystallization process, vice versa during the endothermic melting process). Mass crystallinity changes were estimated: X ∆=

2. Materials and Methods 2.1. Materials Poly(3-hydroxybutyrate (PHB) Biocycle 1000 was supplied

by PHB Industrial (Serrana, SP, Brasil), and is actually a copolymer, containing about 4% units of 3-hydroxyvalerate.

Poly (butylene adipate-co-terephthalate) (PBAT) Ecoflex F Blend C1200, purchased from BASF (São Paulo SP, Brasil), is basically a linear random 1:1 copolyester; however, a BASF sponsored publication[5] mention small amounts of a third polyfunctional comonomer of unknown chemical structure, which may result is some degree of ramification.

2.2. Methods A blend containing 50% of PHB and PBAT was compounded in a corotating twin-screw extruder. Samples of neat components were also processed in this way.

The extrudates were analyzed by differential scanning calorimetry (DSC) in a Mettler-Toledo DSC-1 instrument.

Samples of 5 to 7 mg were heated from 25 °C to 190 °C at 24 °C/min and kept at that temperature for 1 min; then, there were cooled to 25 °C at seven different rates, from 2 °C/min to 64 °C/min; after a 3 min isothermal stage at 25 °C, the samples were reheated to 190 °C at the same rates. The process was conducted under a flow of nitrogen gas of 50 °C/min.

Phase transitions (melt crystallization during the cooling stage, cold crystallization and melting during the second heating stage) were identified and analyzed. Conversion x (crystallized or molten fraction of the total transformed mass) was estimated as a function of time integrating the difference between the DSC output (J) and a suitable virtual baseline (J0): = x(t )

1 t âˆŤ J (t ′) − J 0 (t ′) dt ′ E0 t1

(1)

where = E0 âˆŤtt2 J (t ) − J 0 (t ) dt 1

(2)

c(= t)

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dx = dt

J (t ) − J 0 (t ) E0

E0 ∆H =

∆H mo mS ∆H mo

∆X (∆H m0 ) A + ∆X B (∆H m0 ) B ∆H m0 =A ∆X A + ∆X B

(6)

where (ΔH°M )A and (ΔH°M)B are the latent heats of melting of 100% crystalline components A and B , and ΔXA e ΔXB are the corresponding crystallinity developed during the separate melting processes of the same sample. Equation (6) takes advantage of different crystallinity exhibited by the two components, and is considered a better alternative than the simple average using mass fractions. However, it assumes independent phase transition processes, with no interference on one component on the crystallization of the other, which may be valid only as a first approximation.

The kinetics of the nonisothermal melt crystallization process was correlated by three classical empirical models, Pseudo-Avrami, Ozawa, and Mo. With the first model, which is called Pseudo-Avrami to distinguish from the well-known Avrami model for isothermal crystallization, the relative crystallinity, estimated at constant cooling rate τ is expressed as a function of the time since the onset of the event τ = t – t1 as:

(

)

x = 1 − exp − K Ď„n

(7)

The parameters K = K(đ?œ™) and n = n(đ?œ™) were determined by linear regression of 1   y = ln  ln  = ln K + n ln Ď„  1− x 

(8)

For the Ozawa model the relative crystallinity, interpolated from the original data, is correlated in terms of the cooling rate: x = 1 − exp âˆ’ÎşĎ†âˆ’ m

(3)

(5)

where mS is the sample mass and ΔH°M is the specific latent heat of melting of 100% crystalline polymer. For the neat resins values of 146 J/g and 114 J/g for PHB and PBAT, respectively, taken form the literature (Barham et al, 1984; Gan et al, 2004). For the co-crystallization of the blend, ΔH°M was estimated as a weighted average[13]:

(

The rate of conversion was computed as:

(4)

)

(9)

The parameters đ?œ… = (T) e m= m(T) were determined by linear regression: PolĂ­meros, 29(1), e2019006, 2019

Non-isothermal melt crystallization kinetics of poly(3-hydroxybutyrate), poly(butylene adipate-co-terephthalate) and its mixture 1   y = ln  ln = ln Îş − m ln φ  1− x 

(10)

using the pairs (x, đ?œ™) previously computed for each value of the temperature T. Finally, for the model attributed to Mo[11,14] the rate of cooling is expressed as a function of time: φ = F Ď„âˆ’Îą

(11)

based on interpolated data at constant relative crystallinity, estimated from the original results.

The parameters F = F(x) e Îą = Îą (x) are obtained by linear regression: ln= φ ln F − Îą ln Ď„

(12)

using the pairs (x, đ?œ™) previously computed for each value of the relative crystallinity x.

Parameters of the three models were correlated in terms of the corresponding independent variable and the discrepancy between predicted and experimental data computed, as a way to establish the usefulness of the models for polymer processing applications.

was observed, PHB, PBAT and the blend crystalize at substantially the same temperature. Crystallinity, on the other hand, depends on the nature of the polymer. While the low PBAT crystallinity (around 10 to 15%) is realized completely during the cooling stage (no cold crystallization), PHB crystalizes only partially from melt at moderate or high cooling rates; PHB crystallization is completed during the reheating stage. As a consequence, PHB crystallinity from the melt decreases sharply with the cooling rate.

The PHB/PBAT blend follows an intermediate way, more influenced by high crystalline component (PHB) behavior.

Both the rate of crystallization and the rate of melting increase with the rate of temperature change. This expected behavior often masks insights on the kinetics of phase transformation. For this reason, the specific rate of phase change, C = c/ϕ,, is introduced. Figure 4 shows its maximum value as a function of the heating/cooling rate for melt

3. Results and Discussions Figure 1 shows a typical example of heat flow versus time plot (raw DSC results) during the cooling and reheating stages.

Single melt crystallization peaks (C1) were observed for all three samples during the cooling stage , shaper for PBAT, shallower for PHB. Both PHB and the PHAB/PBAT blend further crystallize (C2) at the beginning of the reheating stage; no cold crystallization of neat PBAT was observed.

An endothermic asymmetric, complex melting peak was observed for neat PHB (at about 30 min) and a shallow, simple melting peak for neat PBAT (at about 27 min).

The blend melts in two separate sub-events, one for the PBAT component and one the PHB. Figure 2 shows the melting behavior (F2) in detail, this time as a function of temperature.

PBAT melts as a shallow peak at about 123 °C, both in the neat resin and in the blend; PHB melts as a much larger peak at about 169 °C (with a shoulder around 156 °C), both in the neat resin and in the blend. This observation suggests that both components of crystalize separately, despite the single crystallization peaks observed for the blend.

Several transition parameters were computed from the DSC output integrated according to the procedure described in the previous section. Numerical results for melt and cold crystallization (C1 and C2) and the second melting (F2), for all material and cooling/heating rates tested, are included as Appendix A, Tables A1-A8. Figure 3 show two melt crystallization parameters plotted as functions of the cooling rate: peak temperature and crystallinity.

Melt crystallization temperature decreases as the rate of cooling increases, a behavior frequently observed in polymers: less time to crystalize (higher rate of cooling) requires larger super‑cooling (lower crystallization temperature) to do the job. No significant dependence on the composition Polímeros, 29(1), e2019006, 2019

Figure 1. Heat flow versus time for samples of PHB. PBAT and the PHB/PBAT blend cooled and reheated at 16 °C/min (exothermic peaks up).

Figure 2. Heat flow versus temperature for the melting events (F2) during reheating at 16°C/min. for PBAT. PHB and the PHB/PBAT bled a 16 °C/min.

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Costa, A. R. M., Ito, E. N., Cavalho, L. H., & Canedo, E. L.

Figure 3. Peak temperature (a) and crystallinity (b) as functions of cooling rate for the melt crystallization of PHB. PBAT and of the PHB/PBAT blend.

Figure 4. Specific melt crystallization (a) and melting (b) maximum rates for PHB. PBAT and for the PHB/PBAT blend as functions of the cooling/heating rate.

crystallization during the cooling stage (C1) and for melting during the reheating stage. Specific rates are less dependent of the material and cooling/heating rate and, for ϕ, > 30°C/min, become virtually independent of both.

There are few studies of phase transition in PHBV/PBAT blends in the technical literature. Vidharte el al.[15] observed the separate melting of the two components, with a slight rise of the melting points. During cooling, a single melt crystallization peak was observed in the blends, suggesting co-crystallization of PBHV and PBAT. A small increase of the melt crystallization temperature may be attributed to the interference of PBAT on PHB crystal growth. Similar results were reported by Bittmann et al.[16] and are fully consistent with the detailed results discussed in the present work.

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3.1. Kinetics modeling: Pseudo-Avrami The relative crystallinity x as a function of time τ measured from the onset of the melt crystallization event (C1) at different cooling rates ϕ, was correlated using Equation (7) for PHB, PBAT and the PHB/PBAT blend. Figure 5 shows a typical plot of Pseudo Avrami, Equation (8). Parameters lnK and n computed from a linear regression of the data in the interval 2-98% relative crystallinity. Numerical results are collected in Tables A9-A11 of the Appendix A. Fitting is excellent, with uncertainty of the parameters below 1%.

This is reflected in a fairly good prediction of the experimental results by the Pseudo-Avrami model, as shown in Figure 6 for neat PBAT. Polímeros, 29(1), e2019006, 2019

Non-isothermal melt crystallization kinetics of poly(3-hydroxybutyrate), poly(butylene adipate-co-terephthalate) and its mixture

Figure 5. Pseudo-Avrami plot. Equation (8). for the melt crystallization of PBAT at different cooling rates (indicated).

Figure 6. Experimental relative crystallinity (circles) and Pseudo-Avrami model predictions (lines) for the melt crystallization of PBAT.

A better, quantitative picture of the fitting is obtained by plotting the discrepancy between predicted and experimental x pre − xexp, as shown in Figure 7 relative crystallinity δ= x for neat PBAT. Pseudo-Avrami overpredicts relative crystallinity during the first half of the crystallization processes for as much as 4% and underpredicts relative crystallinity by up to 6% in last stage of the process. Similar results were obtained with PHB and the PHB/PBAT blend.

These are very good results and Pseudo-Avrami is an excellent choice to correlate experimental data. The problems arise when we try to predict the relative crystallinity for melt crystallization processes conducted at arbitrary rates of cooling. Our purpose, as stated in the Introduction, is to obtain reliable correlations for polymer processing applications, not so much for structural characterization. In this case, an analytical correlation of the model parameters with the cooling rate is needed. Figure 8 shows Pseudo-Avrami parameters as functions of the cooling rate. It is clear that the Polímeros, 29(1), e2019006, 2019

dispersion of the results precludes the development of such correlation with a reasonably uncertainty. Pseudo-Avrami cannot be recommended polymer processing applications.

3.2. Kinetics modeling: Ozawa Relative crystallinity x as a function of temperature T for the melt crystallization event (C1) in PHB, PBAT and the PHB/PBAT blend were interpolated and correlated as a function of cooling rate ϕ using Equation (9). Figure 9 shows a typical Ozawa plot, Equation (10).

Ozawa parameters lnκ and m were computed from a linear regression. Results are collected in Tables A12-A14 of the Appendix A, and shown graphically in Figure 10, along with their uncertainties.

Figure 10 suggest that Ozawa model does not represent experimental data very well and cannot be of use to correlate them, even less to predict them satisfactorily for any application: uncertainties are too high. That is the case in general. However, for neat PHB in the restricted 5/16

Costa, A. R. M., Ito, E. N., Cavalho, L. H., & Canedo, E. L.

Figure 7. Discrepancy between the Pseudo-Avrami model predictions and the experimental results for the melt crystallization of PBAT as function of relative crystallinity.

Figure 8. Pseudo-Avrami model parameters ln K (a) and n (b) for the melt crystallization of PHB. PBAT. and for the PHB/PBAT blend as functions of cooling rate.

Figure 9. Ozawa plot. Equation (10). for the melt crystallization of PHB at different temperatures (indicated).

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PolĂ­meros, 29(1), e2019006, 2019

Non-isothermal melt crystallization kinetics of poly(3-hydroxybutyrate), poly(butylene adipate-co-terephthalate) and its mixture

Figure 10. Ozawa model parameters ln ĸ (a) and m (b) for the melt crystallization of PHB. PBAT. and for the PHB/PBAT blend as functions of temperature.

temperature interval from 75 °C to 90 °C the uncertainty of the parameters is less than 5% and they can be correlated using a simple quadratic expression: = ln κ 0,1140T − 0,0001T 2 = m 0,3030T − 0,0032T 2

(13)

as shown in Figure 11.

Fitting is excellent. It turns out that PHB crystallizes from the melt in the reduced temperature interval of 40‑100 °C when cooled at rates equal or less that 6 °C/min. Figure 12 shows how it performs predicting the values of relative crystallinity versus temperature at low cooling rate.

Ozawa model seems to be perfectly adequate to correlate and predict experimental melt crystallization results for this particular material and cooling rate range. Unpromising models cannot be summarily dismissed without a thorough examination.

Figure 11. Ozawa model parameters for the melt crystallization of PHB in the interval of 75-90 °C.

3.3. Kinetics modeling: Mo The cooling rate ϕ required to reach different values of relative crystallinity at a given time τ was interpolated for the melt crystallization event (C1) in PHB, PBAT and the PHB/PBAT blend. Cooling rates were correlated with time at constant relative crystallinity using Equation (11).

Figure 13 shows a typical Mo plot, Equation (12).

Mo parameters lnF and α were computed from a linear regression. Results are collected in Tables A15-A17 of the Appendix A, and shown graphically in Figure 14, along with their uncertainties.

Figure 14 shows Mo parameters could be easily and precisely correlated with relative crystallinity by simple polynomial expressions. However, it shows also a significant uncertainty associated to the parameters for PHB and the PHB/PBAT blend; PBAT appears to be the exception.

But smoothness and low uncertainty in the parameters not always translates in good predictive behavior. For example, it can be proved (see Appendix B) that Mo model does not Polímeros, 29(1), e2019006, 2019

Figure 12. Experimental relative crystallinity (circles) and Ozawa model predictions (lines) for the melt crystallization of PHB at low cooling rates.

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Costa, A. R. M., Ito, E. N., Cavalho, L. H., & Canedo, E. L.

Figure 13. Mo plot. Equation (12). for the melt crystallization of PBAT at different relative crystallinities (indicated).

Figure 14. Mo parameters ln F (a) and Îą (b) for the melt crystallization of PHB. PBAT. and the PHB/PBAT blend as functions of relative crystallinity.

Figure 15. Discrepancy between the Mo model predictions and the experimental results for the melt crystallization of PBAT as function of relative crystallinity.

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PolĂ­meros, 29(1), e2019006, 2019

Non-isothermal melt crystallization kinetics of poly(3-hydroxybutyrate), poly(butylene adipate-co-terephthalate) and its mixture lead to a true sigmoid shape of the relative crystallinity versus time or temperature curve, and fails at the beginning of the crystallization process: c = dx/dτ → 0 as τ → 0 only if lnF → 0 as x → 0, which is not the case, according to Figure 14.

Discrepancy between experimental results and predictions of Mo model is better expressed in terms of time, which may be estimated inverting Equation (11): 1/ α( x)  F ( x) 

τ pre =    φ 

(14)

Relative discrepancy for a given cooling rate could be defined as a function of relative crystallinity as: τ pre − τexp δτ =

τ½

(15)

where τEXP is the time to reach a relative crystallinity x during a test at a given cooling rate, τPRE is computed from Equation (14), and τ½ is a characteristic time, taken as the half-crystallization time at the cooling rate in question.

Figure 15 and 16 shows the discrepancy versus relative crystallinity for the melt crystallization o neat PBAT and the PHB/PBAT blend, respectively.

Figure 15 shows that, even in the best case, neat PBAT, Mo model can only predict the kinetics of the melt crystallization processes with an uncertainty of ±25% for relative crystallinity between 15% and 85%. If an application can tolerate that level of uncertainty and limited interval of validity, Mo model is excellent choice. Otherwise, it is not appropriate.

Figure 16 shows that the goodness-of-fit of the Mo model to experimental results of the PHB/PBAT blend is unacceptable for any application. This result couldn’t be expected by considering the fitting of Mo parameters, as Figure 14 shows a very smooth behavior. The importance of the discrepancy study, in addition of the parameters fitting – illustrated by a comparison of Figures 14-16 is perhaps the most important contribution of the present work.

4. Conclusions A thorough – albeit limited in scope – study of the PHB/PBAT system by a conventional DSC technique reveals interesting facts of its melting and crystallization behaviour, and of their nonisothermal melt crystallization kinetics.

PHB and PBAT exhibit very different degrees of crystallinity, melting temperatures and peak structures, but these characteristics are independent of the cooling rate and are the same for the neat resins and the blend. Neat PHB crystallizes partially from the melt, even at moderate cooling rates, while PBAT crystallizes completely from the melt at all cooling rates; the blend follows a combined trend that suggests separate crystallizations of the two components. These observations suggest an immiscible, incompatible blend.

None of the three classical kinetic models tested was able to predict the experimental data for the PHB/PBAT blend with acceptable uncertainty. Some models are usable – as predictive tools – for the net components (Ozawa for PHB, Mo for PBAT) over a reduced range of validity. The study confirms the often overlooked fact that good fitting of the model parameters with the available data not always translates in a good predictive behavior, as judged by the discrepancy between model predicted and experimental data.

5. Acknowledgements The authors wish to thank the Conselho Nacional de Pesquisa - CNPq e Coordenação de Aperfeiçoamento de Pessoal Superior - CAPES, Grant # 473622/2013-0, for financial support.

6. References 1. Xiaohui, W., Jun, S., Ying, C., Zhifeng, F., & Yan, S. (2012).

Nonisothermal crystallization kinetics of poly(butylene adipateco-terephthalate) copolyester. China Petroleum Processing and Petrochemical Technology, 14(1), 74-79.

Figure 16. Discrepancy between the Mo model predictions and the experimental results for the melt crystallization of PHB/PBAT blend as function of relative crystallinity.

Polímeros, 29(1), e2019006, 2019

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Costa, A. R. M., Ito, E. N., Cavalho, L. H., & Canedo, E. L. 2. Fukushima, K., Wu, M. H., Bocchini, S., Rasyida, A., & Yang, M. C. (2012). PBAT based nanocomposites for medical and industrial applications. Materials Science and Engineering C, 32(6), 1331-1351. http://dx.doi.org/10.1016/j.msec.2012.04.005.

PMid:24364930.

3. Al-Itry, R., Lamnawar, K., & Maazouz, A. (2012). Improvement of thermal stability, rheological and mechanical properties of PLA, PBAT and their blends by reactive extrusion with functionalized epoxy. Polymer Degradation & Stability, 97(10), 1898-1914. http://dx.doi.org/10.1016/j.polymdegradstab.2012.06.028.

4. Barham, P. J., Keller, A., Otun, E. L., & Holmes, P. A. (1984).

Crystallization and morphology of a bacterial thermoplastic: poly-3-hydroxybutyrate. Journal of Materials Science, 19(9), 2781-2789. http://dx.doi.org/10.1007/BF01026954.

5. Yamamoto, M., Witt, U., Skupin, G., Beimborn, D., & Müller, R. J. (2002). Biodegradable aliphatic-aromatic polyesters: Ecoflex. In Y. Doi, & A. Steinbüchel (Eds.), Biopolymers, polyesters III – applications and commercial products (pp.

299-305). Weinheim: Wiley-VCH Verlag GmbH. 6. Costa, A. R. M., Almeida, T. G., Silva, S. M. L., Carvalho, L. H., & Canedo, E. L. (2015). Chain extension in poly(butyleneadipate-terephthalate). Inline testing in a laboratory internal mixer. Polymer Testing, 42, 115-121. http://dx.doi.org/10.1016/j. polymertesting.2015.01.007.

7. Witt, U., Einig, T., Yamamoto, M., Kleeberg, I., Deckwer, W. D., & Müller, R. J. (2001). Biodegradation of aliphaticaromatic copolyesteres: evaluation of the final biodegradability and ecotoxicological impact of degradation intermediates.

Chemosphere, 44(2), 289-299. http://dx.doi.org/10.1016/ S0045-6535(00)00162-4. PMid:11444312.

8. Avérous, L., & Fringant, C. (2001). Association between plasticized starch and polyesters: processing and performances of injected biodegradable systems. Polymer Engineering and Science, 41(5), 727-734. http://dx.doi.org/10.1002/pen.10768.

9. Gan, Z. H., Kuwabara, K., Yamamoto, M., Abe, H., & Doi, Y. (2004). Solid-state structures and thermal properties of aliphaticaromatic poly(butylene adipate-co-butylene terephthalate) copolyesters. Polymer Degradation & Stability, 83(2), 289-300. http://dx.doi.org/10.1016/S0141-3910(03)00274-X.

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10. Chen, C., Fei, B., Peng, S., Zhuang, Y., Dong, L., & Feng, Z. (2002). Nonisothermal crystallization and melting behavior of poly(3-hydroxybutyrate) and maleated poly(3-hydroxybutyrate).

European Polymer Journal, 38(80), 1663-1670. http://dx.doi. org/10.1016/S0014-3057(02)00046-0.

11. An, Y., Dong, L., Mo, Z., Liu, T., & Feng, Z. (1998). Nonisothermal crystallization kinetics of poly(β-hydoxybutyrate). Journal of Polymer Science. Part B, Polymer Physics, 36(8), 1305-1312. http:// dx.doi.org/10.1002/(SICI)1099-0488(199806)36:8<1305::AIDPOLB5>3.0.CO;2-Q. 12. Schultz, J. M. (2001). Polymer crystallization – the development of crystalline order in thermoplastic polymers. New York: Oxford University Press.

13. Avella, M., Martuscelli, E., Orsello, G., Raimo, M., & Pascucci, B. (1997). Poly(3- hydroxybutyrate)/poly(methyleneoxide) blends: thermal, crystallization and mechanical behaviour.

Polymer, 38(25), 6135-6143. http://dx.doi.org/10.1016/S00323861(97)00166-3.

14. Liu, T., Mo, Z., & Zhang, H. (1998). Nonisothermal crystallization behavior of a novel poly(aryl ether ketone): PEDEKmK. Journal of Applied Polymer Science, 67(5), 815-821. http://dx.doi. org/10.1002/(SICI)1097-4628(19980131)67:5<815::AIDAPP6>3.0.CO;2-W. 15. Vidhate, S., & D’Souza, N. A. (2011). Biodegradable poly(hydroxy butyrate-covalerate) nanocomposites and blends with poly(butylene adipate-co-terephthalate) for sensor applications, miscibility, compatibility of PBHV/PBAT blends, (Doctoral thesis). University of North Texas, Texas.

16. Bittmann, B., Bouza, R., Barral, L., Castro-Lopez, M., & Dopico-Garcia, S. (2015). Morphology and thermal behavior of poly (3-hydroxybutyrate-co-3- hydroxyvalerate)/poly(butylene adipate-co-terephthalate)/clay nanocomposites. Polymer Composites, 36(11), 2051-2058. http://dx.doi.org/10.1002/ pc.23115.

Received: Nov. 19, 2017 Revised: May 19, 2018 Accepted: June 09, 2018

Polímeros, 29(1), e2019006, 2019

Non-isothermal melt crystallization kinetics of poly(3-hydroxybutyrate), poly(butylene adipate-co-terephthalate) and its mixture Appendix A. Crystallization parameters.

The following tables contain a series of parameters for the melt (C1) and cold (C2) crystallization events, and the second melting event (F2), obtained by point-by-point integration of the raw time-temperature-heat flow data reported by the DSC. ϕ (°C/min): heating/cooling rate.

T0.1% (°C): temperature for 0.1% molten/crystallized fraction (a good estimate of the initial point of the event).

T50% (°C): temperature for 50% molten/crystallized fraction (a better estimate of the “characteristic” melting/ crystallization temperatures than the peak values, in particular for events represented by complex peaks) T99.9% (°C): temperature for 99.9% molten/crystallized fraction (a good estimate of the final point of the event) Tp, (°C): peak melting/crystallization temperature. cmax (min-1): maximum melting/crystallization rate. c5-95% (min-1): mean melting/crystallization rate for the central 90% of the polymer. τ½ (min): half crystallization time (time to reach 50% crystallized fraction from the start of the event) ΔH (J/g): latent heat of melting/crystallization ΔX (%): change in crystallinity during the event (estimated from the latent heats).

Numbers in italic correspond to a minor peak.

Table A1. PHB: melt crystallization (C1). ϕ (°C/min) 2 4 6 8 12 16 24 32 48 64

T0,1% 103.6 101.2 99.6 105.7 87.7 100.2 88.0

T50%

T99,9%

90.0 78.3 66.1 80.9 61.8 66.2 58.3

(°C) 78.4 51.2 31.4 50.4 37.7 40.7 32.6

ΔT 25.1 49.9 68.3 55.4 50.0 59.5 55.5

Tp

τ½

89.2 79.8 64.8 78.8 62.2 64.9 53.8

(min) 6.82 5.73 5.59 3.11 2.16 2.12 1.24

cmax

c5-95%

ΔH

ΔX

0.117 0.131 0.135 0.227 0.331 0.376 0.544

(J/g) 67.4 60.1 35.2 58.7 10.0 8.6 10.8

(%) 46.2 41.2 24.1 40.2 6.9 5.9 7.4

(min−1) 0.169 0.202 0.216 0.337 0.470 0.520 1.433

(event not observed)

Table A2. PHB: cold crystallization (C2). ϕ (°C/min) 2 4 6 8 12 16 24 32 48 64

T0,1%

T50%

T99,9%

ΔT

(°C)

Tp

τ½

cmax

(min)

c5-95% (min−1)

ΔH

ΔX

(J/g)

(%)

(event not observed) 38.1

48.8

54.4

18.3

44.3 43.8 52.1 51.2 55.0 58.3

57.3 57.6 62.6 69.4 72.5 80.5

66.8 67.7 69.5 96.4 86.0 101.2

22.5 23.9 17.3 45.2 30.9 42.9

Polímeros, 29(1), e2019006, 2019

49.6 1.77 (event not observed) 58.1 1.07 58.4 0.85 63.4 0.43 69.7 0.56 73.4 0.36 80.9 0.34

0.853

0.547

15.2

10.4

1.390 1.802 2.798 2.392 3.506 3.382

0.896 1.159 1.920 1.289 2.362 2.270

29.1 33.3 32.2 48.6 36.8 43.9

19.9 22.8 22.0 33.3 25.2 30.1

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Costa, A. R. M., Ito, E. N., Cavalho, L. H., & Canedo, E. L. Table A3. PHB: melting (F2). ϕ

T0.1%

T50%

ΔT

T99.9%

(°C/min)

Tp

cmax

ΔH

ΔX

(J/g)

(%)

0.121

64.2

44.0

0.183

76.4

52.3

0.261

79.2

54.2

0.341

73.7

50.4

0.250 0.371 0.499 0.947 1.223 1.724

82.3 83.7 80.7 69.5 70.0 67.8

56.4 57.4 55.3 47.6 48.0 46.4

c5-95%

(°C)

(min−1)

2

148.3

169.6

173.7

25.4

4

138.6

170.0

175.4

36.8

6

137.0

168.4

175.4

38.4

8

139.3

167.8

175.3

36.1

12 16 24 32 48 64

97.0 99.4 102.4 128.1 117.3 127.8

166.6 165.6 165.5 164.6 162.0 163.1

176.5 176.9 180.8 180.2 180.3 183.7

79.6 77.5 78.4 52.1 62.9 55.9

160.2 170.8 161.2 172.5 170.8 160.6 171.1 170.5 170.0 169.9 168.1 166.7 166.9

0.066 0.323 0.103 0.524 0.618 0.315 0.836 0.970 1.202 1.417 1.737 2.331 3.100

Table A4. PBAT: melt crystallization (C1). ϕ

T0.1%

T50%

(°C/min)

T99.9%

ΔT

(°C)

(min) (min 89.5 0.105 25.5 7.05 81.3 0.111 85.7 0.089 30.4 4.05 67.7 0.135 27.7 79.8 2.61 0.546 25.4 75.1 1.89 0.733 31.1 71.1 1.44 1.008 32.5 70.4 1.06 1.278 44.2 61.8 0.97 1.555 30.2 58.4 0.52 2.263 26.1 36.1 0.33 3.344 (isothermal crystallization @ 25°C)

2

105.7

91.6

80.2

4

102.2

86.0

71.8

6 8 12 16 24 32 48 64

96.4 91.4 89.6 88.3 85.7 75.9 53.2

80.8 76.3 72.4 71.4 62.4 59.2 37.6

68.7 65.9 58.5 55.8 41.6 45.7 27.2

τ½

Tp

cmax

ΔH

ΔX

(J/g)

(%)

0.119

11.8

10.4

0.214

13.2

11.6

0.357 0.489 0.654 0.839 0.868 1.544 2.472

14.6 14.9 16.6 15.8 25.1 15.7 5.77

12.8 13.1 14.6 13.9 22.0 13.8 5.1

c5-95% )

−1

Table A5. PBAT: melting (F2). ϕ (°C/min) 2 4 6 8 12 16 24 32 48 64

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T0.1% 100.4 94.4 92.8 86.9 86.2 85.7 83.8 82.6 84.4 85.2

T50%

T99.9%

123.1 121.6 120.7 118.5 118.7 118.5 117.6 117.0 119.6 121.1

(°C) 140.0 146.4 141.6 141.6 144.1 143.3 142.0 142.3 146.8 149.1

ΔT

Tp

cmax

c5-95%

ΔH

ΔX

0.065 0.113 0.167 0.204 0.298 0.395 0.573 0.742 1.090 1.430

(J/g) 7.2 9.3 10.0 11.6 12.4 11.8 13.7 13.0 14.1 11.4

(%) 6.3 8.2 8.8 10.2 10.9 10.4 12.0 11.4 12.4 10.0

(min−1) 39.6 52.0 48.9 54.8 58.0 57.6 58.2 59.7 62.4 63.9

125.7 124.2 123.2 121.2 121.3 121.5 120.7 120.0 122.4 123.8

0.093 0.166 0.245 0.301 0.439 0.588 0.837 1.097 1.594 2.112

Polímeros, 29(1), e2019006, 2019

Non-isothermal melt crystallization kinetics of poly(3-hydroxybutyrate), poly(butylene adipate-co-terephthalate) and its mixture Table A6. PHB/PBAT blend: melt crystallization (C1). ϕ

T0.1%

T50%

(°C/min)

T99.9

ΔT

Tp

(°C)

2

108.1

86.3

66.8

4

103.0

70.5

39.4

6

102.7

75.2

32.1

8 12 16 24 32 48 64

92.8 89.1 90.4 87.4 69.8 49.9

77.1 73.9 67.6 63.7 55.9 36.0

65.6 61.5 52.4 41.7 42.5 27.0

τ½

cmax

(min)

ΔH

ΔX

(J/g)

(%)

0.075

36.9

26.0

0.086

34.4

24.1

0.132

31.3

21.9

0.446 0.684 0.751 0.746 1.668 2.768

9.1 7.1 10.2 14.0 6.4 1.8

6.6 5.0 7.1 9.8 4.5 1.2

c5-95% (min−1)

89.5 0.105 10.90 81.3 0.111 85.7 0.089 63.6 8.14 67.7 0.133 80.4 0.238 70.6 4.59 71.3 0.158 27.2 76.1 1.96 0.681 27.5 73.0 1.27 1.020 38.0 66.1 1.42 1.203 45.7 63.4 0.99 1.263 27.3 55.1 0.43 2.362 22.9 34.2 0.29 3.726 (isothermal crystallization @ 25°C) 41.3

Table A7. PHB/PBAT blend: cold crystallization (C2). ϕ

T0.1%

(°C/min) 2 4 6 8 12 16 24 32 48 64

T50%

T99.9%

ΔT

(°C) 34.6 35.8 39.4 36.3 37.0 45.9 43.3 41.8 45.9

42.4 41.8 49.0 49.4 50.1 57.8 57.4 58.4 64.6

49.4 48.2 57.8 57.2 61.4 68.5 70.4 72.1 83.1

14.8 12.4 18.4 18.9 24.5 22.5 27.0 30.3 37.1

Tp

τ½

(min) (evento não observado) 42.9 1.70 42.1 1.01 49.4 1.21 49.9 0.92 50.6 0.81 58.2 0.49 57.9 0.43 59.1 0.34 65.0 0.29

cmax

c5-95% (min−1)

0.690 1.041 1.093 1.667 1.859 2.708 3.080 4.011 4.389

0.449 0.722 0.714 1.072 1.177 1.763 2.006 2.607 2.793

ΔH

ΔX

(J/g)

(%)

2.4 3.9 8.7 18.9 16.2 21.2 20.4 20.0 20.5

1.7 2.7 6.3 13.3 11.3 14.9 14.3 14.0 14.3

Table A8. PHB/PBAT bend: melting (F2). ϕ

2

4

6 8 12 16 24 32 48 64

T0.1%

ΔT

T50%

T99.9%

98.1

123.3

(°C) 139.8

41.7

PHB

152.9

170.1

174.6

21.7

PBAT

97.5

121.9

137.5

40.0

PHB

145.2

169.3

174.7

29.5

PBAT

97.6

120.0

134.8

37.2

PHB

145.4

168.1

174.1

28.7

PBAT PHB PBAT PHB PBAT PHB PBAT PHB PBAT PHB PBAT PHB PBAT PHB

105.7 144.1 94.3 142.0 92.6 141.6 95.0 142.4 100.2 139.2 95.4 137.1 100.0 139.3

122.1 168.2 119.1 167.2 117.0 165.9 118.8 167.1 117.0 164.6 116.6 162.5 118.6 164.2

137.8 174.5 135.1 175.5 133.7 174.7 134.3 178.3 130.6 177.4 129.6 177.1 133.2 181.3

32.1 29.9 40.8 32.5 41.2 33.1 39.2 35.8 30.5 38.2 34.2 40.0 33.3 42.0

(°C/min) PBAT

Polímeros, 29(1), e2019006, 2019

Tp

cmax

c5-95%

ΔH

ΔX

0.063

(J/g) 7.5

(%) 6.6

0.132

64.6

44.3

0.130

7.3

6.4

0.223

81.8

56.1

0.206

3.1

2.7

0.322

37.4

25.6

0.323 0.430 0.386 0.559 0.503 0.750 0.817 1.059 1.319 1.273 1.875 1.863 2.587 2.302

6.0 58.5 7.27 71.8 8.5 61.4 6.55 69.5 4.6 65.5 4.3 62.4 3.8 56.9

5.3 40.1 6.4 49.2 7.4 42.0 5.7 47.6 4.0 44.8 3.8 42.8 3.3 39.0

(min−1) 126.5 160.5 171.4 124.6 152.2 171.2 122.0 159.8 170.3 122.3 170.4 121.6 169.6 118.6 168.4 120.7 169.9 118.4 167.7 117.4 166.1 119.3 167.6

0.092 0.041 0.320 0.185 0.017 0.512 0.294 0.152 0.728 0.427 0.963 0.560 1.175 0.708 1.550 1.113 2.116 1.686 2.228 2.494 3.120 3.420 3.404

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Costa, A. R. M., Ito, E. N., Cavalho, L. H., & Canedo, E. L. Table A9. Pseudo-Avrami parameters for PHB (C1). ϕ (°C/min) 2 4 6 8 12 16 24

-6.4074 -6.9572 -6.6818 -4.2084 -2.8648 -3.4311 -0.8769

ln K ± 0.0203 0.0080 0.0149 0.0043 0.0080 0.0048 0.0105

% 0.3 0.1 0.2 0.1 0.3 0.1 1.2

3.0333 3.5184 3.4410 3.1465 2.9683 3.5135 2.3405

n ± 0.0113 0.0045 0.0087 0.0038 0.0099 0.0060 0.0186

% 0.4 0.1 0.3 0.1 0.3 0.2 0.8

3.1626 3.6935 3.9119 3.8032 4.0535 4.1928 3.8443 3.2902 3.2293

n ± 0.0121 0.0146 0.0238 0.0269 0.0309 0.2260 0.0426 0.0461 0.3630

% 0.4 0.4 0.6 0.7 0.8 5.4 1.1 1.4 11.2

3.4649 3.1179 3.4215 3.6864 3.8235 4.9705 3.2684 2.8133 2.9836

n ± 0.0052 0.0084 0.0194 0.0218 0.0216 0.0445 0.0241 0.3190 0.0528

% 0.2 0.3 0.6 0.6 0.6 0.9 0.7 11.3 1.8

3.3072 3.0466 3.0423 3.2783 3.4340 3.5352 2.3021 1.7803 2.4581 1.7609 0.9376 0.5830

m ± 0.2530 0.2383 0.0667 0.0098 0.0960 0.0666 0.4836 0.6470 0.5786 0.4251 0.6106 0.6921

% 7.6 7.8 2.2 0.3 2.8 1.9 21.0 36.3 23.5 24.1 65.1 118.7

R2 0.993 0.999 0.997 0.999 0.998 0.999 0.993

Table A10. Pseudo-Avrami parameters for PBAT (C1). ϕ (°C/min) 2 4 6 8 12 16 24 32 48

-6.7143 -5.9399 -4.5618 -3.1850 -2.3419 -1.3028 -0.6570 1.2063 2.9628

ln K ± 0.0221 0.0197 0.0222 0.0173 0.0141 0.0079 0.0179 0.0383 0.0494

% 0.3 0.3 0.5 0.5 0.6 0.6 2.7 3.2 1.7

R2 0.9930 0.996 0.994 0.994 0.995 0.998 0.992 0.992 0.997

Table A11. Pseudo-Avrami parameters for PHB/PBAT blend (C1). ϕ (°C/min) 2 4 6 8 12 16 24 32 48

-9.0492 -7.2279 -6.3746 -3.2308 -1.7596 -3.0387 -0.6115 1.8591 3.1936

ln K ± 0.0122 0.0176 0.0333 0.0150 0.0089 0.0203 0.0113 0.0358 0.0795

% 0.1 0.2 0.5 0.5 0.5 0.7 1.8 1.9 2.5

R2 0.998 0.995 0.987 0.995 0.997 0.993 0.996 0.996 0.994

Table A12. Ozawa parameters for PHB (C1).

T (°C) 100 95 90 85 80 75 70 65 60 55 50 45

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-1.3041 0.4196 1.7162 3.0437 4.1058 4.9685 3.6428 3.3463 4.9965 4.3003 2.7860 2.4554

ln κ ± 0.3460 0.3259 0.0912 0.0134 0.1313 0.1294 0.9389 1.2561 1.3386 0.9834 1.5562 1.7639

% 26.5 77.7 5.3 0.4 3.2 2.6 25.8 37.5 26.8 22.9 55.9 71.8

R2 0.994 0.994 0.999 1.000 0.999 0.999 0.956 0.883 0.900 0.896 0.702 0.415

Polímeros, 29(1), e2019006, 2019

Non-isothermal melt crystallization kinetics of poly(3-hydroxybutyrate), poly(butylene adipate-co-terephthalate) and its mixture Table A13. Ozawa parameters for PBAT (C1).

T (°C) 95 90 85 80 75 70

1.7164 3.1909 4.6403 4.1889 5.5220 5.8429

ln κ ± 1.6134 1.4853 1.3316 0.5798 0.4380 0.1053

% 94.0 46.5 28.7 13.8 7.9 1.8

3.8437 3.8391 3.8933 2.6476 2.6492 2.2798

m ± 1.1796 0.9424 0.7418 0.3230 0.2215 0.0493

% 30.7 24.6 19.1 12.2 8.4 2.2

3.1229 3.1761 2.7735 2.6438 3.5234 2.7207 1.8479 1.9783

m ± 2.3441 0.0576 0.3388 0.3965 0.7064 0.4909 0.5894 1.4165

% 75.1 1.8 12.2 15.0 20.0 18.0 31.9 71.6

1.9880 1.8737 1.8010 1.7500 1.7113 1.6780 1.6477 1.6184 1.5902

α ± 0.2017 0.1836 0.1684 0.1558 0.1450 0.1350 0.1279 0.1256 0.1304

% 10.1 9.7 9.3 8.9 8.4 8.0 7.7 7.7 8.2

0.9512 0.9517 0.9511 0.9497 0.9478 0.9451 0.9414 0.9359 0.9257

α ± 0.0439 0.0362 0.0316 0.0284 0.0261 0.0246 0.0240 0.0253 0.0328

% 4.6 3.8 3.3 2.9 2.7 2.6 2.5 2.7 3.5

1.2073 1.1946 1.1964 1.1957 1.1995 1.2096 1.2143 1.2188 1.2268

α ± 0.1191 0.1051 0.1047 0.1031 0.1005 0.0968 0.0971 0.1005 0.1064

% 9.8 8.7 8.7 8.6 8.3 8.0 7.9 8.2 8.6

R2 0.914 0.893 0.902 0.957 0.986 0.999

Table A14. Ozawa parameters for PHB/PBAT blend (C1).

T (°C) 85 80 75 70 65 60 55 50

3.4662 5.5473 5.9673 6.8120 10.0566 8.7700 6.7435 8.1549

ln κ ± 5.7715 0.1533 0.9008 1.0544 2.2551 1.5671 1.9721 4.9812

% > 100 2.8 15.1 15.5 22.4 17.9 29.2 61.1

R2 0.6396 0.9993 0.9710 0.9569 0.8924 0.9110 0.8309 0.6611

Table A15. Mo parameters for PHB (C1). x (%) 10 20 30 40 50 60 70 80 90

4.7575 4.7716 4.7712 4.7799 4.7960 4.8127 4.8337 4.8634 4.9201

lnF ± 0.4920 0.4479 0.4109 0.3802 0.3537 0.3293 0.3120 0.3064 0.3181

% 10.3 9.3 8.6 7.9 7.3 6.8 6.4 6.3 6.4

R2 0.979 0.981 0.982 0.984 0.985 0.987 0.988 0.988 0.986

Table A16. Mo parameters for PBAT (C1). x (%) 10 20 30 40 50 60 70 80 90

2.2293 2.4551 2.5847 2.6773 2.7512 2.8146 2.8727 2.9304 2.9966

lnF ± 0.1148 0.0946 0.0826 0.0743 0.0683 0.0643 0.0628 0.0661 0.0858

% 5.1 3.8 3.1 2.7 2.4 2.2 2.1 2.2 2.8

R2 0.987 0.991 0.993 0.994 0.995 0.996 0.996 0.995 0.992

Table A17. Mo parameters for PHB/PBAT blend (C1). x (%) 10 20 30 40 50 60 70 80 90

3.1676 3.3480 3.4894 3.5855 3.6797 3.7850 3.8699 3.9569 4.0688

Polímeros, 29(1), e2019006, 2019

lnF ± 0.3013 0.2659 0.2649 0.2607 0.2542 0.2448 0.2456 0.2543 0.2691

% 9.5 7.9 7.5 7.2 6.9 6.4 6.3 6.4 6.6

R2 0.966 0.970 0.970 0.971 0.972 0.975 0.975 0.973 0.970

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Costa, A. R. M., Ito, E. N., Cavalho, L. H., & Canedo, E. L. Appendix B. Differential form of the Mo model.

The integral form of Mo model is given by Equation (11). which expresses the cooling/heating rate ϕ as a function of time τ measured from the start of the crystallization event. at constant relative crystallinity x:

φ = F τ−α

(A1)

F = F(x) and α = α(T) are the model parameters. In this case. x = x(ϕ.τ); thus:

c =

∂x  dx  ≡  ∂τ  d τ φ

(A2)

To estimate the crystallization rate it is convenient to start with:

F =τα φ

(A3)

For constant Mo exponent α. differentiation at constant cooling/heating rate :

∂F dF  dx  α−1 =   = ατ φ ∂τ dx  d τ φ

(A4)

and elimination τ between Equations (A3) e (A4):

dF  dx  1/ α 1−1/ α

  = αφ F dx  d τ φ

(A5)

or

 dx   d ln F  c =   = αφ1/ α F −1/ α    d τ φ  dx 

−1

(A6)

Equation (A6) is the differential form of Mo for constant exponent α. Equation (A6) may be expressed as a product:

c = k (φ) ⋅ f ( x)

(A7)

k = αφ1/ α

(A8)

of a pseudo kinetic constant (function of ϕ only):

and a function f(x) that depends on conversion only:

 dF  f = F 1−1/ α    dx 

−1

(A9)

For crystallinity-dependent Mo exponent. α = α(x). differentiation of Equation (A3):

d α  dx   α dF  dx  d τα =τα φ  + ln τ     =φ τ dx  d τ φ  dx  d τ φ dτ 

(A10)

or

ατα−1φ  dx  c = =  dF dα  d τ φ − τα ln τ dx dx

(A11)

The last step consists in the elimination of τ and lnτ between Equations (A11) e (A3); a compact form of the final results is:

αφ1/ α F −1/ α  dx  c = =   d τ φ d ln F + 1 − ln F  d ln α   dx  ln α  dx

(A12)

Equations (A12) is a differential form of Mo model for exponent α = α(x). It is not possible to define a kinetic constant in this case.

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Polímeros, 29(1), e2019006, 2019

ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.09717

Rheological properties of low-density polyethylene filled with hydrophobic Co(Ni)-Al layered double hydroxides Silvia Jaerger1, Andreas Leuteritz2, Rilton Alves de Freitas1 and Fernando Wypych1*  1

Departamento de Química – DQUI, Universidade Federal do Paraná – UFPR, Curitiba, PR, Brasil 2 Leibniz-Institut für Polymerforschung Dresden e.V. – IPF, Dresden, Germany *wypych@ufpr.br

Abstract Cobalt/aluminum and nickel/aluminum layered double hydroxide (LDH - M+2:Al molar ratio of 3:1) were intercalated with dodecylsulphate (DDS), laurate (LAU), stearate (STE) and palmitate (PAL) and used as filler in low-density polyethylene (LDPE) in percentages between 0.2 and 7.0wt%. After injection molding, the samples were submitted to morphological characterization by scanning electron microscopy (SEM), analysis of thermal behavior by differential scanning calorimetry (DSC) and investigation of rheological properties. All Co/Al-LDPE samples showed the formation of a high temperature polymer crystal domain, induced by the LDH filler. The rheological properties indicated in general a reduction of shear modulus due to incompatibility between some regions of LDH and LDPE, which promoted phase separation. However, interaction with the LDH surface indicated higher affinity of the Ni/Al-LDH for the LDPE compared to Co/Al-LDH, forming permanent networks.

Keywords: layered double hydroxide, low-density polyethylene, nanocomposites, rheological properties.

How to cite: Jaerger, S., Leuteritz, A., Freitas, R. A., & Wypych, F. (2019). Rheological properties of low-density polyethylene filled with hydrophobic Co(Ni)-Al layered double hydroxides. Polímeros: Ciência e Tecnologa, 29(1), e2019007. https://doi.org/10.1590/0104-1428.09717

1. Introduction Low-density polyethylene (LDPE) is widely used in the production of packaging films, but its use can be expanded by improving its properties. Polymeric composites using natural and synthetic layered fillers can be interesting alternatives to improve some polymer properties, especially mechanical, gas barrier, flame retardancy, etc[1-5].

Polymer nanocomposites (PNC) are materials in which the matrix is a polymer and the filler has at least one dimension of nanometer magnitude, normally smaller than 100 nm[6]. Layered compounds are composed of two dimensional nanosized layers, which are stacked along the basal axis, being ideal for use as fillers in PNC[7]. Normally, hydrophilic layered compounds can become hydrophobic after the intercalation of specific organic anionic species.

The most common are the carboxylates and sulphonates[8,9].

The majority of studies devoted to PNC use clay minerals of the 2:1 groups cation exchangers, but they present many disadvantages, which can be overcome by the use of synthetic layered compounds[3-5,10,11].

Layered double hydroxide (LDH) having the general formula [M+21-xM+3x(OH)2]x+[Am-]x/m.yH2O, where M2+ and M3+ represent di and trivalent metals, respectively occupying octahedral sites coordinated by hydroxyl groups and Am-(H2O)y an intercalated hydrated anion, have been recently proposed as alternative materials for use as fillers in the formulation of PNC[5,10,11]. LDHs can be obtained with variable chemical composition, and to overcome the hydrophilic properties of

Polímeros, 29(1), e2019007, 2019

LDH, long carbon chain anionic species are intercalated, such as those derived from ethylene diaminetetraacetic acid[12] , aromatic acid[13], dihydroxybenzoic acid[14], chromotropic acid[15], dodecylsulphonic acid[16], stearic acid[17], oleic acid[18], and dodecylbenzenesulphonic acid[19].One extra advantage of using LDH is the possibility to impart color to the polymeric matrix, where the intercalated anions or layers of structural cations can be colored[20,21]. This will avoid the use pigments or dyes as polymer colorants.

To date, very few studies have been published described the rheology of LDH-polymer (PE) interactions[22-32].

For example, Costa et al.[22,23] studied the rheological properties of Mg/Al based layered double hydroxide-polyethylene and polyethylene grafted with maleic anhydride and demonstrated only strong interactions between the LDH with grafted maleic anhydride matrix. The authors also observed some structural aggregation, ruptured by shearing but reformed after resting.

In our previous study[33] , we described the preparation of a LDPE nanocomposite containing organo-modified Ni/Al and Co/Al-LDH using melt extrusion. The mechanical properties, transmission electron micrographs and differential scanning calorimetric results showed that the addition of these LDHs had a small influence on the polymer properties, even in the case of good dispersions and exfoliation of the LDH layered crystals.

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

Jaerger, S., Leuteritz, A., Freitas, R. A., & Wypych, F. To date, very few studies have been published described the rheology of LDH-polymer (PE) interactions[22-32].

For example, Costa et al.[22,23] studied the rheological properties of Mg/Al based layered double hydroxide-polyethylene and polyethylene grafted with maleic anhydride and demonstrated only strong interactions between the LDH with grafted maleic anhydride matrix. The authors also observed some structural aggregation, ruptured by shearing but reformed after resting.

In the present study, morphological, thermal and rheological experiments were conducted on LDPE nanocomposites containing hydrophobic Ni/Al-LDH and Co/Al-LDH as filler. Rheological analyses investigate the linear viscoelastic properties of the nanocomposites and proved to be a good technique to understand more clearly the nature of particle dispersion in PNC. Under the linear viscoelastic regime, rheological behavior is very sensitive to any change in microstructure of composites when the shearing is carried out under very low shear rates, and it can provide information that can explain the clay particle state dispersion and possible particle-particle and particlepolymer interactions. These possibilities for improvement will be discussed in this manuscript.

2. Materials and Methods 2.1 Materials LDPE, PB608 (MFI 30 g/10 min and density 0.915 g/cm3) was a resin free of additives and was kindly donated by Braskem (Brazil). Hydrophilic Ni/Al-LDH and Co/Al-LDH (M+2:M+3 molar ratio of 3:1), intercalated with the anions carbonate, chloride and nitrate, were prepared by hydrolysis and the hydrophobic materials, intercalated with the anions laurate (LAU), palmitate (PAL), dodecylsulphate (DDS) and stearate (STE), were prepared by ion exchange method and characterized by different techniques, as describe in the previous article[33]. Briefly, first the carbonate derivatives were prepared by urea hydrolysis, serving as matrix for the preparation of the chloride derivative, obtained by treating the sample in an acetate buffer solution containing NaCl. Next, chloride anions were replaced by the specific organic anions by exchange reactions.

2.2 Instrumental techniques The LDH samples were characterized and were added to LDPE, using different percentages in relation to the LDPE mass (0.2, 0.5, 2.0, 5 and 7wt%)[33]. The mixture was fed into a HAAKE MiniLab II micro compounder and then specimens were prepared of each composition by melting the LDPE/filler mixtures at 160 bar, 130 °C for 5 minutes.

The material was then injected at 320 bar into a mold cavity (at 40 °C) to produce ASTM D638 samples (type IV) in a HAAKE MiniJet II injector. For statistical purposes, ten specimens were evaluated for each composition and average values are presented, excluding outliers.

The samples for the scanning electron microscopy (SEM) measurements were obtained by dipping the composites in liquid N2, followed by fracturing and placement on aluminum stubs with conductive glue for gold sputtering.

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The images were recorded with a Tescan VEGA3 LMU microscope at 10 kV. Differential scanning calorimetry (DSC) analysis of the nanocomposites was performed with a Netzsch DSC 200 calorimeter under nitrogen atmosphere using the following running cycle: heating from room temperature to 180 °C (at 10 °C.min–1), stabilization for 5 min at 150 °C, cooling to –80 °C (at 10 °C.min–1), and heating to 150 °C (at 10 °C.min–1). After eliminating the thermal history of the PNC, only the first cooling and second heating curves were recorded.

Rheological studies were performed with the aid of an ARES rheometer from Rheometrics Scientific, USA, using parallel plates. The torque transducers varied from 0.02-2000 g.cm-1 and the temperature of the sample was controlled in a closed chamber under nitrogen atmosphere.

The samples were pressed and molded with diameter of 25 mm and thickness of 1.5 mm. The dynamic oscillatory shear experiment occurred at 190 °C with a frequency sweep of 0.063-100 rad.s-1, where the amplitude was maintained at 10%, which is the range for viscoelastic polymers.

3. Results and Discussions In order to gain a better understanding of the LDH-LDPE nanocomposites’ behavior, SEM (Figure 1) and DSC (Figure 2) measurements were performed with all the Ni/Al and Co/Al samples. After, a complete rheological description of the interactions between LDH-LDPE will be presented (Figure 3-8 and Appendix A).

The SEM images (Figure 1) are shown with lower magnification to have an idea of the whole sample surface and with higher magnification to note the surface details. In the SEM images of the samples of Ni/Al-LDPE (Figure 1a-d), poor dispersion of the filler is observed in micrometric scale, with segregation of agglomerates consisting basically of LDH (see white particles in Figure 1a, c). These aggregates were sometimes detached from the polymeric matrix during cryogenic fracture, as observed in Figure 1b, d, leaving large depressions (layered particles pull-out).

The sample Ni/Al-STE (Figure 1b) presented the best dispersion, since no agglomerations were detected, even in the low magnification image, suggesting low phase separation of LDH Ni/Al-STE and LDPE. In the SEM images of the samples of Co/Al-LDPE (Figure 1e-h), relatively homogeneous phases were observed in micrometric scale, with the segregation of some particles rich in LDH but embedded in the LDPE matrix, in which LDH was also dispersed. This was more frequently observed in the cases of Co/Al-LAU (Figure 1g), Co/Al-PAL (Figure 1h) and Co/Al-STE (Figure 1f). In the Co/Al-STE phase, the XRD patterns also indicated the delamination/exfoliation of the layered crystals in the polymeric matrix[33].

Since these layered particles were particularly evident in the Co/Al-LDPE samples, we suppose this can be attributed to those observed in high melting/crystallization peaks by DSC (Figure 2a, b), also exclusive for the same samples.

As can be observed for Co/Al-STE (Figure 2a), the first cooling presented two exothermic peaks, at 89 and 109°C. Polímeros, 29(1), e2019007, 2019

Rheological properties of low-density polyethylene filled with hydrophobic Co(Ni)-Al layered double hydroxides The first is attributed to the crystallization of LDPE[31] and the second is related to the formation of crystalline domains at high temperature, with the influence of the filler, which acts as a nucleating agent[32]. This high temperature phase

occurs due to the effective interaction between polymeric chains and the LDH particles surface, where different crystal sizes or even different polymorphic phases of HDPE are obtained, due to different chain conformations.

Figure 1. SEM images of cryogenic fracture for the LDPE nanocomposites containing 7% of the fillers: Ni/Al-DDS (a), Ni/Al-STE (b), Ni/Al-LAU (c), Ni/Al-PAL (d), Co/Al-DDS (e), Co/Al-STE (f), Co/Al-LAU (g), and Co/Al-PAL (h). The LDH crystals can be seen as white plates.

PolĂ­meros, 29(1), e2019007, 2019

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Jaerger, S., Leuteritz, A., Freitas, R. A., & Wypych, F. In the second melting, three endothermic events were observed, at 102, 123 and 128 °C. The first is attributed to the melting of LDPE[31] and the second and third to the melting of the phases obtained at high temperature, induced by the LDH filler[11]. In the phase of Co/Al-PAL (Figure 2b), similar behavior was observed with the same position of the melting/crystallization temperatures (Table 1).

One important aspect to be observed is that the proportion between the regular phase of LDPE and the high temperature phase was almost constant with rising filler content. This can be seen when the melting and crystallization enthalpies of both phases are compared (Table 1). Similar high temperature/ low temperature phase ratios were observed for the samples Co/Al-PAL and Co/Al-STE for crystallization and melting.

Another observation is the reduction of the enthalpies compared to the neat LDPE and with increasing content of both fillers, indicating reduction of LDPE chain mobility, hindering formation of polymer crystal domains. This can be attributed to the good interaction between the two phases.

In fact, both evaluated samples did not show any diffraction peak in the X-ray diffraction pattern, even with 4% filler, indicating delamination/exfoliation (data not shown). The DSC splitting of the crystallization/melting peaks were not observed for the samples containing Ni/Al-LDPE, independent of the intercalated organic anion, as reported previously[33].

Figure 2. DSC curves of the Co/Al-STE (a) and Co/Al-PAL (b) LDPE nanocomposites containing 0.2, 0.5 and 7% of the fillers.

The LDPE nanocomposites filled with different proportions of hydrophobic LDH (Ni/Al-LDH and Co/Al-LDH

Figure 3. Frequency dependence of the shear modulus G’ and G” of Ni/Al-LAU (a1 and a2) and Ni/Al-STE (b1 and b2) (0.2, 0.5, 2.0, 5.0 and 7wt%) in LDPE at 190 °C. 4/13

Polímeros, 29(1), e2019007, 2019

Rheological properties of low-density polyethylene filled with hydrophobic Co(Ni)-Al layered double hydroxides intercalated with the anions DDS, LAU, PAL or STE), were also evaluated by shear modulus frequency dependence using rheological analysis at 190 °C. In the first part of the results and discussion section, Ni/Al-LDH filler will be presented while in the second part, the focus will be on Co/Al-LDH and LDPE mixtures, for comparative purposes. The elastic modulus (G’) and viscous modulus (G”) of Ni/Al-LAU and Ni/Al-STE are presented in Figure 3a and 3b, respectively.

The shear modulus for Ni/Al-LDH DDS and PAL are presented in the Appendix A (Figure A1).

At low LDH concentration, no significant effect on LDPE shear moduli was observed. However, with the increase of concentration (≥ 2 wt%) for all Ni/Al-LDH in LDPE, a

reduction of G’ and G” values was observed compared to neat LDPE (without any filler). This reduction of G’ and G” was not expected, and the shear moduli reduction at higher frequencies can be attributed to a phase separation processes.

Such hypothesis will be investigated below.

To improve the comparison between all the fillers and concentrations with LDPE, master curves were constructed, using a vertical shift factor (av) obtained at high frequency range (~100 rad.s-1). The shift factor was used on both G’ and G” raw values (Figure 4). The av values were obtained by the ratio of LDPE with the shear modulus of LDPE-LDH mixtures. This high frequency was chosen to calculate av

Figure 4. Master curves of the frequency dependence of G’ (1) and G” (2) for LDPE filled with Ni/Al-LDH, intercalated with DDS (a1, a2), STE (b1, b2), LAU (c1, c2) and PAL (d1, d2), obtained by the superposition of different LDH loadings, from 0 to 7 wt%. The reference used was neat LDPE. Polímeros, 29(1), e2019007, 2019

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Jaerger, S., Leuteritz, A., Freitas, R. A., & Wypych, F. Table 1. Neat LDPE and LDPE/LDH thermal properties obtained from DSC curves.

Melting Sample

Tm (°C) Tm1 (°C)

Tm2 (°C)

Crystallization ΔHm (J/g)

ΔHm1 (J/g) ΔHm2 (J/g)

Tc (°C) Tc1 (°C)

Tc2 (°C)

Neat LDPE 105 -88.2 88 Co/Al-STE 0.2% 103 128 -80.8 -14.7 89 109 Co/Al-STE 0.5% 103 128 -73.1 -13.1 89 108 Co/Al-STE 7% 103 128 -54.6 -13.6 88 108 Co/Al-PAL 0.2% 103 128 -81.7 -18.2 88 109 Co/Al-PAL 0.5% 103 128 -51.9 -9.7 89 107 Co/Al-PAL 7% 104 128 -40.7 -10.8 88 107 m = melting; c = crystallization; Peak 1 = low temperature; Peak 2 = high temperature.

due to the presence of transient networks, usually observed for polymers and characteristic of chain entanglements.

However, for LDPE or LDPE-LDH, no crossing between G’ and G” was observed, and G” > G’ throughout the frequency sweep, evidencing that in melted state no transient networks were identified for the LDPE at 190 °C. The absence of crossing could reflect the low molar mass of the melted polymers used, and the shift of the crossing between G’ and G” to higher frequencies indicated lower relaxation time (τ) for LDPE at 190 °C. For the fillers Ni/Al-STE and Ni/Al-PAL, almost the same av values were observed. Ni/Al-LAU LDPE nanocomposites presented a shift factor higher than the previous two fillers and the most significant av values were related to Ni/Al-DDS. The significant reduction of G’ and G” due to presence of LDH can be attributed to a phase separation process, as hypothesized above, with polymer concentration reduction in the bulk. Based on the av values calculated at higher frequencies, LDH and LDPE fractions were excluded from the LDPE matrix, and it was found that the phase separation can reduce the concentration of both LDH and LDPE in the bulk.

This LDPE-LDH phase separation was attributed to formation of aggregates that are segregated from the continuous phase of LDPE-LDH nanocomposites.

This effect will be better described in SEM and DSC measurements (Figure 1 and 2). Apparently, some regions of LDH interact strongly with the LDPE, but somehow this fraction separates from the continuous phase. Other possibilities are some incompatibility and inducement of crystallization of LDPE. The same aggregation of PE/LDH was demonstrated by Costa et al.[22,23], who reported that the grafting of polyethylene with maleic anhydride, increased the dispersion and rheological performance, mainly due to hydrophilic interactions.

Despite the phase separation observed, the hydrophobic segments of LDH and LDPE also formed permanent networks, from 0.0631 to 5 rad.s-1, since the elastic modulus at low frequencies increased, suggesting the formation of a new network, not observable in the matrix of LDPE, with longer relaxation time. This increase was also observed by Costa et al.

[22,23] , who attributed it to LDH-LDPE matrix interaction.

In the present case, two different events were observed: one responsible for a significant phase separation process, quantified by av, and a permanent network at lower frequencies.

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Peak ratio

ΔHc (J/g) ΔHc1 (J/g)

ΔHc2 (J/g)

c2/c1

m2/m1

74.5 50.1 48.1 48.4 54.4 39.6 39.5

14.7 12.7 14.0 17.8 9.80 11.5

0.29 0.26 0.29 0.33 0.25 0.29

0.18 0.18 0.25 0.22 0.19 0.27

The first one can be attributed to strong hydrophobic interactions between LDPE and LDH, creating nuclei for aggregation, followed by phase separation of the aggregates.

Another possibility is some hydrophilic incompatibility between LDH and matrix, due to lateral zone interactions of the layered crystals and LDPE, promoting phase separation.

A final hypothesis investigated by DSC is that LDH induces crystallization of the LDPE, with phase separation (Figure 2).

The permanent network can be associated with interaction of LDPE with the hydrophobic surfaces of LDH, observed as a rise at lower frequencies. This increase was directly related to the new network formed due to association of LDPE-LDH, and it is dependent on filler concentration for Ni/Al-LDH. This network was still present at higher temperatures (190 °C) and can be attributed to the hydrophobic interactions between the LDH and LDPE. A compilation of the vertical shift factors (av) used due to Ni/Al-LDH fillers in LDPE is observed in Appendix A (Figure A2).

To complement the permanent network characterization, the complex dynamic viscosity (η*) was calculated, as presented in Equation 4 and Figure 5. The complex viscosity (η*) (Figure 5) was calculated from G’ and G” after vertical factor (av) correction, using the following Equations 1 and 2. = G*

( G '.av )2 + ( G ".av )2

(1)

G* ( ωi ) (2) η* = ωi

where, G* is the complex shear modulus and ωi is the frequency i.

At low frequencies, the rise observed in the calculated η* was much more evident in Ni/Al-DDS (Figure 5a) and Ni/Al-LAU (Figure 5c), and the increase of viscosity was much more significant after av correction, demonstrating the formation of a network in LDPE induced by the filler.

This up-turn also appeared at higher frequencies with increasing LDH concentration, suggesting that the permanent network started to dominate the bulk phase.

To confirm the permanent network, aggregation and hydrophobic interactions induced by the filler for comparative purpose, another LDH (Co/Al), with the same intercalated ions (DDS, LAU, PAL and STE) was used. The G’ and G” of Co/Al-LDH with LAU and STE are presented in Figure 6.

Polímeros, 29(1), e2019007, 2019

Rheological properties of low-density polyethylene filled with hydrophobic Co(Ni)-Al layered double hydroxides

Figure 5. Complex viscosity (η*) obtained from G’av and G”av for LDPE filled with Ni/Al-LDH, intercalated with DDS (a), STE (b), LAU(c) and PAL (d), obtained by the superposition of different filler concentrations, from 0 to 7 wt%. The reference used was neat LDPE, at 190 °C.

Figure 6. Frequency dependence of the shear moduli G’ and G” of Co/Al-LAU (a1 and a2) and Co/Al-STE (b1 and b2) (0.2, 0.5, 2.0, 5 and 7wt%) in LDPE, at 190 °C.

The shear modulus for Co/Al-LDH DDS and PAL are presented in Appendix A (Figure A3).

The G’ and G” moduli were normalized, for comparative purposes, using av at high frequency (~100 rad.s-1) (Figure 7).

Shear moduli reduction was observed for Co/Al-LDH fillers in mixtures with LDPE, compared to Ni/Al-LDH fillers, since in both cases G’ and G” moduli of LDPE were smaller. However, the relation was not Co/Al-LDH dependent. All the fillers presented almost the same behavior at high frequencies and high LDH concentration (7 wt%).

For Co/Al-LDH with different intercalated ions, the up-turn at lower frequencies was much more evident than for Ni/Al-LDH, even at low Co/Al-LDH concentration (0.2-0.5 wt%). With relation to aggregates formed, the phase-separation from the continuous phase was much more evident for Co/Al-STE, even at the lowest concentration

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Jaerger, S., Leuteritz, A., Freitas, R. A., & Wypych, F.

Figure 7. Master curves of the frequency dependence of G’ and G” for LDPE filled with Co/Al-LDH, intercalated with ions of DDS (a1, a2), STE (b1, b2), LAU (c1, c2) and PAL (d1, d2), obtained by the superposition of different filler concentrations, from 0 to 7%.

The reference used was neat LDPE.

(0.2 wt%) (Figure 7). The up-turn observed at low frequency in the samples is consistent with the observation that the sample was delaminated/exfoliated in the polymeric matrix[33].

The higher phase separation of the Co/Al-LDH for the LDPE compared to Ni/Al-LDH was also evident due to lower concentration of LDH, which induced aggregation and shear moduli reduction, clearly observed by the shift-factors (Appendix A, Figure A4). Also, after normalization of G’ and G” by av, no evident differences were observed when using a higher concentration of LDH 8/13

(7 wt%) at higher frequencies (> 10 rad.s-1), suggesting that most of the LDPE-LDH phase separated, and the G’ and G” values of the continuous phase were almost independent of the filler at higher concentration.

In the lower frequency region (0.0631 to 5 rad.s-1), where the permanent network was observed for LDH and LDPE PNC, the samples with the strongest remaining networks in the continuous phase were Co/Al-LAU and Co/Al-PAL mixtures with LDPE (Figure 8c, d). All the samples at low Polímeros, 29(1), e2019007, 2019

Rheological properties of low-density polyethylene filled with hydrophobic Co(Ni)-Al layered double hydroxides

Figure 8. Complex viscosity (η*) obtained from G’av and G”av for LDPE filled with Co/Al-LDH, intercalated with DDS (a), STE (b), LAU (c) and PAL (d), obtained by the superposition of different filler concentrations, from 0 to 7 wt%. The reference used was neat LDPE.

concentration presented almost the same shear thinning behavior at very low frequencies, compared to LPDE samples.

The contrast between the SEM images and rheological clearly demonstrate phase separation, it can be a very powerful association technique to indicate LDH and polymer incompatibility and inhomogeneities.

The second event observed by DSC (Figure 2 and Table 1) can be attributed to hydrophobic interaction between LDH and LDPE also observed by rheology as an up-turn at very low frequency range (0.0631 to 5 rad.s-1). The phase-separated material was still present in the samples up to 190 °C, the temperature of the rheological analysis, suggesting melting of the aggregates at much higher temperatures. Apparently, the aggregates remaining at higher temperatures promoted some nucleation and LDPE organization, induced probably by hydrophobic interactions.

4. Conclusions Pink cobalt/aluminum and green nickel/aluminum layered double hydroxide (LDH - M+2: Al molar ratio of 3:1) were successfully intercalated with dodecylsulphate (DDS), laurate (LAU), stearate (STE) and palmitate (PAL) as pure phases.

First, the carbonate intercalated phases were obtained by urea hydrolysis, where carbonate was replaced by chloride anions from the acetate/NaCl buffer solution, followed by cation exchange reaction with the desired organic anions.

After characterization by several instrumental techniques, the nanometric thick layered crystals were used as fillers in low-density polyethylene (LDPE) in percentages between Polímeros, 29(1), e2019007, 2019

0.2 and 7 wt%, producing colored polymer nanocomposites by injection molding.

Scanning electron microscopy (SEM) images indicated that most of the samples were poorly dispersed in the polymeric matrix. Thermal behavior analyzed by differential scanning calorimetry (DSC) indicated that in the Ni/Al samples, the melting and crystallization were almost unchanged in relation to the neat LDPE, while in the samples Co/Al-PAL and Co/Al-STE, a second phase, and to a lesser extent, a third phase were observed, with higher crystallization and melting temperature than in the presence of neat LDPE. This high temperature phase is attributed to the formation of polymer crystals domain induced by the LDH filler particles.

The measured crystallization and melting enthalpies of both Co/Al-LDPE samples declined with increasing filler content, indicating a reduction of the crystallinity indices, caused by the LDH influence. Rheological properties indicated the formation of a permanent network, reducing shear moduli and hindering the filler dispersion, producing agglomerates, especially for the samples Ni/Al-LAU, Ni-Al-DDS, Co/Al-LAU, Ni/Co-PAL. In fact, as reported previously[33], only samples Ni/Al-STE and Co/Al-STE presented effective delamination/exfoliation.

5. Acknowledgements We thank Conselho Nacional de Desenvolvimento Científico e Tecnológico (CNPq) (Proc. 303846/2014-3), Coordenação de Aperfeiçoamento de Pessoal de Nível Superior (CAPES), Financiadora de Estudos e Projetos (FINEP) and Universidade Federal do Paraná (UFPR) for 9/13

Jaerger, S., Leuteritz, A., Freitas, R. A., & Wypych, F. financial support. S.J. also thanks CNPq – Ciências sem Fronteiras for the Ph.D. grant in the Leibniz-Institut für Polymerforschung Dresden, Germany, where part of this work was developed.

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Rheological properties of low-density polyethylene filled with hydrophobic Co(Ni)-Al layered double hydroxides morphological and rheological analysis of nanocomposites.

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polypropylene/clay hybrid materials. Macromolecule, 34(6), 1864-1872. http://dx.doi.org/10.1021/ma001122e.

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31. Zaman, H. U., & Beg, M. D. H. (2014). Influence of two novel compatibilizers on the properties of LDPE/organoclay nanocomposites. Journal of Polymer Engineering, 34(1), 1-9. http://dx.doi.org/10.1515/polyeng-2013-0144.

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33. Jaerger, S., Zawadzki, S. F., Leuteritz, A., & Wypych, F. (2017).

J. New alternative to produce colored polymer nanocomposites: organophilic Ni/Al and Co/Al layered double hydroxide as fillers into low-density polyethylene. Journal of the Brazilian Chemical Society. http://dx.doi.org/10.21577/0103-5053.20170093.

Received: Sept. 20, 2017 Revised: July 23, 2018 Accepted: Sept. 06, 2018

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Jaerger, S., Leuteritz, A., Freitas, R. A., & Wypych, F. Appendix A. Rheology of LDPE filled with Ni/Al-LDH or Co/Al-LDH intercalated with DDS or PAL and the shift factors (av) used in the master curves.

Figure A1. Master curves of the frequency dependence of G’ (S1, S3) and G” (S2, S4) of LDPE filled with Ni/Al-LDH, intercalated with ions of DDS (S1, S2) and PAL (S3, S4), obtained by the superposition of different filler concentrations, from 0 to 7wt%. The reference used was neat LDPE.

Figure A2. Vertical Shift factors (av) used due to Ni/Al-LDH fillers in LDPE. The avvalues were calculated based on the shear modulus G’ of LDPE at 190 °C. 12/13

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Rheological properties of low-density polyethylene filled with hydrophobic Co(Ni)-Al layered double hydroxides

Figure A3. Master curves of the frequency dependence of G’ (S1, S3) and G” (S2, S4) of LDPE filled with Co/Al-LDH, intercalated with ions of DDS (S1, S2) and PAL (S3, S4), obtained by the superposition of different filler concentrations, from 0 to 7wt%. The reference used was neat LDPE.

Figure A4. Vertical shift factors (av) used due to Co/Al-LDH fillers in LDPE. The avvalues were calculated based on LDPE at 190 °C.

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ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.03618

Alternative use of oily fractions of olive oil Melina Bagni1* , Dolly Granados1 and María Reboredo2 Instituto de Ingeniería Química – IIQ, Facultad de Ingeniería – FI, Universidad Nacional de San Juan – UNSJ, San Juan, Argentina 2 Departamento de Ingeniería de Materiales, Instituto de Investigaciones en Ciencia y Tecnología de Materiales – INTEMA, Consejo Nacional de Investigaciones Científicas y Técnicas – CONICET, Universidad Nacional de Mar del Plata – UNMdP, Mar del Plata, Argentina 1

*mbagni@unsj.edu.ar

Abstract Oily fractions of olive oil not suitable for human consumption used were: lamp oil (AL), olive pomace oil (OO) and clear oil lees (CB). Polyols were obtained by modifications of epoxidation with subsequent hydrolysis (ALH, OOH, CBH) and transesterification (ALHT, OOHT, CBHT) in order to favor, then, the polymerization reactions. The analysis of the physicochemical properties determined show the decrease of the unsaturations in the triglyceride and the increase of the OH concentration in the modified polyols as compared to the initial oils (from 16 to 380 for the AL and the ALHT; from 24 to 448 for the CB and the CBHT and from 3 to 430 for the OO and the OOHT, respectively) .The main objective of this work is to evaluate the stability of these oils over time and to provide an alternative synthesis of polyurethanes from a renewable resource, not previously used for this purpose.

Keywords: biopolymers, oil stability, polyols, polyurethanes, synthesis.

How to cite: Bagni, M., Granados, D., & Reboredo, M. (2019). Alternative use of oily fractions of olive oil. Polímeros: Ciência e Tecnologia, 29(1), e2019008. https://doi.org/10.1590/0104-1428.03618

1. Introduction Polymeric materials from renewable resources have attracted a lot of attention in recent years. The development and utilization of vegetable oils for polymeric materials are currently in the spotlight of the polymer and chemical industry[1]. Vegetable oils have a relatively low cost, are abundant and widely available. They are triglycerides obtained by the esterification of glycerol and fatty acids (predominantly unsaturated fatty acids). These triglycerides have several active sites that can be the starting point for polymerization reactions giving rise to products of the polymer industry. Although they are heterogeneous in composition, the final properties of the polyurethanes will depend to a large extent on the content of the hydroxyls of the polyol (derived from triglycerides) and not of their composition[2].

The objective of this work is to develop green polyols from oily fractions derived from the olive industry not suitable for human consumption without previous refinement that can compete successfully with those polyols of petrochemical origin, in the area of polyurethanes. In addition, the stability of the oils during storage was studied in order to determine the potential existence of physicochemical changes that could later modify the polymerization reactions. When the oils are stored they can undergo two types of chemical modifications. One of them, the enzymatic rancidity (or hydrolytic rancidity), is produced by enzymes (lipases) found in both the embryo and the mesocarp of the fruit, responsible for converting triglycerides to fatty acids and glycerol, which results in an increase in the acidity of stored

Polímeros, 29(1), e2019008, 2019

oils. These fatty acids show an increase until reaching a maximum from which they begin to descend. In turn, oxidative rancidity can occur, when the double bonds of the unsaturated fatty acids are oxidized in the presence of oxygen to peroxides or hydroperoxides that give rise to aldehydes, ketones, etc. Due to this, it is important to evaluate the oil physicochemical changes to predict the shelf life and in which stage of its storage it is convenient to use them as precursors of the polyols[3,4].

The main contribution of this work is to provide an alternative synthesis of polymers from a renewable resource and propose a new option for the final disposal of these oily fractions, since no works have been found on polyurethanes or polyols from these wastes.

2. Materials and Methods 2.1 Materials The oily raw materials used in this work were: Clear oil lees (CB), the oily stream extracted from the bottom of the decanters accompanied by sludges and vegetation water and that has begun to suffer oxidative processes; Lamp oil (AL), oily phase obtained from damaged or overripe olives and finally, Olive pomace oil (OO), residual oil extracted with solvents from the solid by-product generated during the process of olive oil extraction which is known as olive wet husk. All of these oily fractions need to be refined if they pretend to be suitable for human consumption.

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

Bagni, M., Granados, D., & Reboredo, M. Lamp oil was obtained from Nucete Company, La Rioja, Argentina. Olive pomace oil was obtained from Olivsan, San Juan, Argentina, and Clear oil lees was obtained from 6 Marías Establishment, San Juan, Argentina. Hydrogen peroxide (30 wt%) from Anedra Laboratories and formic acid (85 wt%) from Sintorgan Laboratories were used in the hydroxylation reaction. Triethanolamine (TEA >99%) from Biopack Laboratories and lithium hydroxide anhydrous (100%) from Hach Company were used in the transesterification reaction.

In the preparation of the polyurethane polymers, the isocyanate used was a commercial polymeric MDI (4,4-diphenylmethane diisocyanate) Rubinate 1680 from Huntsman Polyurethanes, USA.

2.2 Stability of oils The evolution with time of oils was studied by means of Fourier Transform Infrared Spectroscopy (FTIR) and by the determination of the Acid Value and the Iodine Index. The FTIR spectra were obtained using a Nicolet 6700 spectrometer from Thermo Scientific in absorbance mode, operated at a resolution of 2 cm-1, in a complete working range of 4000-500 cm-1. Acid value measurement was performed according to the Free Fatty Acids Method, AOCS Official Method Ca 5a-40. Iodine Index was determined by the Wijs method American Oil Chemists’ Society, AOCS Official Method Cd 1-25. The spectra found by Fourier Transform Infrared Spectroscopy were performed on samples without storage time, with one year and three years of storage, while the acid value and the iodine index tests were performed on oils without storage time and on oils with 2, 2 and a half and 3 years of storage and on oils without storage time and on oils with 1, 2 and 3 years of storage, respectively.

2.3 Synthesis of polyols Epoxidation with subsequent hydrolysis (Hydroxylation): Hydrogen peroxide together with formic acid were added to a stirred reactor at a temperature of 40 °C. The initial ratio used was H2O2/HCOOH (weight) = 1/1.9. The oil was added dropwise with an initial ratio oil/H2O2 (weight) of 3.5/1 and the temperature was maintained between 40 and 50 °C for 3 hours. Then the final product was cooled to room temperature. The liquid separated clearly into two phases. The upper phase was recovered. During this stage, hydrogen peroxide reacts with formic acid to form performic acid. Once the oil is added, it goes through an intermediate stage of epoxidation. The epoxy ring is unstable under acidic conditions so it opens up giving rise to the formation of hydroxyl groups[5].

Transesterification of the hydroxylated oil: The hydroxylated oil, as well as that obtained in the epoxidation stage, (ALH, OOH and CBH), triethanolamine and lithium hydroxide were placed together in a reactor with mechanical agitation.

The temperature was raised to 150 °C in 30 minutes and then maintained at this value for 2.5 hours. This is a complex stage consisting in the transesterification of the triglyceride molecules with triethanolamine to give rise to a mixture of different polyalcohols containing one or two chains of hydroxylated fatty acids (ALHT, OOHT, CBHT). No further purification steps were performed on the obtained polyol.

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The physicochemical characterization of the oils and their modifications were carried out through measurements of viscosity, density, iodine value, hydroxyl value, water content and volatile material, acid value, saponification index, Fourier Transform Infrared Spectroscopy (FTIR) and Nuclear Magnetic Resonance (NMR). The FTIR spectra were obtained under the same conditions used to determine the stability of the oils.

The NMR measurements were made using a spectrometer brand BRUKER model AVANCE II. It operates at a frequency of 300.13MHz for protons. A probe of CP/MAS was used with rotors for HRMAS of 4mm in diameter and volume of 57 µL. The protons spectra were recorded using the technique of direct polarization and rotation to the magic angle. The pulse of proton excitation (π/2) was 3.8µs of duration, the repetition time was 3s and the speed of rotation of the sample 3kHz. There were 32 averages per spectrum. The samples were tested as received without any extra treatment.

2.4 Synthesis of polyurethane Synthesis tests were carried out by reacting the obtained polyols (ALHT, CBHT, OOHT) with diisocyanate. The molar ratio used was NCO/OH = 1.1. The system was manually stirred for 30 seconds and the temperature was recorded during the test.

3. Results and Discussions 3.1 Stability of oils Figure 1 shows the infrared spectra of the CB oils according to the passage of time (only one of the fractions is shown since the spectra for the three oils are similar).

The main chemical change suffered by the oils during storage was the increase of the acid groups as a consequence of the enzymatic hydrolysis, responsible for breaking down triglycerides into free fatty acids and glycerol (enzymatic rancidity). In the spectra of CB and OO the appearance of a small peak at 1710 cm-1 (acidic groups) adjacent to the ester peak (1750 cm-1) in the sample stored for 3 years was detected, which in oils without storage was imperceptible.

In the AL sample it was observed that the peak corresponding to the acid group conserved the same size. Table 1 shows the results of acid value for the three samples analyzed in different periods of time studied. Total correspondence with the information provided by the FTIR spectra can be observed.

On the other hand, it could be affirmed that the samples have suffered oxidative rancidity due to the presence of oxygen during storage. Oxygen is responsible for oxidizing the double bonds present in the fatty acids, giving rise to the formation of peroxides or hydroperoxides. Table 2 shows the decrease in iodine values during storage, which can directly translate into a decrease in unsaturation.

Therefore, it can be concluded that it is convenient to carry out the hydroxylation and transesterification reactions to obtain polyols to samples without storage time, in order to have the highest number of double bonds in the samples prone to hydroxylation.

Polímeros, 29(1), e2019008, 2019

Alternative use of oily fractions of olive oil 3.2 Characterization of oils and their modifications The main differences in the FTIR spectra between untreated oils versus the hydroxylated and transesterified samples (OO versus OOHT) can be seen in Figure 2

Figure 1. (a) FTIR spectra of clear oil lees at different periods of storage.CB, CB1 and CB3, without storage, with 1 year and 3 years of storage, respectively; (b) Peak at 1750 cm-1 corresponding to the ester group and a peak at 1710 cm-1 corresponding to acid groups in the clear oil lees.

(only one of the oily fractions is shown since the spectra are similar) and lie in the appearance of a broad peak to 3400 cm-1 corresponding to the OH groups incorporated in the treated samples at the expense of the decrease of the peak corresponding to the unsaturations present at 3010 cm-1[6].

Table 3 shows the increase in OH values experimentally found by analytical techniques[7]. These values agree with the information provided by the spectra of each sample.

Similar behavior has been observed in tung oil samples[5].

The stretching vibration at 1746 cm-1 is attributed to the ester carbonyl funcional group of the triglycerides. The existence of peaks at 965 and 735 cm-1 indicates that both cis and trans conformations are present in the fatty acids that form the triglyceride[6].

In turn, NMR spectra corroborate these chemical modifications (Figures 3, 4 and 5). In these, the presence of new resonances in the hydroxylated and transesterified samples can be observed. They appear in the range 5.09‑5.29 ppm and are assigned to protons of OH groups. The signal of the unsaturated groups is still present (at 5.40 ppm) but decreased by less than half in intensity with respect to the unmodified sample. The signal of the protons of the CH2 groups neighboring to the unsaturated groups is also diminished (also in the middle), which is consistent with the decrease of unsaturated groups[5]. In the AL sample at 8 ppm (Figure 3b), a broad signal appears which normally corresponds to protons of carboxylic acid groups and which is not present in the CB and OO samples, since the lamp oil has the highest acidity index (see Table 1). The two

Figure 2. FTIR spectra of olive pomace oil and hydroxylated and transesterified olive pomace oil.

Table 1. Comparison of the acidity index of oils.

Samples AL CB OO

without storage 9.7 ± 0.1 1.1 ± 0.2 2.0 ± 0.1

Acidity Index % 2 years of storage 10.7 ± 0.3 1.0 ± 0.1 1.6 ± 0.4

2 ½ years of storage 10.5 ± 0.1 5.2 ± 0.1 2.5 ± 0.1

3 years of storage 10.4 ± 0.3 5.6 ± 0.2 2.7 ± 0.2

Iodine Index (Wijs) 1 years of storage 83 ± 2 84 ± 1 81 ± 3

2 years of storage 77 ± 2 79 ± 3 75 ± 2

3 years of storage 76 ± 7 69 ± 8 74 ± 9

Table 2. Iodine index for different storage times.

Samples AL CB OO

without storage 80 ± 2 84 ± 1 83 ± 1

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Bagni, M., Granados, D., & Reboredo, M. Table 3. OH value for oils and their modifications.

Samples Value OH mgKOH/g

AL 16

ALHT 380

CB 24

CBHT 448

OO 3

OOHT 430

Figure 5. NMR spectra of clear oil lees and hydroxylated and transesterified clear oil lees.

Figure 6. Polyurethane obtained from OOHT (a), from CBHT (b) and from ALHT (c).

Figure 3. (a) NMR spectra of lamp oil and hydroxylated and transesterified lamp oil; (b) Resonance present at 8 ppm in AL, corresponding to acid groups.

iodine index due to the opening of fatty acid unsaturations to achieve their incorporation[7].

3.3 Synthesis of polyurethanes Preliminary tests of polyurethane synthesis were performed by the reaction of the obtained polyols (ALHT, CBHT, OOHT) with diisocyanate (RUBINATE 1680), as it can be seen in Figure 6.

Figure 4. NMR spectra of olive pomace oil and hydroxylated and transesterified olive pomace oil.

multiplets at 4.4 and 4.1 ppm corresponds to the four glycerol methylene protons in the triglyceride molecule, while the small multiplet at 5.3 ppm corresponds to the central H in the glycerol moieties.

The physicochemical characterization determined that after the modification reactions (epoxidation with subsequent hydrolysis and transesterification) an increase of hydroxyl groups (greater number of hydroxyls) was reached, which is reflected in an increase in viscosity and reduction of 4/5

The system was manually stirred for 30 seconds.

Opalescence, foaming and gelation times and the reaction temperature were recorded throughout the test (see Table 4).

The polymerization reaction between the OOHT polyol with the diisocyanate is the most exothermic and has the shortest opalescence, foaming and gelation times. The curing reaction occurs faster and its final appearance is porous, which could be attributed to the water content remaining in the oil. The final volume is more than twice the initial volume for the ALHT and OOHT samples while the CBHT only increases by 66%.

The reactions between the ALHT and CBHT polyols with diisocyanate present similar opalescence, foaming and gelation times. The main difference found between these samples is their final appearance. The sample obtained from ALHT is more compact, while the sample obtained from CBHT has a greater number of small pores. The samples are in the characterization stage, which includes tests of FTIR, DSC (differential scanning calorimetry), TGA (thermogravimetry), among others.

Polímeros, 29(1), e2019008, 2019

Alternative use of oily fractions of olive oil Table 4. Characteristics observed during the synthesis of polyurethane.

Samples

Initial Volume (mL)

Final Volume (mL)

Opalescence (s)

ALHT CBHT OOHT

12 12 12

25 20 25

10 20 5

4. Conclusions The study of the stability of the oils allowed to know in which stage of the storage it is convenient to chemically modify the oils to obtain the polyols. Therefore, it can be concluded that the oils must be treated without storage time since it is when they present a greater amount of unsaturations that can be hydroxylated and thus contribute to the increase of the content of OH groups in the polyols.

By means of epoxidation and transesterification reactions, it was possible to increase the content of OH groups in order to obtain polyols that give rise to thermoset polyurethanes.

The three oily fractions adopted are suitable for carrying out polymerization reactions with diisocyanates.

The three polymeric materials obtained present well differentiated aspects. Subsequent tests of physicochemical characterization will allow to associate the resulting properties of the polymers with the characteristics of the chemically modified oily fractions.

5. Acknowledgements The authors gratefully acknowledge the CONICET for the doctoral fellowship of M.M. Bagni and for the financial support, as well as Mr. Oscar Casemayor and Dr. Diana Fasce for the chemical and analytical measurements.

6. References 1. Samarth, N. B., & Mahanwar, P. A. (2015). Modified vegetable oil based additives as a future polymeric material. Open

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Temperature (°C) 4 min 48 48 58

Foaming time (s)

Gelling time (s)

112 120 90

180 180 120

Final Appearance (3 h) Rubbery Pasty Rigid

Journal of Organic Polymer Materials, 5(1), 1-22. http:// dx.doi.org/10.4236/ojopm.2015.51001.

2. Petrović, Z. S., Zhang, W., Zlatanić, A., Lava, C. C., & Ilavskyý, M. (2002). Effect of OH/NCO molar ratio on properties of soy-based polyurethane networks. Journal of Polymers and the Environment, 10(1-2), 5-12. http://dx.doi. org/10.1023/A:1021009821007.

3. Martinez C. (2018). Determination of the acidity index in oils and edible fats. Retrieved in 2018, May 28, from https:// es.scribd.com/doc/97574878/Determinacion-del-indice-deacidez-en-aceites-y-grasas-comestibles 4. Shahidi, F., & John, J. A. (2013). Oxidative rancidity in nuts.

In L. Harris (Ed.), Improving the safety and quality of nuts (pp.

198-229). Cambridge: Woodhead Publishing. http://dx.doi. org/10.1533/9780857097484.2.198 5. Vlachos, N., Skopelitis, Y., Psaroudaki, M., Konstantinidou, V., Chatzilazarou, A., & Tegou, E. (2006). Applications of Fourier transform-infrared spectroscopy to edible oils. Analytica Chimica Acta, 573-574, 459-465. http://dx.doi.org/10.1016/j. aca.2006.05.034. PMid:17723561.

6. Bagni, M., Reboredo, M., & Granados, D. (2015). Use of oily wastes to obtain polymers. In Latin American Congress of Engineering and Applied Sciences (p. 201). Mendoza: National University of Cuyo.

7. Mosiewicki, M., Casado, U., Marcovich, N., & Aranguren, M. (2009). Polyurethanes from tung oil: polymer characterization and composites. Polymer Engineering and Science, 49(4), 685-692. http://dx.doi.org/10.1002/pen.21300.

Received: May 29, 2018 Revised: July 11, 2018 Accepted: July 17, 2018

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ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.01018

Sulfonated poly(ether ether ketone)/hydroxyapatite membrane as biomaterials: process evaluation Cristiane Agra Pimentel1* , José William de Lima Souza1, Flávia Suzany Ferreira dos Santos1, Mayelli Dantas de Sá1, Valéria Pereira Ferreira1, Gislaine Bezerra de Carvalho Barreto1, José Filipe Bacalhau Rodrigues1, Wladymyr Jefferson Bacalhau de Sousa1, Cláudio Orestes Britto Filho1, Francisco Kegenaldo Alves de Sousa1 and Marcus Vinicius Lia Fook1 Laboratório de Avaliação e Desenvolvimento de Biomateriais do Nordeste – CERTBIO, Unidade Acadêmica de Engenharia de Materiais – UAEMa, Centro de Ciências e Tecnologia – CCT, Universidade Federal de Campina Grande – UFCG, Campina Grande, PB, Brasil

1

*cristiane.pimentel@certbio.ufcg.edu.br

Abstract Poly(ether ether ketone) (PEEK) has excellent properties, such as high biocompatibility and an elastic modulus similar to bone, which makes it a suitable biomaterial. When modified with sulfuric acid (H2SO4) and hydroxyapatite (HA), its workability and bioactivity is enhanced, and this makes it of great importance in medicine. This study investigates a better combination of process parameters to manufacture sulfonated PEEK/HA (SPEEK/HA) membranes for biomaterials.

Chemical, thermal, and biological analyses were carried out on all samples. The sulfonated structure was observed to enhance wettability, adhesion, and cell viability. Furthermore, an increase in the degree of sulfonation facilitated their workability as required for biomaterials; making them suitable for osseointegration. Besides, the SPEEK/HA membranes presented cell adhesion, confirming the viability to use as biomaterial. This study presents a cheap alternative method to easily process biomaterials of improved workability.

Keywords: biomaterial, chemical modification, hydroxyapatite, membrane, SPEEK. How to cite: Pimentel, C. A., Souza, J. W. L., Santos, F. S. F., Sá, M. D., Ferreira, V. P., Barreto, G. B. C., Rodrigues, J. F. B., Sousa, W. J. B., Britto Filho, C. O., Sousa, F. K. A., & Fook, M. V. L. (2019). Sulfonated poly(ether ether ketone)/ hydroxyapatite membrane as biomaterials: process evaluation. Polímeros: Ciência e Tecnologia, 29(1), e2019009. https://doi.org/10.1590/0104-1428.01018

1. Introduction Poly(ether ether ketone) (PEEK) has attracted interest in medicine. This polymer has been used in orthopedics, neurosurgery, and traumatology because of its favorable mechanical, chemical, and tribological properties[1]. The benefits of this material include excellent mechanical resistance, high biocompatibility, improved biological inertia, low friction coefficient, elastic modulus similar to bone, reusability, and resterilization[2,3]. However, its high processing temperature makes it difficult to be worked into shape.

Recently, research has been done to chemically modify PEEK with sulfuric acid (H2SO4) to improve its workability and processing. Studies were initially geared towards its use in fuel cells, which convert chemical energy to electrical energy[4]. However, more recently, it has been studied to be used as b​​ iomaterials. Zhao et al.[5] studied sulfonated PEEK (SPEEK) applications in orthopedic implants. Kalambettu and Dharmalingam[6] studied the fabrication of SPEEK membranes incorporated with hydroxyapatite (HA), and Montero et al.[7] researched biofilms fabricated with SPEEK. Although SPEEK is commonly applied in fuel cells and biomaterials, further studies for other possible applications are necessary[5], as it possesses some limitations. Some of

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these are its inability for direct bone apposition, which could result in poor osseointegration[8]; and a possible presence of acid residue, which makes it risky for medical applications[6].

SPEEK is a relevant biomaterial ​​ in terms of medicine devices and scientific prospects because of its simple processing and requirement of ordinary equipment. It is also easy to be worked into shape and incorporated with drugs[7].

This paper focuses on a detailed evaluation of the influence of process parameters on the performance of a SPEEK/HA membrane. The chemical, thermal, and wetting properties of SPEEK/HA were evaluated. In addition, cell viability and adhesion were studied to be applied in biomaterials.

2. Materials and Methods 2.1 SPEEK/HA membranes preparation Two grams of PEEK (Victrex USA Inc – Vicote 702) were dissolved in 50 mL of H2SO4 98% (VETEC Química Fina Ltda). The mixture was heated up to 50 °C and mechanically stirred for 3 h. After 1.5 h of stirring, 0.6 g of HA (Labsynth) was added (during stirring). After 4 days, the solution was poured on a plate and frozen at -80 °C.

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

Pimentel, C. A., Souza, J. W. L., Santos, F. S. F., Sá, M. D., Ferreira, V. P., Barreto, G. B. C., Rodrigues, J. F. B., Sousa, W. J. B., Britto Filho, C. O., Sousa, F. K. A., & Fook, M. V. L. It was placed into a freeze drier for 24 h and allowed to stand for another 24 h before it was decanted to isolate the membrane. The product was then filtered, washed with distilled water until complete removal of H2SO4, and dried at 50 °C in an oven for 3 h.

2.2 Fourier Transform Infrared Spectroscopy (FTIR) FTIR analysis was carried out in a Perkin Elmer spectrophotometer (Spectrum 400), and the spectra were recorded in absorbance mode, to detect any chemical bonding at the SPEEK/HA membranes and to calculate the degree of sulfonation (DS). The DS of PEEK was calculated using the following Equation 1: %DS

 SPEEK peak height  1 − PEEK peak height  × 100 (1)    

where the normalized SPEEK peak height (1484 cm-1) is compared with the PEEK peak height (1494 cm-1) to obtain the %DS. The PEEK sulfonation can be confirmed by the division of the aromatic C=C absorption band at 1494 cm-1 and the appearance of a new band at 1484 cm-1. Therefore, it is necessary to measure the heights of the characteristic bands of both materials[4].

2.3 Thermogravimetric Analysis (TGA) The thermal properties of SPEEK/HA membranes were examined using TGA. Thermal analysis was carried out in a nitrogen atmosphere between 25 °C (room temperature) and 600 °C at a heating rate and flow rate of 10 °C/min and 20 mL/min, respectively, using a TG 50H (Shimadzu) analyzer.

2.4 Wettability analysis The wettability analysis was done by the static drop method. A contact angle goniometer (CERTBIO) was used to measure the contact angle of the membrane. Deionized water was used as the liquid. A total of three points of the contact angle was measured in each sample within 30 s after the drop.

2.5 Cytotoxicity study The cytotoxicity analysis was performed according to ISO 10993-5: 2009 (Biological evaluation of medical devices -- Part 5: Tests for in vitro cytotoxicity)[9]. An L929 fibroblast cell line (ATCC NCTC clone 929) was grown in RPMI culture medium (RPMI 1640 Medium, Gibco - Invitrogen Corporation, Grad Island, USA) and was supplemented with 10% bovine fetal serum (Gibco, Life Technologies) and 1% antibiotic-antimycotic (Gibco, Life Technologies).

These cells were preserved in CO2 incubators at 37 °C in 5% atmosphere. The cell suspension (100 μL per well) was added to a 96-well plate at 1 × 105 cells/mL in the RPMI 1640 culture medium. The plate was transferred to a CO2 oven (5%) at 37 °C and incubated for 24 h. The culture medium was aspirated from all wells, and then 170 μL RPMI 1640 culture medium and 30 μL sample extract were added to each well. The plate was incubated again in a CO2 oven (5%) at 37 °C for 24 h. The culture medium was aspirated 2/8

from all wells and 100 μL of MTT solution (1 mg/mL) was added. The plates were incubated again for 3 h in a CO2 oven (5%) at 37 °C. The supernatant was discarded and 100 μL of isopropyl alcohol was added per well. Optical density was read on a microplate reader (VictorX3 - PerkinElmer) at 570 nm with 650 nm reference filters. Cell viability was calculated as a percentage of the modified z-Score test for outliers detection. High density polyethylene (HDPE) and natural latex were used as a negative and positive control, respectively.

2.6 Cell adhesion To evaluate the cell adhesion of the membrane, the samples were preserved in 70% ethanol for 24 h. They were then washed thrice in sterile PBS and dried at 40 °C for 24 h.

They were placed into a 48-well tissue culture plate, and 500 μL of OFCOL II cell suspension was added per well at 1 × 105 cells/mL in RPMI 1640 culture medium (RPMI 1640 Medium, Gibco - Invitrogen Corporation, Grad Island, USA) supplemented with 10% bovine fetal serum (Gibco, Life Technologies) and 1% antibiotic-antimycotic (Gibco, Life Technologies). The plate was transferred to the CO2 incubator (5%) at 37 °C and incubated for 7 days. After incubation, the culture medium was aspirated from all the wells and the samples were washed with PBS. The PBS was aspirated, and formaldehyde solution was added to each well at 10% for 10 min for cell attachment. The formaldehyde was removed and the samples were washed with PBS. The PBS was removed and the samples were dried at 40 °C for 24 h. Cell adhesion was evaluated through the surfaces of the samples by scanning electron microscopy (SEM) (WORLD PHENOM - PRO-X 800‑07334 model); the samples for which were coated with gold for better visualization.

2.7 Design of Experiment (DOE) After defining and validating the methodology of other studies, this research chose DOE to study the influence of process parameters on SPEEK/HA membrane properties.

The studied parameters (inputs) were PEEK sulfonation time, HA addition time, and freeze-drying time. Duplicates of a 23 factorial design without center points were conducted, and Minitab 18 was used to analyze the results.

Table 1 shows the experimental planning matrix with minimum and maximum values for ​​ each variable.

Table 2 shows the samples and input parameters, including the duplicates (E9 to E16). All experiments were performed randomly.

The output parameters to evaluate the experimental design were based on characterizations that could greatly influence the membrane properties of biomaterials. The characterizations are FTIR analysis, defined by %DS; TGA, defined by the Table 1. Experimental planning matrix.

Input Parameters Freeze-drying time Sulfonation time HA addition time

Unit H H H

-1 24 2.0 0.5

+1 48 3.0 1.5

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Sulfonated poly(ether ether ketone)/hydroxyapatite membrane as biomaterials: process evaluation Table 2. Sample identification per experiment.

Experiment 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16

Sample E1 E2 E3 E4 E5 E6 E7 E8 E9 E10 E11 E12 E13 E14 E15 E16

Freeze-drying time 24 48 24 48 24 48 24 48 24 48 24 48 24 48 24 48

weight loss during a change from 300 to 400 °C, which is due to the removal of sulfonic acid groups[7]; and wettability analysis, defined by the water contact angle. A higher %DS increases the wettability, which would facilitate workability of the membrane and osseointegration[10-12]. Table 3 presents the output parameters per characterization at experimental planning.

3. Results and Discussions 3.1 FTIR analysis FTIR spectra of the membranes are shown in Figure 1 and 2.

PEEK was used as the target material and reference for identifying functional groups.

The sulfonation of PEEK is confirmed by the division of the aromatic C=C absorption band at 1494 cm-1 and the appearance of a new band at 1484 cm-1. This peak was used to calculate %DS during DOE (Figure 1a). A sharp peak at 1024 cm−1 confirms the presence of a H2SO4 group, which is due to the S=O group in all samples and duplicates (Figure 1a). A peak at 3425 cm−1 confirms the presence of the OH group, which is bonded to an SO3H group. The peak intensities in all samples are different because of varied process parameters. The highest peak intensity corresponded to a higher degree of sulfonation, as observed in E8, E12, and E15 (Figure 1 and 2). Previous studies have reported similar bands[7,13,14].

A shift of the peak from 3447 cm−1 to 3425 cm−1 was observed for samples where HA was deposited on the SPEEK surface (Figure 1). The coordinate bond of the HA polar group and OH group of the SPEEK was weak as a result of the stretching of OH bonds[15], and this led to a low peak frequency.

3.2 TGA TGA curves are shown in Figure 3 and 4. PEEK was used as the target material. SPEEK/HA membranes of all samples were thermally stable up to 350 °C and exhibited two distinct weight loss stages, depending on the process Polímeros, 29(1), e2019009, 2019

Sulfonation time 2 2 3 3 2 2 3 3 2 2 3 3 2 2 3 3

HA addition time 0.5 0.5 0.5 0.5 1.5 1.5 1.5 1.5 0.5 0.5 0.5 0.5 1.5 1.5 1.5 1.5

Table 3. Output parameter per characterization.

Characterization 1 - FTIR 2 - Wettability 3 - TGA

Unit % Degree %

Output parameter Degree of sulfonation (DS) Water contact angle Weight loss - ranged from 300 to 400 °C (removal of sulfonic acid groups)

parameters used, as shown in Figure 3 and 4, whereas PEEK samples had one weight loss at 550 °C (Figure 3 and 4).

The mass loss at 100 °C is due to the removal of water molecules absorbed by the material. The first stage of thermal degradation of samples during a shift from 300 to 400 °C is due to the removal of SO3H. The second stage of thermal degradation ranging from 500 to 600 °C is due to polymer degradation. All these results are corroborated by literature[4,7].

In the final degradation stage around 550 °C, the samples approach a final mass at 600 °C, except samples E5 and E13, as seen in Figures 3 and 4; however this does not happen in the PEEK sample. Zaidi et al.[16] explained that when H2SO4 (95-98%) is used for sulfonation, material degradation and crosslinking reactions are avoided.

TGA was used to estimate the weight loss of the samples on DOE by assuming that the first stage of thermal degradation is entirely caused by elimination of sulfonic acid groups.

3.3 DOE results Table 4 presents the DOE results, outlining all samples and parameters. The mass loss of samples E5 and E13 could not be obtained probably because of reaction problems, such as problems at sulfonation reaction.

The main objective of DOE was to obtain the best combination of output parameters for optimum performance of the samples as biomaterials. The best combination was a high degree of sulfonation, high wettability, and higher mass loss at 350 °C. An increase in %DS facilitated the wettability of the materials, improving their workability as required for biomaterials; making them suitable for osseointegration[17].

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Pimentel, C. A., Souza, J. W. L., Santos, F. S. F., Sá, M. D., Ferreira, V. P., Barreto, G. B. C., Rodrigues, J. F. B., Sousa, W. J. B., Britto Filho, C. O., Sousa, F. K. A., & Fook, M. V. L.

Figure 1. FTIR spectra of (a) PEEK, E1 to E4 samples; (b) PEEK, E9 to E12 duplicates.

Figure 2. FTIR spectra of (a) PEEK, E5 to E8 samples; (b) PEEK, E13 to E16 duplicates.

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Sulfonated poly(ether ether ketone)/hydroxyapatite membrane as biomaterials: process evaluation

Figure 3. TGA curves of (a) PEEK, E1 to E4 samples; (b) PEEK, E9 to E12 duplicates.

Figure 4. TGA Curves of (a) PEEK, E5 to E8 samples; (b) PEEK, E13 to E16 duplicates.

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Pimentel, C. A., Souza, J. W. L., Santos, F. S. F., Sá, M. D., Ferreira, V. P., Barreto, G. B. C., Rodrigues, J. F. B., Sousa, W. J. B., Britto Filho, C. O., Sousa, F. K. A., & Fook, M. V. L. Some samples had their best combination impaired, such as E1 and E14, all of them presented a higher %DS and a contact angle inconsistent, it was observed probably by acid residues. In future studies, it is important to detail the influence of acid residues.

Table 5 shows the Minitab optimization analysis.

Considering the adjustments and confidence intervals, the best combination of input parameters to obtain the desired outputs were a freeze-drying time of 24 h, sulfonation time of 3 h, and HA addition time of 1.5 h. Samples E7 and E15 had these combinations. Almasi et al.[13], Kalambettu and Dharmalingam[6], and Zhou and Lee[18] conducted a similar study with another combination. This is a novel research.

In this way, the average cell viability shown satisfactory and contributes to their application as a biomaterial. But when it was studied these applications, it is important to detail the influence of acid residues.

3.5 Cell adhesion Cell adhesion analysis was performed through SEM with a gold-metallized surface and results are shown in Figure 6. Figure 6a shows the structure of a SPEEK/HA membrane without cell adhesion. Cell adhesion can be observed in Figure 6b, where the SPEEK/HA membranes show a layer of coating, which is in agreement with the literature[5,6]. This confirms the possibility of their use as a biomaterial.

3.4 Cell viability Considering the optimum result of E7/E15, the cell viability of the membrane was analyzed for use in biomaterials.

Figure 5 shows the membrane cytotoxicity according to BS EN ISO 10993-5:2009—Tests for in vitro cytotoxicity[9]. Cell viability is specified when calculated values are a​​ bove 70%.

Cytocompatibility results of optimum samples toward the fibroblast-like L929 cells show an average cell viability of 86%, which is not much higher than the specification (70%).

This percentage is probably due to extensive agglomeration of the HA filler particles in SPEEK/HA membrane. Cells are known to be very sensitive to surface energy and chemistry[13].

Besides, it should be considered the adverse conditions to sulfonated PEEK at laboratory as well as the acid residues.

Figure 5. Cell viability of L929 cell lines by MTT method.

Table 4. DOE results correlating samples and output results.

Samples E1 E2 E3 E4 E5 E6 E7 E8 E9 E10 E11 E12 E13 E14 E15 E16

%DS 83.96 54.04 62.88 67.93 49.96 61.60 70.13 66.12 68.64 48.70 48.40 47.83 41.81 75.21 61.78 54.20

Contact angle (degrees) 32.25 28.53 29.51 23.92 28.41 31.97 25.14 24.71 22.23 27.64 18.00 26.05 32.19 24.55 24.03 31.23

% Mass loss at 350 °C 11.51 16.86 7.77 10.30 ----12.64 12.95 10.71 11.62 15.68 8.21 7.05 ----11.17 10.93 11.50

Table 5. Minitab optimization analysis.

1 Output

Freeze-drying time -1 Adjustment

%DS Contact angle %Mass loss

70.66 24.63 10.94

Solution

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Sulfonation time 1 Standard adjustment error 8.03 2.37 3.04

HA addition time 1 Confidence bounds 95% (52.15; 89.18) (19.17; 30.09) (3.93; 17.95)

%DS

Contact angle

% Mass loss

70.66

24.63

10.94

Compound desirability 0.62

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Sulfonated poly(ether ether ketone)/hydroxyapatite membrane as biomaterials: process evaluation

Figure 6. SEM micrographs of membranes (sample E7) (a) SPEEK/HA before cell adhesion; (b) SPEEK/HA after cell adhesion and gold metallization.

4. Conclusions This study reported a process evaluation of SPEEK/HA membrane application as biomaterials. The chemical deposition technique used did not involve high temperatures, and the sulfonated structure improved wettability, cell adhesion, and growth. This technique of surface treatment is relatively cheap and easy to perform. Hence, this method can be adopted to modify PEEK membranes for better application as biomaterials.

5. Acknowledgements This work was supported by Coordination for the Improvement of Higher Level Personnel – CAPES.

6. References 1. Wang, M., Tang, S. J., McGrady, L. M., & Rao, R. D. (2013).

Biomechanical comparison of supplemental posterior fixations for two-level anterior lumbar interbody fusion. Journal of Engineering in Medicine, 227(3), 245-250. http://dx.doi. org/10.1177/0954411912465057. PMid:23662340.

2. O’Reilly, E. B., Barnett, S., Madden, C., Welch, B., Mickey, B., & Rozen, S. (2015). Computed-tomography modeled polyetheretherketone (PEEK) implants in revision cranioplasty.

Journal of Plastic, Reconstructive & Aesthetic Surgery, 68(3), 329-338. http://dx.doi.org/10.1016/j.bjps.2014.11.001.

PMid:25541423.

3. Siddiq, A. R., & Kennedy, A. R. (2015). Porous polyetheretherketone (PEEK) manufactured by a novel powder route using nearspherical salt bead porogens: characterisation and mechanical properties. Materials Science and Engineering C, 47(1), 180-188. http://dx.doi.org/10.1016/j.msec.2014.11.044. PMid:25492187.

4. Aguiar, K. R., Batalha, G. P., Peixoto, M., Ramos, A., & Pezzin, S. H. (2012). Produção de membranas híbridas zirconizadas de SPEEK/Copolissilsesquioxano para aplicação em células a combustível do tipo PEM. Polímeros: Ciência e Polímeros, 29(1), e2019009, 2019

Tecnologia, 22(5), 453-459. http://dx.doi.org/10.1590/S010414282012005000060.

5. Zhao, Y., Wong, H. M., Wang, W., Li, P., Xu, Z., Chong, E. Y., Yan, C. H., Yeung, K. W., & Chu, P. K. (2013). Cytocompatibility, osseointegration, and bioactivity of three-dimensional porous and nanostructured network on polyetheretherketone.

Biomaterials, 34(37), 9264-9277. http://dx.doi.org/10.1016/j. biomaterials.2013.08.071. PMid:24041423.

6. Kalambettu, A., & Dharmalingam, S. (2014). Fabrication and in vitro evaluation of sulphonated polyetheretherketone/ nano-hydroxyapatite composites as bone graft materials.

Materials Chemistry and Physics, 147(1-2), 168-177. http:// dx.doi.org/10.1016/j.matchemphys.2014.04.024.

7. Montero, J. F., Tajiri, H. A., Barra, G. M., Fredel, M. C., Benfatti, C. A., Magini, R. S., Pimenta, A. L., & Souza, J. C. (2017).

Biofilm behavior on sulfonated polyetheretherketone (sPEEK).

Materials Science and Engineering C, 70(1), 456-460. http:// dx.doi.org/10.1016/j.msec.2016.09.017. PMid:27770916.

8. Kurtz, S. M., & Devine, J. N. (2007). Applications of polyetheretherketone in trauma, arthroscopy, and cranial defect repair. In S. Lovald & S. M. Kurtz (Eds.), PEEK biomaterials handbook, (pp. 243-260). New York: Springer .

9. International Organization for Standardization – ISO. (2009).

BS EN ISO 10993-5: biological evaluation of medical devices: tests for in vitro cytotoxicity. Genebra: ISO. 10. Conceição, T. F., Bertolino, J. R., Barra, G. M., Mireski, S. L., Joussef, A. C., & Pires, A. T. (2008). Preparation and characterization of polyetheretherketone derivatives. Journal of the Brazilian Chemical Society, 19(1), 111-116. http://dx.doi. org/10.1590/S0103-50532008000100016.

11. Jiang, R., Kunz, H. R., & Fenton, J. M. (2005). Investigation of membrane property and fuel cell behavior with sulfonated polyetheretherketone electrolyte: temperature and relative humidity effects. Journal of Power Sources, 150(4), 120-128. http://dx.doi.org/10.1016/j.jpowsour.2005.03.180.

12. Xing, P., Robertson, G. P., Guiver, M. D., Mikhailenko, S. D., Wang, K., & Kaliaguine, S. (2004). Synthesis and characterization of sulfonated polyetheretherketone for proton 7/8

Pimentel, C. A., Souza, J. W. L., Santos, F. S. F., SĂĄ, M. D., Ferreira, V. P., Barreto, G. B. C., Rodrigues, J. F. B., Sousa, W. J. B., Britto Filho, C. O., Sousa, F. K. A., & Fook, M. V. L. exchange membranes. Journal of Membrane Science, 229(1-2), 95-106. http://dx.doi.org/10.1016/j.memsci.2003.09.019.

13. Almasi, D., Izman, S., Assadian, M., Ghanbari, M., & Abdul Kadir, M. R. (2014). Crystalline HA coating on PEEK via chemical deposition. Applied Surface Science, 314(30), 10341040. http://dx.doi.org/10.1016/j.apsusc.2014.06.074.

14. Jaafar, J., Ismail, A., & Mustafa, A. (2007). Physicochemical study of polyetheretherketone electrolyte membranes sulfonated with mixtures of fuming sulfuric acid and sulfuric acid for direct methanol fuel cell application. Materials Science and Engineering A, 460(15), 475-484. http://dx.doi.org/10.1016/j. msea.2007.02.095.

15. Janaki, K., Elamathi, S., & Sangeetha, D. (2005). Development and characterization of polymer ceramic composites for orthopedic applications. Artificial Organs, 22(3), 169-178.

Retrieved in 2018, February 12, from http://medind.nic.in/ taa/t09/i3/taat09i3p169.pdf

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16. Zaidi, S. J., Mikhailenko, S. D., Robertson, G., Guiver, M., & Kaliaguine, S. (2000). Proton conducting composite membranes from polyetheretherketone and heteropolyacids for fuel cell applications. Journal of Membrane Science, 173(1), 17-34. http://dx.doi.org/10.1016/S0376-7388(00)00345-8.

17. Mahjoubi, H., Buck, E., Manimunda, P., Farivar, R., Chromik, R., Murshed, M., & Cerruti, M. (2017). Surface phosphonation enhances hydroxyapatite coating adhesion on polyetheretherketone and its osseointegration potential. Acta Biomaterialia, 47(1), 149-158. http://dx.doi.org/10.1016/j.actbio.2016.10.004.

PMid:27717913.

18. Zhou, H., & Lee, J. (2011). Nanoscale hydroxyapatite particles for bone tissue engineering. Acta Biomaterialia, 7(7), 2769-2781. http://dx.doi.org/10.1016/j.actbio.2011.03.019. PMid:21440094.

Received: Feb. 12, 2018 Revised: July 19, 2018 Accepted: Sept. 12, 2018

PolĂ­meros, 29(1), e2019009, 2019

ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.00218

Predicting LDPE/HDPE blend composition by CARS-PLS regression and confocal Raman spectroscopy Daniel José da Silva1*  and Hélio Wiebeck1 Departamento de Engenharia Metalúrgica e de Materiais, Escola Politécnica, Universidade de São Paulo – USP, São Paulo, SP, Brasil

1

*dankuruta@hotmail.com

Abstract Industries and the scientific community currently focus on creating new ways to recycle and to reuse polymer waste that leads to serious socio-environmental risks. However, the quality of recycled polyethylenes depends strongly on their purity degree, but the distinction between Low Density Polyethylene (LDPE) and High Density Polyethylene (HDPE) by a fast and consistently good methodology is still an unsolved issue for the current recycling processes.

In this study, confocal Raman spectroscopy and Competitive Adaptive Reweighted Sampling - Partial Least Squares (CARS-PLS) linear regression have been successfully applied to quantify the concentration of LDPE/HDPE blends.

The effects of several regression parameters (pretreatment method, Monte Carlo sampling runs, k-fold and maximal number of latent variables for cross-validation) on the CARS-PLS model training and prediction performance were analyzed. The CARS-PLS-based models show root-mean-squared prediction error of 4.06 - 8.87 wt% of LDPE for the whole composition range of HDPE/LDPE blend.

Keywords: CARS-PLS regression, polymer blends, polyethylene, confocal Raman spectroscopy.

How to cite: Silva, D. J., & Wiebeck, H. (2019). Predicting LDPE/HDPE blend composition by CARS-PLS regression and confocal Raman spectroscopy. Polímeros: Ciência e Tecnologia, 29(1), e2019010. https://doi.org/10.1590/01041428.00218

1. Introduction Polyethylenes (PEs) are the main thermoplastic polymers consumed by the current civilization and, consequently, the largest polymer fraction found in urban solid wastes.

The reason for the great versatility of their mechanical properties is the control of the degree of polymeric branches during the ethene polymerization by a low-cost production[1].

However, this characteristic of PEs results in several difficulties in manufacturing their recycled products with attractive properties by mechanical recycling[2,3].

HDPE/LDPE blends have been widely used by the plastic industry to adjust processability and mechanical properties of the polyethylene resins[4]. However, the unknown and uncontrolled composition of these polymeric blends and recycled polyethylene wastes hinders the processing and production of material goods with satisfactory performance and quality.

In several countries, Low Density Polyethylene (LDPE) and High Density Polyethylene (HDPE) are the main representatives in the family of PEs due to their higher degree of production than that observed for other polyethylenes commercially available[5,6], such as Linear Low Density Polyethylene (LLDPE), Ultra High Molecular Weight Polyethylene (UHMWPE) and Ultra Low Density Polyethylene (ULDPE).

LDPE and HDPE are semi-crystalline thermoplastics, frequently distinguished by their densities

Polímeros, 29(1), e2019010, 2019

(δLDPE = 0.91-0.93 g/cm3 and δHDPE = 0.95-0.97 g/cm3)[7], which are closely linked to their differences in the number of polymer branches[8-10]. LDPE is commonly processed by extrusion, blow molding and injection molding. This polyethylene has high impact resistance and flexibility among the PEs, as well as interesting electrical properties to be used as an electrical insulator. Consequently, LDPE has been applied to the production of flexible packaging, wiring and cable coating. HDPE is used in several segments: buckets, bowls, trays, toys and pots are obtained by injection processing; packaging for detergents and cosmetics are made by blowing processing; insulation of telephone wires, decorative tapes, garbage bags and grocery bags are obtained by extrusion[7].

The determination of the fractional composition of LDPE/HDPE blends is not a simple task because the chemical structures of their polymer chains are only based on carbon and hydrogen. Wide-Angle X-Ray Scattering (WAXS), Dynamic Mechanical Thermal Analysis (DMTA) and Differential Scanning Calorimetry (DSC) have shown limitations to estimate the composition of this polymer blend due to the effects of co-crystallization for blends with more than 10 wt% of LDPE[11]. Contrary to these characterization techniques, confocal Raman spectroscopy is a quick, nondestructive and inexpensive method since it does not require expensive inputs or time-consuming methods for sample preparation and analysis[12].

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

Silva, D. J., & Wiebeck, H. In this contribution, we evaluated the potential of Partial Least Squares linear regression modified by Competitive Adaptive Reweighted Sampling (CARS-PLS) to analytically determine the compositional fraction of LDPE/HDPE blends by prediction models based on confocal Raman data.

1.1 Mathematical and computational fundamentals PLS linear regression is an mathematical method that externally correlates an instrumental data set (X matrix) and an interest property set (Y matrix) by linear equations[13,14]: = X ( n×k ) T( n×h ) * P(Th×k ) + E( n×k )

(1)

= Y( n×m ) U ( n×h ) * Q(Th×m ) + F( n×m )

(2)

where n refers to the number of observations of a property; k is the number of responses measured for each sample; m corresponds to the number of properties to be predicted

by PLS regression; h is the number of latent variables (LVs); T and U are the matrices of scores for X and Y data matrices, respectively; P and Q matrices present the inputs for X and Y, in that order; E and F matrices contain the residual errors for the prediction model.

To maximize the covariance between X and Y matrices, the scores are obtained from the linear combinations of the elements from the instrumental data set, using weight coefficients (w) and a given number of LVs[15]. In the conventional PLS regression, the elements in the T matrix (t), i.e. the scores, are estimated by: tnh = ∑xnk wkh k

(3)

where x is an elements in the X matrix.

Keeping the minimum modulus for F elements and the matrices of scores (T and U) internally correlated by U = T (i.e., X scores are assumed to be the most appropriate predictors for Y matrix)[13], the interest property set is predicted by PLS linear regression using: = Y( n×m ) X ( n×k ) * B( k ×m ) + G( n×m )

(4)

where G is a matrix of random errors and B is the matrix of model regression with the linear coefficients. Computationally, the PLS linear model is implemented by the NIPALS algorithm detailed in Figure 1.

The CARS algorithm was projected to interactively find the optimal subset, i.e., points in an instrumental data set (X matrix with the spectra data) to build the PLS regression with the lowest value of Root Mean Square Error of Cross Validation (RMSECV)[16]. At every sampling run, the CARS algorithm builds a PLS model with a randomly selected variable subset from the calibration set (Monte Carlo sampling method). The Exponentially Decreasing Function (EDF) and Adaptive Reweighted Sampling (ARF) are subsequently applied as a two-step method for wavelength selection to remove the wavelength (elements in the X matrix) that present the poorest weight coefficients (w) by a simulation of the “survival of the fittest” principle.

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Figure 1. Flow chart of the NIPALS algorithm[13].

In CARS-PLS, the importance of each x element is evaluated by a normalized weight calculated by: wCARS =

bi

∑ik=1 bi

(5)

While the ARF method keeps the x element with the largest weights, the EDF method induces the reduction of the number of x elements to build the PLS models with the small absolute regression coefficients by force. In each sampling run, EDF uses the following exponential model: ri = αe−βi

(6)

where: 1/ ( N −1)

k α =  2 β=

ln ( k / 2 )

( N − 1)

(7)

(8)

2. Materials and Methods 2.1 Materials Virgin LDPE (MFI of 2.6 g/10 min, 190 °C, 2.16 kg – ASTM D-1238) and HDPE (MFI of 0.3 g/10 min, 190 °C, 5 kg – ASTM D-1238) were obtained from Braskem and Petroquímica Triunfo S.A., respectively.

Polímeros, 29(1), e2019010, 2019

Predicting LDPE/HDPE blend composition by CARS-PLS regression and confocal Raman spectroscopy 2.2 Sample preparation Twenty-one LDPE/HDPE polymer blends with several concentrations (from 0/100, 5/95, 10/90, 15/85, 20/80, and so forth until the concentration of 100/0 wt%) were prepared by the extrusion process (Dynisco Laboratory Mixing Extruder, barrel diameter = 19 mm, diameter orifice = 3.12 mm, temperature profile = 180 and 190 °C, screw speed = 220 rpm). To simulate a mechanical recycling process and to ensure mixing, the samples were extruded three times before the spectroscopic analysis.

2.3 Apparatus and software Raman spectra of the HDPE/LDPE extruded pellets were obtained using confocal Raman Microscope Alpha300 R (WITEC, laser of 532 nm and 45 mW), and collected from 210 to 3785 cm-1 at room temperature with a spectral resolution of 3 cm–1.

All Raman data were smoothed using the SavitzkyGolay method[17] (polynomial order of 5, window points of 10) and previously normalized. CARS-PLS regression of the pre-processed spectra were carried out on MATLAB software (version R2015a) using libPLS 1.95 toolbox[18].

2.4 CARS-PLS regression analysis Forty-two spectra were used as a cross validation set, while sixteen spectra were used as an independent prediction set. The root mean squared errors were measured by[19]: RMSECV / RMSEP =

∑in=1 ( yi − yˆi ) n

2

(9)

shifts of the polymer chains of the polyethylenes (Table 1).

In sum, the Raman shifts at 1070, 1135 and 1300 cm−1 are from C-C stretching and -CH2- twisting. The medium Raman shift at 1445 cm-1 is associated to three -CH2- vibrational modes from the PE crystal structure (one wagging and two scissoring vibrational modes)[21]. The strong Raman shifts at 2845 cm-1 and 2883 are from the asymmetric and the symmetric stretching of the CH2 units, respectively[22].

The weak Raman shift at 480 cm-1 is from the molecular rotations of the C-C ramifications with four to nine carbons in gauche state[23].

In addition, Raman shifts derived from optical effects were identified on the Raman spectra of the LDPE/HDPE blends: 2725 and 2430 cm-1 are overtones and combinations of wavenumbers in the range of 1400-1495 cm-1 (-CH2- bonds)[24]; 2935 cm−1 (smooth shoulder) is reported to be from the Fermi resonance between the CH2 symmetric stretching and the overtone from the CH2 bond[25]. The Raman spectra of LDPE and HDPE are very similar, but it was observed that the maximum intensity of the Raman band at 1460 cm-1 increases with the reduction of LDPE in the polymer blend, while the opposite behavior is observed for the Raman shifts at 1370 and 1416 cm-1. These spectral characteristics are the basis for operation of the multivariate calibration to quantify the composition of the LDPE/HDPE blends using confocal Raman spectra data[20].

Table 2 presents the optimal predictive models built by the CARS-PLS algorithm using several statistical pretreatment methods for the Raman data (the number of latent variables for cross-validation, type of cross-validation and number of runs were maintained constant, as described

where n is the spectrum number; yi are the reference concentrations of the samples and yˆi are the concentrations predicted by the calibration set (RMSECV), or independent validation test (RMSEP), respectively.

The fitting degree between the predicted results and reference values was obtained by the correlation coefficient (R)[20]: R =

1−

∑in=1 ( yi − yˆi )

2

∑in=1 ( yi − yˆ mean )

2

(10)

where ŷ mean is the average polymer concentration of all samples in the cross validation and external test sets.

3. Results and Discussions The Raman spectra from the processed samples are shown in Figure 2 to represent all the compositions of the LDPE/HDPE blends (0-100 wt% of LDPE) and the characteristic Raman

Figure 2. Confocal Raman spectra from the HDPE/LDPE blends.

Table 1. Main Raman shifts of the LDPE and HDPE. Raman shift (cm-1) 1070, 1135 and 1300 1175 1372 1445 2845 2883

Bond -C-C- and -CH2-CH2-CH2-CH2-CH2-CH2-

Polímeros, 29(1), e2019010, 2019

Vibrational mode Stretch and twisting Rocking Wagging One wagging and two scissoring modes Asymmetric stretching Symmetric stretching

Phase Crystalline and anisotropic regions Crystalline Amorphous Crystalline Amorphous and crystalline Amorphous and crystalline

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Silva, D. J., & Wiebeck, H. in the table label). The pretreatment step is essential to reduce the negative effects of the Raman signal instability caused by sample fluorescence and laser instability. Here, we evaluated the Pareto, mean-centering and autoscaling methods; note that all Raman data were previously processed

by Savitzky‑Golay smoothing and normalization procedure before their pretreatment. The results indicate that the predictive model based on mean-centering (PLS-C) best fits the external test set (Rpred = 0.979), and also show the lowest prediction error (RMSPE = 4.062 wt% of LDPE).

Independently of the data pretreatment method, the CARS‑PLS regression has an excellent calibration performance, since the calibration errors in Table 2 are extremely low, lesser than 0.9 for all prediction models.

The data pretreatment by Pareto method minimizes the relative importance of large values, but it does not cause significant changes to the original spectral data. In the autoscaling method, the objective is give equal importance for all the spectral data, while the mean centering pretreatment consists of removing the offsets from the spectral data[26,27].

Table 3 summarizes the predictive PLS-based models with the lowest results of RMSECV and RMSEP, obtained by the CARS search algorithm and several K-fold values for cross-validation. Their correlation coefficients of calibration (Rcalib) and prediction (Rpred) were detailed as well. The K-fold cross-validation technique randomly divides the calibration dataset into K mutually exclusive subsets with the same size, i.e., with the same number of spectra. While K-1 subsets are applied to the training of the predictive model, one subset is used to calculate RMSECV and Rcalib (model testing).

Leave-one-out, which was used to analyze the pretreatment method effects on the CARS-PLS predictive models, is a specific case of K-fold cross-validation, where K is equal to the total number of spectra data (N). In this mathematical

approach, N calculations are performed, incurring expressive computational cost when N is high.

As can be seen in Table 3, there is no improvement in prediction and calibration performance of the CARS-PLS models using more than 5-fold, in which the RMSEP is equal to that obtained by leave-one-out cross-validation (RMSEP = 4.062 wt%). In this fold condition, the fitting degrees do also not display fluctuations for either calibration (Rcalib = 0.999) or prediction (Rpred = 0.979) datasets, while the optimal number of latent variables is 19.

According to Table 4, RMSECV decreased and Rcalib increased as the maximal numbers of latent variables for cross-validation increased, being the minimum result at 0.039 wt% for the PLS-40 model built with 34 LVs. However, the RMSEP results indicate a direct effect on the prediction performance of the CARS-PLS models due to an increase of the maximal LVs, since RMSEP falls from 8.017 wt% to 5.521 wt% of LDPE when the maximal number of LVs is enhanced from 5 to 10, respectively.

In a linear PLS regression, a projected vector subspace is assembled by a linear relationship between the latent variables in the spectral dataset. For this reason, the optimum number of LVs should be identified to obtain the best calibration performance for the PLS model[28]. The advantage of the CARS-PLS method is the possibility to conduct a sophisticated and automatic search to optimize this parameter without the need of excessive manual searches, required in the conventional linear and nonlinear PLS regressions[16].

In order to investigate the influence of the number of Monte Carlo sampling runs on the CARS-PLS model performance, predictive models with 50 to 10000 runs were built and they are shown in Figure 3 (it was kept constant the other parameters, i.e. pretreatment method, cross-validation

Table 2. Optimal predictive models obtained by CARS-PLS regression with the confocal Raman spectra pretreated by several methods (constant parameters: maximal number of latent variables for cross-validation = 20; cross-validation = leave-one-out; Monte Carlo sampling runs = 100).

Model

Pretreatment

LVoptimal

RMSECV (wt% of LDPE)

Rcalib

RMSEP (wt% of LDPE)

Rpred

PLS PLS-C PLS-A PLS-P

Centering Autoscaling Pareto

20 19 20 15

0.147 0.329 0.158 0.765

1.000 0.999 1.000 0.999

6.899 4.062 6.504 5.146

0.939 0.979 0.946 0.966

Table 3. Optimal predictive models obtained by CARS-PLS regression using several K-fold values for cross-validation (constant parameters: pretreatment = centering; maximal number of latent variables for cross-validation = 20; Monte Carlo sampling runs = 100).

Model

K-Fold for cross-validation

LVoptimal

RMSECV (wt% of LDPE)

Rcalib

RMSEP (wt% of LDPE)

Rpred

PLS-C2 PLS-C5 PLS-C15 PLS-C20

2 5 15 20

20 19 19 19

0.803 0.329 0.329 0.329

0.999 0.999 0.999 0.999

4.997 4.062 4.062 4.062

0.968 0.979 0.979 0.979

Table 4. Optimal predictive models obtained by CARS-PLS regression using several maximal numbers of latent variables for cross‑validation (constant parameters: pretreatment = mean centering; cross-validation = leave-one-out; sampling runs = 100).

Model

Maximal number of LVs

LVoptimal

RMSECV (wt% of LDPE)

Rcalib

RMSEP (wt% of LDPE)

Rpred

PLS-5 PLS-10 PLS-30 PLS-40

5 10 30 40

5 8 21 34

6.911 2.951 0.145 0.039

0.948 0.991 1.000 1.000

8.017 5.521 6.083 6.139

0.918 0.961 0.953 0.952

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Polímeros, 29(1), e2019010, 2019

Predicting LDPE/HDPE blend composition by CARS-PLS regression and confocal Raman spectroscopy algorithm and maximal number of latent variables for crossâ&#x20AC;&#x2018;validation). All show good correlation coefficients for the predicted values of the LDPE relative concentration from the external validation set (Rpred higher than 0.9) and excellent correlation coefficient for calibration (Rcalib = 0.999).

The highest calibration and prediction error was identified for the CAR-PLS model built with 50 runs, probably due to the low steps for searching the main Raman shifts to set up a predictive model with a robust predictive performance.

The results of the CARS-PLS models assembled with more

Figure 3. Reference vs. predicted LDPE relative amounts of the LDPE/HDPE blends obtained by CARS-PLS regression using several numbers of Monte Carlo sampling runs: (a) 50; (b) 100; (c) 500; (d) 1000; (e) 5000; and (f) 10000. (Constant parameters: pretreatment method = mean centering; cross-validation = leave-one-out; maximal number of latent variables for cross-validation = 20).

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Silva, D. J., & Wiebeck, H. than 100 sampling runs evidently show that the RMSECV reduction does not necessarily improve the predictive ability of the PLS model (reduction of the RMSEP value). Moreover, all the statistical errors of the CARS-PLS models are constant when more than 5000 sampling runs (RMSEP = 0.281 and RMSECV = 4.806 wt% of LDPE) are used.

All CARS-PLS-based models presented more significant prediction performance with the interval containing the Raman shift at 2883 cm-1 (both amorphous and crystalline polyethylene phases) and 1445 cm-1 (only from the PE crystalline phase). In a previous work with Interval PLS linear regression[20], we identified that the Raman signal at 2845 cm-1, which regards the CH2 asymmetric stretching in amorphous and crystalline phases, enables to obtain prediction models with the smallest RMSEP values (2.68‑6.94 wt% of LDPE). The most plausible justification is associated to the intensity and width of Raman shifts (1370, 1416 and 1460 cm-1), which are not just related to the content of the polymer chemical groups, but also to the macromolecular organization of the polymeric chains.

The difference between the branching degree of LDPE and HDPE affects the methylene polymer conformations in the amorphous and crystalline regions, directly influencing their molecular rotations and vibrations, intimately connected to the Raman signal detected by this vibrational spectroscopy.

4. Conclusions A modified PLS linear regression was used to predict the composition quantification of LDPE/HDPE blends.

The predictive PLS-based models presented the lowest prediction error of 4.062 wt% of LDPE with a good fitting coefficient of 0.979 in the whole content range, 0-100 wt% of HDPE. The CARS-PLS parameters display a significant role in the RMSECV and RMSEP of the predictive models.

In the conditions evaluated, the mean centering method for Raman data pretreatment favors the best prediction performances, while the autoscaling method benefits the lowest calibration errors. The increase in the K-fold and the maximal numbers of LVs for cross-validation caused a reduction of the RMSECV values, but RMSEP is not directly related with these regression variables. The optimal number of sampling runs was 100; above this value, the CARS-PLS models have a decrease in their potential to determine the LDPE relative amount in the polymeric blend.

5. Acknowledgements The authors would like to acknowledge the Coordination for the Improvement of Higher Education Personnel (CAPES) for their financial support.

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modes of vibration. Spectrochimica Acta. Part A: Molecular and Biomolecular Spectroscopy, 28(8), 1485-1496. http:// dx.doi.org/10.1016/0584-8539(72)80118-1.

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Received: Jan. 30, 2018 Revised: June 23, 2018 Accepted: Sept. 26, 2018

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ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.05218

Extraction and characterization of cellulose microfibers from Retama raetam stems Abdelkader Khenblouche1* , Djamel Bechki1, Messaoud Gouamid2, Khaled Charradi3, Ladjel Segni4, Mohamed Hadjadj5 and Slimane Boughali1 Laboratory of New and Renewable Energy in Arid Zones – LENREZA, University of Ouargla, Ouargla, Algeria 2 Department of Chemistry, Faculty of Mathematics & Material Sciences, University of Ouargla, Ouargla, Algeria 3 Laboratory of Nanomaterials and Systems for Renewable Energies – LaNSER, Research and Technology Center of Energy, Techno-park Borj-Cedria, Hammam-Lif, Tunis, Tunisia 4 Laboratory of Process Engineering, Faculty of Applied Sciences, University of Ouargla, Ouargla, Algeria 5 Laboratory of Valorization and Promotion of Saharian Resources – VPRS, University of Ouargla, Ouargla, Algeria 1

*khenblouche89@gmail.com

Abstract Cellulose is the most abundant renewable resource in nature, it has various industrial applications due to its promising properties. Retama raetam is a wild plant belonging to the Fabaceae family, largely abundant in arid area which makes it a good candidate for industrial utilization. In the present study, highly crystalline cellulose microfibers (77.8% CrI) were extracted from Retama Raetam stems as a novel renewable source. The samples underwent a dewaxing process, then the microfibers were extracted using 7 wt% sodium hydroxide followed by a bleaching treatment. The extracted cellulose microfibers were characterized by Scanning electron microscopy, Fourier transform infrared spectroscopy, X-ray Diffraction and thermo-gravimetric analysis.

Keywords: cellulose, microfibers, Retama raetam, extraction, characterization.

How to cite: Khenblouche, A., Bechki, D., Gouamid, M., Charradi, K., Segni, L., Hadjadj, M., & Boughali, S. (2019).

Extraction and characterization of cellulose microfibers from Retama raetam stems. Polímeros: Ciência e Tecnologia, 29(1), e2019011. https://doi.org/10.1590/0104-1428.05218

1. Introduction Over the last few decades, The use of natural fibers instead of synthetic fibers as reinforcement materials for polymer composites has gained considerable attention because of their unique characteristics, such as renewability, biodegradability, processing flexibility, low density, high specific strength and low-cost[1,2]. In addition, natural fibers have applications in various fields such as bioenergy industries, automobiles, paper manufacturing and textile owing to their properties and broad availability[1,3]. As a result, nowadays, the subject of many researchers worldwide focuses on the need to find alternative fiber sources[4]. Among all natural fibers, cellulose has attracted much interest as it is the most abundant renewable resource in nature and the degradation of cellulosic biomass is an important part of the biosphere’s carbon cycle[5]. Its existence as the common material of plant cell walls was first investigated by Braconnot in 1819[6] and Payen in 1838[7]. It is a polydispersed linear polymer with a microfibrillar structure composed of poly-b (1→4)-D-glucose units with a syndiotactic configuration[8,9], found in the cell walls as a network of microfibrils embedded in a non-cellulosic matrix[10]. Several plants are rich in cellulose, i.e. cotton,

Polímeros, 29(1), e2019011, 2019

wood, bamboo, hemp, flax, and jute …etc[11]. In addition, cellulose fibers have been extracted from several sources such as; milkweed stems[12], hop stems[13], rice husk[14], Cissus quadrangularis root[15], dichrostachys cinerea bark[1] and many more. Although several sources of natural fibers were investigated in detail, the isolation of cellulose fibers from R. Raetam has not been reported yet.

Retama raetam, locally known as R’tem, is a wild plant of the Fabaceae family. It is common to North and East Mediterranean regions[16]. It is largely abundant in arid area; this abundance makes it a good candidate for industrial utilization. Moreover, the Retama species contributes to the bio-fertilization of poor grounds because of their aptitude to associate with fixing nitrogen bacteria Rhizobia[17]. Therefore, the genus of the Retama is included in a re-vegetation program for degraded areas in semi-arid Mediterranean environments[18].

In this research, Natural micro-sized cellulose fibers were extracted from R. Raetam stems using alkali and bleaching treatments, the resultant cellulose microfibers were characterized using FTIR, SEM, XRD and TGA.

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

Khenblouche, A., Bechki, D., Gouamid, M., Charradi, K., Segni, L., Hadjadj, M., & Boughali, S.

2. Materials and Methods

2.4 Fourier transform infrared spectroscopy

2.1 Materials Stems of R. raetam subject of this study were collected in Ouargla, Algeria, in 2015. Acetone, Ethanol, Sodium hydroxide & Hydrogen peroxide were purchased from Sigma-Aldrich and were used without further purification.

To analyze the chemical changes of the samples and investigate functional groups in the extracted cellulose we used Fourier Transmission Infra-Red Spectroscopy. The FTIR spectra were recorded on a Cary 660 FTIR Spectrometer (Agilent Technologies, USA) in a wavelength range of 4000–600 cm-1 with a resolution of 4 cm-1.

2.2 Preparation of samples

2.5 X-ray Diffraction (XRD) analysis

Adult stems were cleaned with water and air dried, broken to the size of about 1 cm long and 1 mm width, grinded into powder with a Retsch SM100 Comfort cutting mill (Retsch GmbH, Haan, Germany), and sieved using a sieve size of 0.25 mm.

The crystallinity of cellulose microfibers was investigated

by X-ray diffraction (XRD) analysis, using a powder X-ray diffractometer (D8 Advance A25 Bruker AXS GbmH., Germany) with Cu Kα radiation (1.5406 A˚) at 40 kV and 25 mA, in the range of 2θ = 5-60° at a scanning rate of 0.02° s−1. The crystallinity index (CrI) was calculated according to Segal equation[19]:

2.3 Microfibers extraction The extraction of cellulose microfibers was performed using classical chemical treatments with adaptations in dewaxing, alkali and bleaching treatment processes.

The totally chlorine-free extraction procedure can be described as follows: 2.3.1 Dewaxing About 20g of powdered stems were first dewaxed in a Soxhlet reflux with a 2:1 (v/v) mixture of Acetone/Ethanol at 63 °C for 7 h, the main purpose of this step is to remove off waxes and extractives, the sample was then placed in a Buckner funnel and vacuum dried at room temperature for 3 h to remove traces of residual solvents.

2.3.2 Alkali treatment The alkali treatment was performed to purify the cellulose by removing lignin and hemicellulose from R. Raetam fibers.

The extractive-free sample was treated with an alkali solution (7 wt% NaOH) with a solvent to solid ratio of 10:1 at 80 °C for 3 h under mechanical stirring. This treatment was performed trice, after each treatment the solid was filtered and washed with distilled water until neutral pH. 2.3.3 Bleaching A subsequent bleaching treatment was carried out to remove residual lignin and whiten the microfibers. The sample was immersed in a hydrogen peroxide solution (11%, v/v), the pH was adjusted to 11 using 7 wt% NaOH, the system was vigorously stirred for 3 h at 45 °C. For a more effective discoloration, the bleaching process was performed twice under the same conditions, after each treatment, the microfibers were filtered and washed with distilled water.

CrI =100 × ( I 200 − I am ) / I 200

(1)

where I200 is the diffraction intensity at 2θ = 22-23˚; and IAM is the minimum diffraction intensity at 2θ = 18-20˚.

The crystallite size was calculated as per the Scherrer equation[20] Lh= , k ,l

( 0.94 ×λ ) / (β× cos θ )

(2)

where λ is X-rays wavelength; β is the full width at half maximum in radians; and θ is the Bragg angle.

2.6 Morphological structure A scanning electron microscope (SEM) (Quanta 250 FEG, FEI, USA) with an accelerating voltage of 15 kV was used to investigate the microstructure and the surface morphology of the obtained cellulose microfibers.

2.7 Thermogravimetric Analysis (TGA) In order to study the thermal stability of the extracted cellulose microfibers, thermogravimetric analysis (TGA) was performed using a Mettler Toledo TGA/DSC 3+ instrument.

The scan was carried out from 25 to 600 °C at a heating rate of 10 °C/min and under nitrogen atmosphere.

3. Results and Discussion 3.1 Extraction method and cellulose yield A stepwise totally chlorine-free procedure for the isolation of cellulose microfibers from retama raetam was proposed in this paper based on that adopted by Sun et al.[21]

Table 1. Comparison of the crystallinity index of cellulose microfibers from various sources.

Source Coconut palm leaf sheath Sugarcane bagasse Agave fibers Mengkuang leaves Soy hulls Commercial microcrystalline cellulose Sisal fibers Wheat straw Retama Reatam Hibiscus sabdariffa Palmyra palm Fruits

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Crystallinity index (CrI) 47.7% 50% 64.4% 69.5% 69.6% 73.91% 75% ±1 77.8% 77.8% 78.95% 81.9%

Reference Uma Maheswari et al.[22] Jonjankiat et al.[23] Reddy et al.[24] Sheltami et al.[3] Alemdar and Sain[25] Kale et al.[20] Morán et al.[26] Alemdar and Sain[25] This work Sonia and Priya Dasan[27] Reddy et al.[28]

Polímeros, 29(1), e2019011, 2019

Extraction and characterization of cellulose microfibers from Retama raetam stems with modifications. These modifications involve increasing sodium hydroxide and hydrogen peroxide concentrations as well as improving treatment time and/or temperature to enhance non-cellulosic components removal. Therefore, the isolation of R. Raetam cellulose microfibers was successfully achieved without any additional harsh acid treatments (Table 1), which makes the suggested extraction process not only eco-friendly and cost-saving, but also yielding cellulose microfibers of higher crystallinity and smaller diameters as further confirmed by XRD and SEM results.

Cellulose microfibers yield was gravimetrically determined (calculated as the percentage of the extracted cellulose microfibers over the initial raw sample weight) and was found to be 52.1%. This yield value is higher than that reported in literature for cellulose microfibers extracted

from Hibiscus sabdariffa fibers (38.6%)[27] and comparable to yield values of 52, 52 and 55% for cellulose microfibers extracted from Coconut palm leaf sheath[22], African Napier grass[29] and Palmyra palm fruits[28], respectively.

3.2 FT-IR spectroscopic analysis Infrared spectroscopy is currently one of the most important analytical techniques available to scientists[30].

It presents a relatively easy method of obtaining direct information on chemical changes that occur during chemical treatments[31]. Furthermore, FT-IR analysis was conducted to investigate the presence of different functional groups in the isolated samples. Figure 1. shows the IR spectra of (a) untreated sample, (b) alkali treated sample and (c) bleached cellulose microfibers. As summarized in Table 2,

Figure 1. FT-IR spectra of (a) untreated sample, (b) alkali treated sample and (c) bleached cellulose microfibers.

Table 2. The main observed IR bands and their assignments.

Spectra (a), (b), (c) (a), (b), (c) (b), (c) (a), (b), (c) (a), (b) (a), (b), (c) (a), (b), (c) (a), (b), (c) (a), (b), (c) (a) (a), (b), (c) (a), (b), (c)

Wavenumber (cm-1) 876-897 1029-1031 1051-1054 1144-1159 1224-1236 1316-1327 1370-1374 1417-1429 1644-1646 1734 2896-2922 3323-3332

Polímeros, 29(1), e2019011, 2019

Assignment C–O–C stretching at the b-(1→4)-glycosidic linkages C–O–C pyranose ring skeletal vibration C–O–C pyranose ring skeletal vibration C–O–C stretching at the b-(1→4)-glycosidic linkages –COO vibration of acetyl groups / C–O stretching of the aryl group C–C and C–O skeletal vibrations O–H bending CH2 bending OH bending of the absorbed water C=O stretching C–H stretching H-bonded OH groups stretching

Ref. Oh et al.[32] Sun et al.[33] Sun et al.[33] Oh et al.[32] Reddy et al.[28] Gao et al.[34] Gao et al.[34] Gao et al.[34] Alemdar and Sain[25] Sain and Panthapulakkal[35] Kondo and Sawatari[36] Kondo and Sawatari[36]

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Khenblouche, A., Bechki, D., Gouamid, M., Charradi, K., Segni, L., Hadjadj, M., & Boughali, S. the cellulose spectrum is similar to those reported in literature for cellulose fibers[37].

The band at 3332 cm-1 relates to the stretching of H-bonded OH groups, and the one at 2901 cm-1 to the C–H stretching[36], we observe that the band around 3332 cm-1 is narrower and has a higher intensity for cellulose, which demonstrated that the extracted cellulose contained more –OH groups than in untreated sample[10]. The band at 1644 cm-1 is associated with –OH bending of the absorbed water[25]. Typical bands assigned to cellulose at 1159 cm-1 and 897 cm-1 are due to C–O–C stretching at the b-(1→4)-glycosidic linkages[32].

The presence of these peaks showed the increase in the percentage of cellulosic components after the removal of non-cellulosic materials by chemical treatments[38].

The absorption peak at ~1734 cm-1 on the spectrum of the untreated sample (a) is attributed to the C=O stretching of the acetyl and uronic ester groups of polysaccharides[9,35], it is also related to the p-coumeric acids of lignin and/or hemicellulose[25], the absence of this peak after successive chemical treatments indicates the removal of most of lignin and hemicelluloses from the microfibers. Another indicator of lignin and hemicellulose removal during the chemical treatments is the significant decrease in the intensity of the peak around 1236 cm-1 which is related to the –COO vibration of acetyl groups in hemicellulose and/or the C–O stretching of the aryl group in lignin[28]. Noticeable peaks on spectrum (c) at 1429 cm-1 relates to the CH2 bending and at 1370 cm-1 to the O–H bending. The absorbance at ~1316 cm-1 is attributed to the C–C and C–O skeletal vibrations[34]. The C–O–C pyranose ring skeletal vibration occurs in 1054 cm-1 and 1031 cm-1[33].

3.3 X-ray Diffraction measurements Figure 2 exhibits the XRD data of the extracted cellulose microfibers. The cellulose amorphous phase is characterized by the low diffracted intensity at a 2θ value of 19.12˚, whilst the peaks at 15.14˚, 16.25˚, 22.75˚ and 34.39˚ are attributable to the crystallographic planes of (1-10), (110), (200) and (004), respectively, which are characteristic of the typical cellulose I structure[29,39]. The crystallinity index (CrI) of the bleached microfibers was determined using Segal equation and was found to be 77.8%. Obviously, as illustrated in Table 1, this CrI value is higher than the value of 73.91% reported by Kale et al. for commercial microcrystalline cellulose from wood pulp[20]. Moreover, it is higher than the CrI values reported in the literature for cellulose microfibers isolated using different methods from various sources such as Sugarcane bagasse (50%)[23], Agave fibers (64.4%)[24], mengkuang leaves (69.5%)[3] and Soy hulls (69.6%)[25].

Furthermore, the crystallite size (Lh,k,l) which was found to be 3.62 n‎ m is comparable to the reported sizes of cellulose crystallites (4 to 7 nm generally)[40]. However, the calculated Lh,k,l value is much lower than that reported for hydrolyzed commercial microcrystalline cellulose (10.32 nm).

The higher CrI and the lower Lh,k,l values suggest that the adopted stepwise chlorine-free extraction treatments were effective in removing most of amorphous domains leading to break the bundles of cellulose fibers to form smaller cellulose crystallites[41,42].

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Figure 2. XRD diffractogram of the extracted R. Raetam cellulose microfibers.

3.4 Morphological properties of chemically purified cellulose microfibers Figure 3 shows SEM micrographs of the chemical‑purified cellulose microfibers. After they had been subjected to alkaline solution treatment and bleaching, the cellulose microfibers were separated into individual micro-sized fibers. These micro-sized cellulose fibers were reported to be composed of strong hydrogen bonding nanofibers[43]. The diameter of the microfibers is about 6-7 μm but the exact determination of their length is difficult.

As Table 3. demonstrates, the extracted R. Raetam cellulose microfibers are smaller in diameter as compared to those isolated by different extraction methods from various sources such as sisal fibers, agave fibers, coconut palm leaf sheath, soy hull and wheat straw[22,24-26]. Moreover, they are comparable to cotton and sugarcane bagasse microfibers extracted by sulfuric and nitric acid hydrolysis, respectively[23,44]. This morphology and smaller diameter would enable R. Raetam cellulose microfibers to be used for various applications ranging from reinforcing agents in biodegradable composites, to gel forming food and cosmetic additives[24,45].

3.5 Thermal stability Investigating thermal properties of cellulose microfibers is a key factor for their applicability as reinforcing fillers in biocomposites[38]. Figure 4a and 4b shows, respectively, the thermogravimetric analysis (TGA), and the derivative thermogram (DTG) curves for both the raw and bleached samples. TGA curves show an initial weight loss below 155 °C, this initial drop (4.6% for raw sample and 6.5% for cellulose microfibers) was due to the evaporation of moisture bounded on the surface of the samples, chemisorbed water bounded inside the samples and/or the compounds of low molecular weight such as extractives presented in the raw sample[38,46], the presence of the absorbed water was affirmed previously through the FT-IR results. The main broader cellulose thermal degradation (50.84%) occurs over 179 °C and involves synchronous multi-processes such as dehydration, depolymerization and decomposition Polímeros, 29(1), e2019011, 2019

Extraction and characterization of cellulose microfibers from Retama raetam stems Table 3. Comparison of the diameters of cellulose microfibers extracted from various sources by different extraction methods.

Source Sisal fibers Coconut palm leaf sheath Soy hull Wheat straw Agave fibers Hibiscus sabdariffa Sisal fibers Cotton Sugarcane bagasse Palmyra palm fruit Napier grass fibers Retama Raetam Stems Jatropha Curcus L seed shell

Extraction method Alkali, peroxide and HNO3/HAc treatments Chlorination, alkali and HNO3/HAc treatments Alkali treatment and HCl Acid hydrolysis Alkali treatment and HCl Acid hydrolysis Chlorination, alkali and HNO3/HAc treatments Steam explosion and oxalic acid hydrolysis Chlorination and alkali treatments Sulfuric acid hydrolysis Nitric acid hydrolysis Chlorination, alkali and HNO3/HAc treatments Chlorination, alkali and HNO3/HAc treatments Alkali and alkaline peroxide treatments Chlorination, alkali and HNO3/HAc treatments

diameter (μm) 12.8-31 10-15 10-15 10-15 8-14 10.04 7-11.2 5-10 5-10 4-11 8.3 6-7 0.23-1.04

Ref. Morán et al.[26] Uma Maheswari et al.[22] Alemdar and Sain[25] Alemdar and Sain[25] Reddy et al.[24] Sonia and Priya Dasan[27] Morán et al.[26] Chatterjee et al.[44] Jonjankiat et al.[23] Reddy et al.[28] Reddy et al.[29] This work Puttaswamy et al.[45]

Figure 3. SEM images of the extracted R. Raetam cellulose microfibers.

of glycosidic units[47]. The raw sample showed separated pyrolysis processes within a wider temperature range, including thermal depolymerization of hemicellulose up to 273 °C, decomposition of cellulose up to 348 °C, and Polímeros, 29(1), e2019011, 2019

the degradation of lignin up to 536 °C in addition to its simultaneous decomposition with other degradation stages due to its complex structure[46]. DTG curves exhibited maximum decomposition rates at DTGmax=294 °C and 5/8

Khenblouche, A., Bechki, D., Gouamid, M., Charradi, K., Segni, L., Hadjadj, M., & Boughali, S.

Figure 4. TGA (a) and DTG (b) curves of the untreated sample and the bleached cellulose microfibers.

Table 4. DTGmax and char yields of cellulose microfibers from different sources.

Samples Date seeds cellulose microfibers Bamboo cellulose microfibers Rice hulls microcrystalline cellulose bean hulls microcrystalline cellulose Onion skin cellulose microfibers Retama Raetam cellulose microfibers Cotton silver microcrystalline cellulose Jute microcrystalline cellulose

Tmax (◦C)

Char (%)

Reference

300 328 283 281 333 311 340 280

11 13 23 25 26 42 57 61

Nabili et al.[48] Chen et al.[49] Adel et al.[50] Adel et al.[50] Reddy and Rhim[51] This work Das et al.[52] Das et al.[52]

311 °C for the raw and bleached samples, respectively.

A shoulder can be clearly observed at 261 °C on the left side of the main peak of the raw sample DTG curve, which was due to initial decomposition of hemicellulose and non‑cellulosic components[47], while the broadening at 245 °C on the microfibers DTG curve, could be an indicator of a broad distribution of molecular mass from cellulose or a residual content of hemicellulose which withstood the extracting procedures[26]. Finally, the formation of a charred residue took place (46% for raw sample and 42% for cellulose microfibers). The higher charred residue of the raw sample is due to the fact that the non-cellulosic components could induce higher char formation[38].

However, R. Raetam cellulose microfibers presented relatively high char yield when compared to literature (Table 4), indicating higher non‑volatile carbonaceous material generated on pyrolysis[53] and could indicate also a good thermal stability of the extracted cellulose microfibers[20,54].

4. Conclusion The main goal of this work was to investigate the viability of Retama Raetam as a novel, renewable and low-cost source of cellulose microfibers. The successful isolation of cellulose microfibers was achieved with a yield of 52.1% by stepwise chemical treatments. The FTIR results revealed that the chemical treatments removed most of lignin and 6/8

hemicellulose from the sample. The extracted cellulose microfibers were highly crystalline native cellulose I, with a crystallinity of 77.8% and a crystallite size of 3.62 n‎ m.

The diameter of the micro-sized fibers was about 6-7 μm.

TGA/DTG curves show a maximum decomposition peak at 311 °C and a high char yield. These findings proved that R. Raetam is a candidate renewable source for the production of cellulose microfibers and should stimulate further research on the use of these fibers for various applications such as cellulose nanocrystals preparation, reinforcement agent in green biocomposites and bio-fillers for polymer matrices.

5. References 1. Baskaran, P., Kathiresan, M., Senthamaraikannan, P., & Saravanakumar, S. (2018). Characterization of new natural cellulosic fiber from the bark of dichrostachys cinerea. Journal of Natural Fibers, 15(1), 62-68. http://dx.doi.org/10.1080/15 440478.2017.1304314.

2. Mohammed, L., Ansari, M. N. M., Pua, G., Jawaid, M., & Islam, M. S. (2015). A review on natural fiber reinforced polymer composite and its applications. International Journal of Polymer Science, 15. http://dx.doi.org/10.1155/2015/243947.

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Received: Aug. 03, 2018 Revised: Sept. 21, 2018 Accepted: Oct. 23, 2018

PolĂ­meros, 29(1), e2019011, 2019

ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.02718

Evaluation of the dispersion of carbon nanotubes in an elastomeric polyurethane and fatigue test Felipe Luiz Queiroz Ferreira1, Magnovaldo Carvalho Lopes1, Ana Paula Mendes Lopes1, Rodrigo Lassarote Lavall1 and Glaura Goulart Silva1*  1

Centro de Tecnologia em Nanomateriais, Instituto de Ciências Exatas, Universidade Federal de Minas Gerais – UFMG, Belo Horizonte, MG, Brasil *glaura@qui.ufmg.br

Abstract Two series of polyurethane (PU) and carbon nanotubes (CNT) based composites with 0.0, 0.25, 0.5 and 1.0 mass% of CNT were obtained from diluting a commercial masterbatch with 30 mass% CNT and using two different dispersion methods. The quality of the dispersions was assessed using optical microscopy, and scanning and transmission electron microscopies. These tests showed that high controlled shear stress is necessary to produce composites with nanoscale dispersion: the elastic modulus improved by an average of 38% in the case of the high-shear dispersed materials in comparison with the neat polymer. A specific fatigue test conducted by dynamic mechanical analysis was first used in this work to compare the neat PU with the CNT/PU nanocomposites. The number of cycles to failure increased from 2700 for the neat polymer to 3200 for the 0.5 mass% CNT based nanocomposite; the elongation at failure increased by 145% in the test conditions.

Keywords: carbon nanotubes, elastomeric polyurethane, mechanical properties, fatigue.

How to cite: Ferreira, F. L. Q., Lopes, M. C., Lopes, A. P. M., Lavall, R. L., & Silva, G. G. (2019). Evaluation of the dispersion of carbon nanotubes in an elastomeric polyurethane and fatigue test. Polímeros: Ciência e Tecnologia, 29(1), e2019012. https://doi.org/10.1590/0104-1428.02718

1. Introduction The extraordinary properties of carbon nanotubes (CNT)[1-3] have motivated large efforts to apply them as reinforcing agents in polymer composites[4,5]. The success, however, is being limited by the difficulties to disperse the nanomaterial and promote a good interaction with polymer chains at the interfaces. Various strategies using different processing equipment/methods, have been tested to improve dispersion: high shear mixers, three roll-mill, or planetary mixers in the case of resin to thermosets[6-10]; and extrusion and injection in the case of thermoplastics[11,12]. Companies commercialize masterbatches with high concentrations of CNT up to 50 mass%. The high concentration of CNT increases significantly the viscosity of the masterbatches, leading liquid resins to be stored and commercialized as solid pellets. Masterbatches of CNT are available for thermoplastic (polyethylene, polypropylene, etc.) and for thermoset resins (epoxy and polyurethane).

Abbasi et al.[13] produced composites of polypropylene (PP) and carbon nanotubes from a 20 mass% commercial masterbatch using a twin-screw extruder for dilution. They reported that processing at low temperature and high speed increased the shear mixing, which in turn broke off the agglomerates of CNT, resulting in a better dispersion of the nanotubes and increase in the electrical conductivity of the composites. Pötschke et al.[14] used a Haake mixer coupled with a twin-screw extruder to produce composites of polycarbonate and carbon nanotubes from a commercial

Polímeros, 29(1), e2019012, 2019

masterbatch with 15 mass%. The authors obtained SEM images showing the CNT covered by the polymer and randomly oriented. Cryofracture image of those composites after dilution shows that the cracks propagated through the matrix. The carbon nanotubes, however, appear to have bridged the cracks in the matrix, reinforcing the strength of the composite. Mansour et al.[15] produced composites of thermoplastic polyurethane (TPU) and carbon nanotubes from a 10% mass masterbatch to evaluate mechanical properties such as hardness and compression strength. They showed that the use of carbon nanotubes increased hardness by up to 13% and modulus by up to 256%. Additional reports featuring nanocomposites of thermoset polyurethane based on commercial masterbatchs have not been found. As such, to the best of our knowledge, this is the first work on this subject.

Polyol resins are used as precursors of polyurethane elastomers (PU). This highly viscous liquid at room temperature can receive CNT in polyol masterbatches, which reacts with isocyanates producing PU. PU elastomers are versatile materials that can be employed in industrial equipment such as bend stiffeners in the off shore oil industry and conveyor belts in the mining industry[16,17]. For these industrial applications, there are complex requirements of mechanical properties, especially tear strength, abrasion resistance, and fatigue behavior. Therefore, they can benefit from new formulations that could show real gain in different

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

Ferreira, F. L. Q., Lopes, M. C., Lopes, A. P. M., Lavall, R. L., & Silva, G. G. properties simultaneously. The introduction of small amounts of CNT has the potential to increase performance for this kind of highly demanded materials. This was already shown by our previous work[9,10,18], as well as that of other authors[19-21].

The PU elastomer chains are composed by the polyol segments and the isocyanate segments, which contribute to the flexibility, mechanical resistance, and fatigue resistance of these materials[22-24]. Conventional fatigue tests are conducted in servo-hydraulic testing machines able to apply cyclic forces to specimen and measure the number of cycles to failure. This type of testing is frequently time-consuming and expensive. A new fatigue test for the PU elastomers and its nanocomposites using a dynamic mechanical analyzer (DMA) was tested in this work. Similar fatigue tests were employed for metals[25] or other polymers such as PTFE and SBR[26,27], but it is the first time that the fatigue of a PU elastomer nanocomposite was evaluated by DMA. Two types of dispersion methodologies were thus investigated: the dispersion proposed by the supplier of a commercial masterbatch of CNT and the dispersion optimized by our group, employing three-roll mills[9,10]. The quality and mechanical properties of these resulting dispersions were compared and a fatigue test by DMA was used to investigate the influence of CNT, under cyclic effort, in the lifetime of the material.

2. Materials and Methods 2.1 Materials A masterbatch of multiwalled carbon nanotubes was purchased from Arkema (France)[28]. Supplier reports that the masterbatch contains 30 mass% of CNT in polyether polyol resin (Graphistrength C PU1-30). Plastiprene (Brazil) kindly provided polytetramethylene ether glycol (PTMEG), toluene diisocyanate (TDI), 1,4-butanediol, and 4,4-methylene-bis-ortho-chloroaniline (MOCA) for use in the synthesis of the elastomeric polyurethane.

2.2 Preparation of MWCNT/PU composites Two series of composites were prepared by adding different amounts of the masterbatch in PTMEG and reacting with TDI and 1,4 butanediol. The prepolymer synthesized was cross-linked with MOCA[9,10]. The final concentrations of the composites were 0.25, 0.50 and 1 mass% of carbon nanotubes. Two types of dispersion methodologies were tested to disperse the masterbatch in the PTMEG. In the first instance, as proposed by the supplier, composites were prepared by adding the masterbatch under mechanical stirring (IKA, RW20, 350 RPM, 100 °C, overnight). In the second one, composites were prepared by adding masterbatch using a high shear mechanical mixer (IKA, T25, 20,000 RPM, 5 min) and subsequently with a three roll mill (Exakt, 80E, 5 runs, gap 5:10, 300 RPM). Two reference samples (without carbon nanotubes) were also processed using the two methods.

The procedures of masterbatch dispersion are summarized in Table 1.

2.3 Characterization The characterization of the samples included a study of the morphology of the masterbatch using scanning electron microscopy (SEM) in a Quanta 200-FEG / FEI microscope; the cryo-fracturing of a masterbatch pellet with liquid nitrogen to analyze the surface; and the examination of carbon nanotubes extracted from the masterbatch with acetone by transmission electron microscopy (TEM) on a Tecnai G2-20 / FEI microscope. Additionally, thermogravimetric analyses of the masterbatch were performed in a synthetic air atmosphere with a heating rate of 10 °C min-1 from room temperature to 1000 °C, using the Q5000 TA instruments apparatus. Samples of the masterbatch dilutions in PTMEG (0.10 mass%), prior to cure, were characterized by optical microscopy (OM) in transmission mode using an Olympus BX50F microscope with a magnification of 100x. The same OM analysis was also done for composite fragments in order to evaluate the dispersion of the nanotubes in the final material. Furthermore, the mechanical properties of the composites were tested in an Emic DL10.000 universal testing machine; elastic modulus was asserted using ASTM D638 and replicates of five dog bone specimens, type IV, with 115 mm (length), 6 mm (width) and 3 mm (thickness).

The secant modulus at 6% elongation was used for evaluation.

The tests were performed at 23 ± 2 °C and at relative humidity of 50 ± 10% using an extensometer, a load cell of 50 kgf, and a displacement velocity of 50.0 cm.min-1.

Finally, a fatigue test conducted using the tensile mode DMA was developed for this work. The test was performed under isotherm at 50 °C, the specimens were oscillated with a frequency of 5 Hz and a force of 6 N until rupture.

The specimen’s dimensions were 20 mm (length) × 6 mm (width) × 0.5 mm (thickness) and they were notched to induce crack propagation. Five specimens of each sample were tested and the two extreme results were excluded, thus, the data presented is in triplicates. Several experiments, using different temperatures, were conducted with specimens of different dimensions, with different frequencies and forces, and without/with notch to find the test parameters that revealed the fatigue behavior of the neat elastomeric PU and nanocomposites studied in this work. The extensive preliminary work (not reported here) showed that within the low limit of force that the DMA allows (6 N), it is necessary to introduce the notch and prepare thin specimens in order to get a clear fatigue response in this kind of PU.

3. Results and Discussions Figure 1a-c shows: (a) TGA analyses with a mass loss of 29% at approximately 550 °C, which is associated to the decomposition of carbon nanotubes confirming the nanotube

Table 1. Procedures of masterbatch dispersion.

Method 1 Method 2

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Dispersion methods Dispersion with mechanical stirring, 350 RPM, 100 °C, overnight.

Dispersion with high shear, 20.000 RPM for 5 min and three-roll mill, 5 runs, gap 5:10, 300 RPM.

Polímeros, 29(1), e2019012, 2019

Evaluation of the dispersion of carbon nanotubes in an elastomeric polyurethane and fatigue test content reported by the supplier[28]; (b) the appearance of the masterbatches in the form of pellets with a high concentration of carbon nanotubes; and (c) the cryo-fractured image of a pellet, in which CNT can be visualized pulling off the matrix, respectively. Figure 1d shows a TEM image of the carbon nanotubes extracted from the masterbatch with acetone, displaying nanotubes with an average of 10 nm of external diameter. The characterization of the masterbatch and the MWCNT in the material indicated that this commercially available system has desirable features for the employment as additive in elastomeric PU developed for industrial demands[9,10,13] Figure 2 shows optical microscopy images of diluted samples of the masterbatchs in PTMEG with 0.10 mass% of CNT, prior to cure, produced through the two procedures presented in this work (Table 1). The method of dispersion 1 using only mechanical stirring was recommended by the supplier, however, it was not enough to disperse the carbon nanotubes that remained agglomerated in small grains, as indicated in Figure 2a by the arrows highlighting the aggregates.

Method 2, showcased the need of high and controlled shear stress to produce composites with nanoscale dispersion, Figure 2b. The morphology of Figure 2b typically results from the shearing of the CNT aggregates by the roll mill; it

allows the formation of nanotube networks in all directions, consequently permitting improvements in the mechanical and electrical properties of the nanocomposite[29].

Figure 3a shows optical images of composite fragments with 0.25 and 0.50 mass% of CNT and (b) SEM images of composite fragments with 0.50 mass% of CNT, which show how efficient dispersion, obtained by processing method 2, was maintained even after curing. The larger black spots of CNT aggregates, which are present after method 1 dispersion, are observable at the top of Figure 3a, some of which have been indicated by arrows. Much smaller aggregates are observed in composites produced by method 2 as indicated by the arrows in the lower half of Figure 3a. SEM images furthermore confirm these observations: for the method 1 dispersion, the high magnified image (top-right of Figure 3b) shows the inside of an agglomerate of approximately 10 μm diameter. The SEM images in two magnifications on the bottom of Figure 3b show the typical dispersion achieved with the high shear method 2 with carbon nanotubes better distributed and pulled out of the PU matrix.

The composites processed by the most efficient dispersion method (method 2), as demonstrated by the morphological study, were tested for tensile strength and Figure 4b shows

Figure 1. Commercial masterbatch: (a) TGA curve; (b) inset - masterbatch sample; (c) SEM image of the cryofracturated pellet; (d) TEM image of the carbon nanotube extracted from the masterbatch.

Figure 2. Optical microscopy images for the dispersion (prior to cure) of masterbatch on PTMEG to a 0.10 mass% composite: (a) Method 1: dispersion by mechanical stirring; and (b) Method 2: dispersion with high shear and three-roll mill.

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Ferreira, F. L. Q., Lopes, M. C., Lopes, A. P. M., Lavall, R. L., & Silva, G. G.

Figure 3. (a) Optical and (b) SEM images of CNT/PU composite fragments prepared with commercial masterbatch dilution by two different methods. Method 1 (top) is mechanical stirring and Method 2 (bottom) is high shear /roll mill processing. The concentrations of CNT in PU are indicated in the bottom of the figure.

Figure 4. Results of mechanical testes for PU and nanocomposites prepared by Method 2: (a) Engineering stress-strain representative curves (from a minimum of five replicates); (b) Secant modulus at 6% elongation for elastomeric neat PU and composites.

the secant elastic modulus at 6% elongation. The elastic modulus increased up to 38% in comparison with the reference sample. Increases in modulus were also observed

by Lopes et al.[10], a study in which up to 47% increase in the stiffness of the composites was obtained with addition of 0.5% MWCNT in PU. Mansour et al.[15] obtained similar increases in composite stiffness with 1 mass% of nanotubes in a thermoplastic polyurethane formulation.

DMA analysis presented in Figure 5 shows that the storage modulus of the composite is similar to the neat PU at room temperature. Nevertheless, for temperatures above 50 °C, it is noted that storage moduli of 1 mass% composites are superior to the reference by up to 20% at 120 °C, for example. Tan δ curves show a large event associated to glass transition from -50 °C to 100 °C. The main peak of tan δ is found at 50 °C and a shoulder peak can be seen at – 30 °C. Neat PU and 1 mass% CNT based composite both show 4/7

no variation in the peak temperatures at the tan δ curve.

The DMA curves obtained in this study are very similar to the ones showed in the work of Lopes et al.[10]. In that case, the same morphology of elastomeric PU employed in the present work was produced with the addition of MWCNT by high shear mixing. The difference is that herein the CNT was added with the help of a commercial masterbatch. Therefore, if high shear dispersion methodology is applied, the results obtained with elastomeric PU/MWCNT are reproducible.

In an effort to assess this important characteristic of an elastomeric material such as PU, fatigue behavior tests

by DMA were also performed in this study. The industrial applications of elastomeric PU in oil, gas and mining industries are very demanding on the material because they are used in harsh conditions of impact, high temperature and need long life cycles of sometimes over a decade. Figure 6a, b presents images of the test pieces used in these experiments as well as Polímeros, 29(1), e2019012, 2019

Evaluation of the dispersion of carbon nanotubes in an elastomeric polyurethane and fatigue test

Figure 5. DMA results. Storage modulus, Loss modulus and Tan δ of neat elastomeric polyurethane and nanocomposite with 1.0 mass% of CNT prepared with a commercial masterbatch dilution by Method 2.

Figure 6. (a) Fatigue samples; (b) tensile probe apparatus; and (c) fatigue behavior for neat PU and nanocomposites.

the tensile device used in the tests. In Figure 6c, the fatigue behavior of the samples was examined based on the results of the % of elongation at failure as a function of number of cycles to failure. The results showed a dispersion of data allowing only a qualitative comparison between the neat PU and the nanocomposites. We have decided to present the three more consistent results which are marked in the image with an ellipsis rather than averaging. Only after 2400 cycles do the notched-thin specimens start to fail in conditions of 50 °C, tensile of 6 N and 5 Hz. The neat PU showed approximately 2700 cycles up to failure with about 10% elongation. It is clear that the composites presented higher % of deformation before failure, with significant increases of up to 245% in elongation for 1.0 mass% composite in comparison with the neat polymer. Composites with 0.25 and 0.50 mass% of nanotubes show an increase in the number of cycles to failure of 7 and 19%, respectively. The average of the results showed in Figure 6 are summarized in Table 2.

The fatigue test employed in this work showed a clear tendency for expanded lifespan and increased ductility for the samples with additions of carbon nanotubes. Although the conditions of this test by DMA make it difficult to compare with other references, the observation of the results allows Polímeros, 29(1), e2019012, 2019

Table 2. Average of fatigue test results.

Sample

Cycles to failure

0% 0.25% 0.50% 1.00%

2722 ± 81 2910 ± 146 3232 ± 44 2556 ± 60

Elongation to failure / % 11 ± 2 27 ± 7 25 ± 10 38 ± 4

for confirmation of effective gains in the fatigue resistance of the nanocomposites. Loos et al.[19] have been studied the effects of carbon nanotube (CNT) inclusion on cyclic fatigue behavior using a convectional tension fatigue test and the tensile properties of polyurethane (PU) composites. They have shown a similar tendency as what was observed in the present work, i.e. the incorporation of CNT extended the life of PU before fatigue in the high-stress amplitude and low-cycle regime by up to 248%. Loos et al.[19] state that the micrographs indicate the key mechanisms for enhancement in fatigue life as CNT crack-bridging and pull-out. Moreover, the effective enhancement in the properties of PU obtained in this study is attributed to successful dispersion of the CNTs and improved interaction between filler and matrix, suggesting an energy absorbing mechanism responsible for the increase in the fatigue life.

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Ferreira, F. L. Q., Lopes, M. C., Lopes, A. P. M., Lavall, R. L., & Silva, G. G.

4. Conclusions In this study carbon nanotubes were used as efficient mechanical reinforcement for elastomeric polyurethane.

This is seen by the increase in elastic modulus and storage modulus. It was also observed in the significant contribution of the carbon nanotubes for the increased fatigue strength, i.e. elongation and number of cycles at failure. This reinforcement was possible because of an efficient dispersion method for a commercial masterbatch achieved using techniques that generated high shear and produced composites with nanoscale dispersion. The nanocomposites PU/CNT have mechanical characteristics that make them possible substitutes in applications that demand high performance materials.

5. Acknowledgements The authors thank the Petrobras Company for support.

They also thank Daniel Freller from Plastiprene, a Brazilian producer of polyurethane parts. Pró-Reitoria de Pesquisa of UFMG, CNPq, Fapemig and the Microscopic Center (UFMG) are also acknowledged for their support.

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Exceptionally high young’s modulus observed for individual carbon nanotubes. Nature, 381(6584), 680-687. http://dx.doi. org/10.1038/381678a0.

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5. Ribeiro, B., Botelho, E. C., & Costa, M. L. (2015). Estudo das propriedades elétricas e térmicas de compósitos nanoestruturados de poli(sulfeto de fenileno) reforçados com nanotubos de carbono. Polímeros Ciência e Tecnologia, 25(1), 94-100. http:// dx.doi.org/10.1590/0104-1428.1728.

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7. Yedra, Á., Gutiérrez-Somavilla, G., Manteca-Martínez, C., González-Barriuso, M., & Soriano, L. (2016). Conductive paints development through nanotechnology. Progress in Organic Coatings, 95, 85-90. http://dx.doi.org/10.1016/j. porgcoat.2016.03.001.

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9. Lopes, M. C., Trigueiro, J. P. C., Castro, V. G., Lavall, R. L., & Silva, G. G. (2016). Otimização do processo de dispersão 6/7

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Polímeros: Ciência e Tecnologia, 26(1), 81-91. http://dx.doi. org/10.1590/0104-1428.2087.

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12. Müller, T. M., Krause, B., Kretzschmar, B., & Pötschke, P. (2011).

Influence of feeding conditions in twin-screw extrusion of PP/ MWCNT composites on electrical and mechanical properties.

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Dynamic mechanical characterization of polyurethane/multiwalled carbon nanotube composite thermoplastic elastomers.

Polymer-Plastics Technology and Engineering, 56(14), 15051515. http://dx.doi.org/10.1080/03602559.2016.1277243.

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Polyurethanes: versatile materials and sustainable problem solvers for todays challenges. Angewandte Chemie International Edition, 52(36), 9422-9441. http://dx.doi.org/10.1002/ anie.201302766. PMid:23893938.

17. Aquino, F. G., Sheldrake, T., Clevelario, J., Pires, F., & Coutinho, F. M. B. (2010). Estudo do envelhecimento de poliuretanos aplicados na indústria de petróleo. Polímeros: Ciência e Tecnologia, 20(1), 33-38. http://dx.doi.org/10.1590/ S0104-14282010005000006.

18. Lima, A. M. F., Castro, V. G., Borges, R. S., & Silva, G. G. (2012). Electrical conductivity and thermal properties of functionalized carbon nanotubes/polyurethane composites.

Polímeros Ciência e Tecnologia, 22(2), 117-124. http://dx.doi. org/10.1590/S0104-14282012005000017.

19. Loos, M. R., Yang, J., Feke, D. L., Manas-Zloczower, I., Unal, S., & Younes, U. (2013). Enhancement of fatigue life of polyurethane composites containing carbon nanotubes.

Composites. Part B, Engineering, 44(1), 740-744. http://dx.doi. org/10.1016/j.compositesb.2012.01.038.

20. Shokry, S. A., El Morsi, A. K., Sabaa, M. S., Mohamed, R. R., & El Sorogy, H. E. (2015). Synthesis and characterization of polyurethane based on hydroxyl terminated polybutadiene and reinforced by carbon nanotubes. Egyptian Journal of Petroleum, 24(2), 145-154. http://dx.doi.org/10.1016/j.ejpe.2015.05.008.

21. McClory, C., McNally, T., Brennan, G. P., & Erskine, J. (2007).

Thermosetting polyurethane multiwalled carbon nanotube composites. Journal of Applied Polymer Science, 105(3), 1003-1011. http://dx.doi.org/10.1002/app.26144.

22. Rueda-Larraz, L., d’Arlas, B. F., Tercjak, A., Ribes, A., Mondragon, I., & Eceiza, A. (2009). Synthesis and microstructure–mechanical property relationships of segmented polyurethanes based on a PCL–PTHF–PCL block copolymer as soft segment. European Polímeros, 29(1), e2019012, 2019

Evaluation of the dispersion of carbon nanotubes in an elastomeric polyurethane and fatigue test Polymer Journal, 45(7), 2096-2109. http://dx.doi.org/10.1016/j. eurpolymj.2009.03.013.

23. Jimenez, G., Asai, S., Shishido, A., & Sumita, M. (2000).

Effect of the soft segment on the fatigue behavior of segmented polyurethanes. European Polymer Journal, 36(9), 2039-2050. http://dx.doi.org/10.1016/S0014-3057(99)00241-4.

24. Pichon, P. G., David, L., Méchin, F., & Sautereau, H. (2010).

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25. Nikulin, S. A., Markelov, V. A., Gusev, A. Y., Nechaykina, T. A., Rozhnov, A. B., Rogachev, S. O., & Zadorozhnyy, M. Y. (2013). Low-cycle fatigue tests of zirconium alloys using a dynamic mechanical analyzer. International Journal of Fatigue, 48, 187-191. http://dx.doi.org/10.1016/j.ijfatigue.2012.10.019.

26. Aindow, T. T., & O’Neill, J. (2011). Use of mechanical tests to predict durability of polymer fuel cell membranes under

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humidity cycling. Journal of Power Sources, 196(8), 38513854. http://dx.doi.org/10.1016/j.jpowsour.2010.12.031.

27. Rublon, P., & Favier, A. (2015). Effect of antioxidants on the fatigue crack growth behavior of filled SBR compounds.

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29. Hollertz, R., Chatterjee, S., Gutmann, H., Geiger, T., Nuesch, F. A., & Chu, B. T. T. (2011). Improvement of toughness and electrical properties of epoxy composites with carbon nanotubes prepared by industrially relevant processes. Nanotechnology, 22(12), 125702. http://dx.doi.org/10.1088/0957-4484/22/12/125702.

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Received: Mar. 29, 2018 Revised: Oct. 18, 2018 Accepted: Oct. 29, 2018

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ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.01717

Effects of mercerization in the chemical and morphological properties of amazon piassava Viviane Rebelo1, Yuri da Silva1, Saulo Ferreira2, Romildo Toledo Filho2 and Virginia Giacon1*  Programa de Pós-graduação em Engenharia Civil, Faculdade de Tecnologia – FT, Universidade Federal do Amazonas – UFAM, Manaus, AM, Brasil 2 Departamento de Engenharia Civil, Instituto Alberto Luiz Coimbra de Pós-graduação e Pesquisa de Engenharia – COPPE, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brasil 1

*giacon@ufam.edu.br

Abstract The objective of this work was to investigate the effects of mercerization on chemical, morphological and thermal properties of Amazon Piassava Fibers. The effect of this treatment was studied using XRF, SEM, XRD and TGA. The fibers have been treated in 5% and 10% NaOH for 60 min. The XRF results for treated and untreated fibers showed that there is a decrease in the amount of SiO2 by increasing the NaOH concentration. It has been possible to observe through SEM in untreated fiber that the surface presents a well arranged pattern of silicon rich star-like protrusions.

For the two concentrations, SEM allowed to notice that the removal of deleterious surface impurities and fiber roughness was enhanced. The removal of organic material after treatment can be observed in the TGA analysis. XRD analysis indicate an increase in the crystallinity index, 0.19 to 0.31 after the treatment for 10% concentration solutions.

Keywords: alkaline treatment, mercerization, piassava fibers, superficial modification.

How to cite: Rebelo, V., Silva, Y., Ferreira, S., Toledo Filho, R., & Giacon, V. (2019). Effects of mercerization in the chemical and morphological properties of amazon piassava. Polímeros: Ciência e Tenologia, 29(1), e2019013. https:// doi.org/10.1590/0104-1428.01717

1. Introduction The increasing global demand for consumer goods has exerted a rising pressure on the Earth’s resource consumption.

This has led to increased interest in the development of sustainable materials[1].

Among such renewable resources are the lignocellulosic materials or vegetable fibers. Brazil, specifically the Amazon region, is rich in renewable resources and possesses wide variety of vegetable fibers[2].

As a result, the use of natural fibers as a viable reinforcement in composite materials is increasing and has gained a significant research interest, possibly due to features like high specific modulus, non-toxicity, biodegradability, low cost and abundance[3]. Natural fibers like flax, hemp, jute and sisal have been well known as potential adequate reinforcements for engineering fiber composites[4-8].

Nevertheless, the Piassava Fiber (Leopoldinia Piassaba) in despite of their availability in Amazon, has been used mainly as raw material in craft and as home utensils as brooms or brushes and there is a lack of specific studies focused on Piassava Amazon fiber uses and their properties in composites materials.

On the other hand, the issue of fiber-matrix interaction in composites materials has received increasing attention.

However, natural fibers have structural compounds (cellulose, hemicelluloses, lignin and other substances) that allow moisture absorption from the environment. The interaction between the hydrophilic fibers and their hydrophobic matrix

Polímeros, 29(1), e2019013, 2019

causes fiber swelling within the matrix[9]. This results in the weakening of bonding strength at the interface, therefore leading to dimensional instability, matrix cracking and poor mechanical properties of the composites[10,11].

In order to enhance the effectiveness of interfacial bonding between fiber and matrix, fiber surface needs to be modified. There are many types of treatments, such as, for instance, sodium hydroxide (NaOH) or mercerization, which is being widely used to modify the cellulosic molecular structure[5,12]. According to John, 2007, This modification promotes the access to penetrate chemicals causing reaction with water molecules and move out them from the fibre structure, while the remaining reactive water molecules form fibre-cell-O-Na bonds[11]. This process increase the fibres moisture resistance property due to reduction of hydropylic hydrowyl groups[10,13]. Besides, another important modification done by mercerization is the increase in the surface roughness caused by the disruption of hydrogen bonding in the network structure. Consequently, there is improvement in the composite mechanical properties[6,14].

The main objective of the present study consists of clarifying the effect of mercerization on chemical, morphological and thermal properties of the Amazon Piassava Fibers (Leopoldinia piassaba), filling the gap of information about the surface modifications after the treatment, which could improve the interface adhesion between Amazon Piassaba Fiber and polymer matrix

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

Rebelo, V., Silva, Y., Ferreira, S., Toledo Filho, R., & Giacon, V.

2. Materials and Methods 2.1 Materials The Piassava fibers used in the present investigation were purchased from O. A. NunesNeta - ME (Brazilian company), located in Manaus.The fibers presented a moisture content of 13%, measures in laboratory. The NaOH was obtained

by Nuclear Company, supplier Instrumental Technical Ltda, Manuas-Am.The fibers were cut at a nominal length of 1 cm.

2.2 Methods 2.2.1 Mercerization For mercerization treatment, a commercial NaOH solution was used.

The fibers were cut at a nominal length of 1 cm and soaked in two different NaOH concentrations (5% and 10% by weight) for 60 min. Then, the fibers were rinsed several times with water to remove the excess NaOH solution, until achieving pH 7, and afterwards dried at room temperature conditions of 36 °C for 24hrs to remove excess water.

The effects of the mercerization in the fibers were investigated using the following techniques: XRF, SEM, XRD and TGA. 2.2.2 X Ray Fluorescence (XRF) The measurements of the sample elemental composition were performed through the X-ray fluorescence (XRF) technique using an Epsilon3-XL spectrometer (PANalytical) with 15 W maximum power. A semiquantitative analysis of the spectra was performed with the Omnian package.

The piassava fiber’s microstructure was investigated using an SEM Hitachi TM3000. The microscope was operated under an accelerating voltage of 15 kV. A pre-coating with a thin layer of approximately 20 nm of gold was done to make the fiber conductive and suitable for analysis.

2.2.4 X Ray Diffraction (XRD) X-ray diffraction studies were performed under ambient condition in the equipment EMPYREAM PIXcel 3D (Panalytical), using Cu Ka radiation (1.5406Å), operating at 40 Kv and current of 40Ma. The scan was taken at a range of 5° < 2θ < 45° at a rate of 1°/min and step size 0.0001°.

Crystallinity index (Ic) and percentage crystallinity (% Cr) were calculated using Equations 1 and 2[15,16],:

Icr =

I 22 x100 I 22 + I18

I 22 − I18

I 22

2.2.5 Thermo Gravimetric Analysis (TGA) Thermal analysis of fibers (TGA/DTGA) was performed in natural and treated fibers. The tests were performed in a SDT Q600 simultaneous TGA/DTA/DSC from TA Instruments.

Samples weighing 10 mg were submitted to a heating rate of 10 °C/min until reaching 600 °C in a platinum crucible using 100 ml/min of nitrogen as the purge gas.

3. Results and Discussions The XRF results are shown in Table 1. After the mercerization process, there is a decrease of the SiO2 content from 34.9% of the natural fiber to 9.6% and 6.6% for 5 and 10% concentration solutions, respectively. The treatment processes can enhance removal of surface impurities and increase fiber roughness. This is advantageous for fiber to matrix adhesion as it facilitates mechanical interlocking from the increased surface area available for contact with the matrix and improves the composites mechanical properties[19].

Moreover, Table 1 shows that after the treatment there is presence of sodium in the fiber. This is possible since the treatment changes the fiber structure and there is the reaction of sodium hydroxide with cellulose. This reaction is shown in Equation 3[3].

Cellulose − OH + NaOH → Cellulose − O − Na + + H 2O + Surface Impurities

2.2.3 Scanning Electron Microscopy (SEM)

Cr % =

The method was for empirical measurements to allow rapid comparison of samples, where I22 and I18 are the crystalline and amorphous intensities at 2θ scale to 22° and 18°, respectively[17,18].

(3)

Figure 1 shows SEM micrographs of the fiber surface morphology and the protrusions, which are Si rich particles[20,21].

The surface topography of piassava (Fib 2b) is rougher than before treatment (Figure 2a) and the surface looks cleaner. This occurs because the mercerization removes lignin, pectin, waxy substances, and natural oils that cover the external surface of the fiber cell wall, revealing the fibrils,

(1)

(2)

Figure 1. SEM of piassava fiber untreated: protrusions.

Table 1. Chemical elements presentin the untreated and treated Piassava fiber samples obtained by the X-ray fluorescence.

Element Na Mg Al Si

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Fiber in Natura (w.%) 0.000 0.541 21.109 34.965

5% NaOH (w.%) 8.435 0.026 6.491 9.664

10% NaOH (w.%) 22.842 0.169 8.194 6.668

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Effects of mercerization in the chemical and morphological properties of amazon piassava resulting in a rough surface topography of the fiber[3,21]. Then, with the higher solution concentration there is a roughness increase (Figure 3a and 3b).

Figure 4a shows the transversal cross-section of the untreated piassava sample. After treatment (Figure 4b), the cell wall thicknessvaried from 3 μmto 1 μm, which can be attributed to the removal of organic materials, suchas lignin and other substances[3].

The results of the thermogravimetric analysis are presented in Figure 5. It is possible to observe in the untreated fiber decomposition behavior a small low temperature weight loss (11.88%), between 25 and 112 °C, that can be attributed to water in the form of absorbed moisture or combined water[22]. The onset of the fiber thermal degradation began at about 180 °C, with the decomposition of hemicellulose.

According to the literature[11], the TGA curve of hemicellulose presented three stages of decomposition (from 25 to 180 °C, from 180 to 280 °C and 280 to 500 °C). The peak at 361 °C could be associated with the decomposition of cellulose.

This value is higher than the reported for the pure cellulose sample[11] and the piassava (Attaleafunifera)[10]. The thermal decomposition of lignin ranged from 150 to 450 °C[11] and is dependent on specie. A clear peak for lignin was not observed in this curve.

From Figure 5 it is also possible to observe the TGA and DTGA curves for the mercerized fibers. The mercerized (5-10% NaOH) fibers showed lower decomposition temperatures compared to the untreated fiber. For the main fiber decomposition region (200-400 °C), the treated fibers had a higher weight loss than the untreated fiber, indicating lower thermal stability.This is probably due to the removal of organic materials such as lignin,pectin, waxy substances, and natural oils, evident in the curves of DTGA for mercerized fibers, where the peak around 250 °C was not visible, indicating the removal of hemicellulose from the fiber[3,23].

To study the effectiveness of chemical treatment in piassaba fibers, XRD analysis was performed whose result is shown in the diffractogramin Figure 6.The figure shows

Figure 2. SEM of piassava fiber (a) untreated; (b) 5% NaOH treated.

Figure 3. SEM of piassava fiber treated 10% NaOH, view of fibrils (a) x250; (b) x800.

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Rebelo, V., Silva, Y., Ferreira, S., Toledo Filho, R., & Giacon, V.

Figure 4. SEM of transversal cross-section piassava fiber (a) untreated; (b) 10% NaOH.

Figure 5. TGA and DTGA curves for treated and untreated fiber.

Figure 6. XRD curves of piassava fibers treated and untreated.

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Effects of mercerization in the chemical and morphological properties of amazon piassava Table 2. XRD data of raw and chemical treated piassava fibers.

Sample

I22 (at 2θ scale)

I18 (at 2θ scale)

C.I.

%Cr.

Untreated 5% NaOH 10% NaOH

3093.94 3982.73 3835.82

2499.56 2790.97 2615.45

0.19 0.30 0.31

55.31 58.79 59.45

that piassava fiber (untreated and treated) there are three broad diffraction peaks at 2 θ, the most intense being the 22° corresponding to the crystalline part of the fiber, and halos at 16° and 35º referring to the amorphous part of the fiber, such as the hemicellulose and lignin present in the fibers microfibrils.The result obtained is very similar to those of other lignocellulosic fibers, and reflects the crystalline lattice of cellulose I [24-28].

Table 2 shows the results of percentage crystallinity (% Cr) and crystallinity index (C.I) of untreated and treated piassava fiber. The untreated fiber at 2θ scale gave peaks at 22.0 and 18.0 with relative intensity is 3093.94 and 2499.56.

Percentage crystallinity (%Cr) and crystallinity index (C.I) of untreated piassava fiber are 55.31 and 0.19 respectively whereas percentage crystallinity of alkali treated fibers (5% and 10%) were 58.79 and 59.41. Whereas crystallinity index of alkali treated fibers are 0.30 and 0.31 respectively.

The increase of crystallinity index in alkaline treated piassava fibers indicated that the chemical treatment induced the crystallinity and it increase due to the removal of amorphous materials like hemicellulose, lignin, and some other[23,29-31].

Through alkalization, an increase of composite quality is to be expected due to the improved fibre–matrix adheshion[5].

4. Conclusions The effects of mercerization on the amazon piassava (Leopoldinia piassaba) were presented. The results of the XRF showed that there is a higher sodium concentration when the fiber is treated with the 10% NaOH solution. The fibers present silicon rich star-like protrusions that are removed after the mercerization process. XRD revealed increase in the crystallinity index due to the removal of amorphous materials like hemicellulose, lignin, and some other.

The fiber surface appears cleaner and rougher after the treatment. Thermal analysis of treated fibers displays lower thermal stability compared to the untreated fibers.

These surface modifications could improve the interface adhesion between fiber and matrix, resulting in composites with enhanced mechanical properties.

5. Acknowledgements The authors thank the Brazilian Agency Coordenação de Aperfeiçoamento de Pessoal de Nível Superior – CAPES for the support.

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10. Li, X., Tabil, L. G., & Panigrahi, S. (2007). Chemical treatments of natural fiber for use in natural fiber- reinforced composites: a review. Journal of Polymers and the Environment, 15(1), 25-33. http://dx.doi.org/10.1007/s10924-006-0042-3.

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12. Mohan, T. P., & Kanny, K. (2012). Chemical treatment of sisal fiber using alkali and clay method. Applied Science and Manufacturing, 43(11), 1989-1998. http://dx.doi.org/10.1016/j. compositesa.2012.07.012.

13. John, M. J., & Anandjiwala, R. D. (2007). Recent developments in chemical modification and characterization of natural fiberreinforced composites. Polymer Composites, 29(2), 187-207. http://dx.doi.org/10.1002/pc.20461.

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15. Segal, L., Creely, J. J. Jr, Martin, A. E. Jr, & Conrad, C. M. (1959). An empirical method for estimating the degree of 5/6

Rebelo, V., Silva, Y., Ferreira, S., Toledo Filho, R., & Giacon, V. crystallinity of nature cellulose using the X-ray diffractometer.

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17. Park, S., Baker, J. O., Himmel, M. E., Parilla, A., & Johnson, D. K. (2010). Cellulose crystallinity index: measurement techniques and their impact on interpreting cellulose performance. Biotechnology for Biofuels, 3(1), 10. http:// dx.doi.org/10.1186/1754-6834-3-10. PMid:20497524.

18. Kalia, S., & Sheoran, R. (2011). Modification of ramie fibers using microwave-assisted grafting and cellulase enzyme-assisted biopolishing: a comparative study of morphology, thermal stability, and crystallinity. Journal of Polymer Analysis and Characterization, 16(5), 307-318. http://dx.doi.org/10.1080/ 1023666X.2011.587946.

19. Mokaloba, N., & Batane, R. (2014). The effects of mercerization and acetylation treatments on the properties of sisal fiber and its interfacial adhesion characteristics on polypropylene.

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Part A, Applied Science and Manufacturing, 43(3), 353-362. http://dx.doi.org/10.1016/j.compositesa.2011.12.004.

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23. Zimmermann, E. G. M., Turell, T. C., Zattera, A. J., & Santana, R. M. C. (2014). Influence of chemical treatment of banana fiber composite of poly (ethylene-co-vinyl acetate) with and without blowing agent. Journal of Polymer Science and Technology, 24(1), 58-64. http://dx.doi.org/10.4322/polimeros.2013.071.

24. Rong, M. Z., Zhang, M. Q., Liu, Y., Yang, G. C., & Zeng, H. M. (2001). The effect of fiber treatment on the mechanical

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Crystal transition from cellulose I to cellulose II in NaOH treated Agave americana L. fibre. Carbohydrate Polymers, 86(3), 1221-1229. http://dx.doi.org/10.1016/j.carbpol.2011.06.037.

26. Goswami, P., Blackburn, R. S., El-Dessouky, H. M., Taylor, J., & White, P. (2009). Effect of sodium hydroxide pre-treatment on the optical and structural properties of lyocell. European Polymer Journal, 45(2), 455-465. http://dx.doi.org/10.1016/j. eurpolymj.2008.10.030.

27. Široký, J., Manian, A. P., Široká, B., Abu-Rous, M., Schlangen, J., Blackburn, R. S., & Bechtold, T. (2009). Alkali treatments of lyocell in continuous processes. I. Effects of temperature and alkali concentration on the treatments of plain woven fabrics.

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Grafting of polyacrylonitrile onto guar gum under microwave irradiation. Journal of Applied Polymer Science, 92(3), 15691575. http://dx.doi.org/10.1002/app.20099.

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30. Gonçalves, A. P. B., Miranda, C. S., Guimarães, D. H., Oliveira, J. C., Cruz, A. M. F., Silva, F. L. B. M., Luporini, S., & José, N. M. (2015). Physicochemical, mechanical and morphologic characterization of purple banana fibers. Materials Research, 18(2, Suppl 2), 205-209. http://dx.doi.org/10.1590/15161439.366414.

31. Geethamma, V. G., Joseph, R., & Thomas, S. (1995). Short coir fibre reinforced natural rubber composites: effects of fibre length, orientation and alkali treatmant. Journal of Applied Polymer Science, 55(4), 583-594. http://dx.doi.org/10.1002/ app.1995.070550405.

Received: Jan. 23, 2018 Revised: Aug. 14, 2018 Accepted: Dec. 28, 2018

Polímeros, 29(1), e2019013, 2019

ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.04018

Gamma irradiation effects on polycaprolactone/zinc oxide nanocomposite films Marcos Paula1, Ivo Diego2, Ronaldo Dionisio3, Glória Vinhas2 and Severino Alves1*  Laboratório de Terras Raras, Departamento de Química Fundamental, Universidade Federal de Pernambuco – UFPE, Recife, PE, Brasil 2 Laboratório de Materiais Poliméricos e Caracterização, Departamento de Engenharia Química, Universidade Federal de Pernambuco – UFPE, Recife, PE, Brasil 3 Instituto Federal de Pernambuco – IFPE, Vitória de Santo Antão, PE, Brasil 1

*salvesjr@ufpe.br

Abstract Polycaprolactone (PCL) to which has been added zinc oxide nanoparticles (ZnO NPs) produces nanocomposites (PCL/ZnO NCs). These nanocomposites can be used in biomedical applications and in the food packaging sector.

However, for these materials to be used in these applications, they need to be sterilized. For this, gamma irradiation is the most common method. Thus it is important to evaluate the effects of gamma irradiation on the properties of PCL and PCL/ZnO that have been exposed to gamma irradiation. PCL/ZnO NCs films were obtained by solvent casting and exposed to gamma irradiation at 25 kGy and evaluated by Fourier transform infrared spectra (FT-IR), X-ray diffraction (XRD), thermogravimetric analysis (TGA), differential scanning calorimetry (DSC) scanning electron microscopy (SEM) and mechanical properties. Mechanical properties and crystallinity showed marginal variations for the irradiated samples. The results obtained demonstrate that gamma irradiation at 25 kGy, did not cause profound changes in nanocomposite properties.

Keywords: polycaprolactone, gamma irradiation, nanocomposites, materials properties.

How to cite: Paula, M., Diego, I., Dionisio, R., Vinhas, G., & Alves, S. (2019). Gamma irradiation effects on polycaprolactone/zinc oxide nanocomposite films. Polímeros: Ciência e Tecnologia, 29(1), e2019014. https://doi. org/10.1590/0104-1428.04018

1. Introduction Conventional synthetic polymers are widely used in our society due to characteristics such as low density, low cost of production, and facility for molding, among others.

Due to their chemical constitution, however, they have a low rate of degradation, leading to environmental problems[1-3].

An alternative to these materials are the polymeric materials constituted by poly (α-hydroxy acids), which are biodegraded by microorganisms in CO2, water, and cellular components among other materials[4,5]. One of the main representatives of this class of polymers is polycaprolactone, which is a synthetic thermoplastic polymer. PCL is a hydrophobic, semi‑crystalline polymer that can attain a degree of crystallinity of 69%[6,7]. It is commercially available, biocompatible, achieves crystallization and a has a melting point at low temperatures, is soluble at room temperature in solvents such as chloroform, dichloromethane, carbon tetrachloride, benzene, toluene, cyclohexanone and 2-nitropropane, which facilitates processing by solvent casting methods[6,8]. The principal way to obtain PCL mentioned in the literature occurs through the ring-opening polymerization of the ε-caprolactone monomer, which takes place in an inert atmosphere with heat supply and in the presence of a catalyst, resulting in a high molar mass polymer with improved mechanical properties[6,7].

The main property of PCL is its biocompatibility, which

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makes the material very interesting for applications such as for implants[9], controlled drug delivery systems[10], scaffolds in tissue engineering[11], and food-packaging[8]. However, PCL has low thermo-mechanical properties and high gas permeability[1]. The strategy adopted to control this is to add inorganic nanomaterials to the polymer[12], obtaining a nanocomposite and add also other polymers[13,14]. Among the nanomaterials that can be mixed with PCL, the nanoparticles of ZnO can be used[15]. ZnO is a biocompatible inorganic agent, semiconductor with a band-gap in the UV region, which makes it an efficient absorber of UV radiation[16].

When added to polymeric matrices, the nanocomposites (NCs) exhibit antibacterial activity, low toxicity, lower the gas permeability, gener- reactive oxygen species, cell adhesion and proliferation[12,17]. The particles with nanometric dimensions have a strong tendency to agglomerate due to high surface energy, which acts to prevent their dispersion in the polymer matrix, compromising the properties expected for the material[18]. This dispersion can be controlled by modifying the surface of nanoparticles (NPs)[19]. Neverthelles Elen and co-workers using a melt-blending technique reported a proper dispersion of ZnO NPs in the PCL matrix, without the use of modifiers for ZnO NPs[8]. Given the abovementioned characteristics, the nanocomposites produced

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

Paula, M., Diego, I., Dionisio, R., Vinhas, G., & Alves, S. from PCL with ZnO NPs can be employed in biomedical applications and in the food packaging sector, but need to be sterilized in order to be used[20,21]. The most common method of sterilization occurs by exposing these materials to gamma irradiation at the dose of 25 kGy[21]. Exposure of polymer materials to gamma irradiation can lead to changes in the molecular structure, altering both their physical and chemical properties[22,23]. The gamma rays produce excited and ionized species that lead to changes in the properties of the product, having as the main effect the scission and crosslinking of the polymer chains[24]. These effects depend on the type of radiation, polymer structure, applied dose, dose rate and irradiation conditions[23]. The scission of chains causes a reduction in the molar mass of the polymer[25], while cross-linking produces an increase in the molecular mass, increasing its mechanical properties and decreasing its solubility in organic solvents[26], which may not be desirable for applications in food packaging and biomedical use. It is important to evaluate the effects of gamma irradiation on the physical and chemical properties of polycaprolactone and its nanocomposites with ZnO NPs irradiated at 25 kGy. However, no studies have been reported on the effects of gamma irradiation on PCL/ZnO NCs. The objective of this paper is to report on the effects of gamma irradiation at 25 kGy on PCL and PCL/ZnO NCs films, prepared with different concentrations of ZnO by the solvent casting method. The films obtained were exposed to gamma irradiation at 25 kGy in the presence of air, at room temperature and evaluated by experimental techniques such as Fourier transform infrared spectra (FT-IR), X-ray diffraction (XRD), thermogravimetric analysis (TGA), differential scanning calorimetry (DSC) and scanning electron microscopy (SEM).

2. Materials and Methods 2.1 Materials All the reagents used were analytical grade and used as received. Nanosized ZnO powder was acquired from Aldrich with particle size <100 nm. PCL was obtained from CapaTM (PCL 6500). Chloroform was purchased from Dinâmica.

2.2 Preparation of PCL and PCL/ZnO NCs films PCL and PCL/ZnO NCs films in different concentrations (0.5; 1; 2 and 5%) were obtained by solvent casting.

Pre‑determined amounts of ZnO NPs were added to 5 mL of Chloroform and sonicated for 30 minutes. Then, a solution of 2.5g of PCL in 50 mL Chloroform was added. The mixture was stirred for 48 hours. After the stirring period, the mixture was poured into a Petri dish, and the solvent was removed by slow evaporation in air at room temperature.

range of 4000-650 cm-1. Sixteen scans were performed with a resolution of 4 cm−1 for each spectrum.

2.5 X-ray diffraction Diffractograms were recorded at room temperature in the range of 5° to 80° in a Bruker D8 Advanced X-ray diffractometer with Cu Kα (0.15 nm), at the speed of 0.02°/min.

2.6 Thermogravimetric analysis The thermogravimetric analyses were performed in a SHIMADZU DTG-60H instrument, in the range from room temperature to 600 °C, using a platinum crucible with about 10 mg of sample, under a dynamic nitrogen atmosphere (100 mL/min) and with a heating rate of 10°C/min.

2.7 Differential scanning calorimetry Heat flux curves were obtained in a differential scanning calorimeter, model 1 Star* system (Mettler Toledo) under nitrogen atmosphere with the following programming: 1) heating from 0 °C to 80 °C, at a rate of 10 °C/min; 2) cooling to 0 °C, at a rate of 20 °C/min; and 3) heating from 0 °C to 80 °C, at a rate of 10 °C /min. The crystallization percentage Xc (%) of the materials was calculated based on the following equation: , where , represents the heat needed for a melting point for 100% crystalline PCL. The value used for the heat of fusion of the fully crystalline PCL was 139.3 J/g[27].

2.8 Scanning Electron Microscopy (SEM) The samples were prepared on carbon tape on an aluminum support and coated with a 10-20 nm gold film, using a Bal-Tec SCD 050 sputter coater. Images were recorded by a scanning electron microscope (Tescan Mira3) operating at a voltage of 10 KV.

2.9 Transmission Electronic Microscopy (TEM) TEM image for the NC film were obtained in a transmission electron microscope (Jeol, model JEM-2100), with a voltage of 200 kV. Drops of the films of NC suspended in dichloromethane were deposited on carbon grids, with subsequent evaporation of the solvent.

2.10 Mechanical measurements Tensile properties were determined following the ASTM D-882 with an Instron machine EMIC, DL-500 N, crosshead speed of 5 mm/min, at room temperature. For the mechanical test, three samples of each composition were analyzed.

2.3 Irradiation of samples

3. Results and Discussion

PCL and PCL/ZnO NCs films were exposed to gamma irradiation from a source of 60Cobalt (rate of 2,157 kGy/h), at a dose of 25 kGy in the presence of air at room temperature.

3.1 Preparation of PCL and PCL/ZnO nanocomposite films

2.4 Fourier transform infrared spectroscopy The vibrational spectra in the infrared region (IR) were obtained in a Frontier-Perkin Elmer (with attenuated total reflectance universal accessory-UATR) spectrometer in the 2/7

PCL/ZnO NCs films have a white color with an average thickness of 0.15 mm before and after exposure to gamma radiation (Figure 1). The characterization of the thermal and physical properties for the films of PCL and PCL/ZnO NCs, before and after irradiation at 25 KGy, will be discussed in the following sections.

Polímeros, 29(1), e2019014, 2019

Gamma irradiation effects on polycaprolactone/zinc oxide nanocomposite films

Figure 1. Films of PCL and PCL/ZnO NCs before and after exposure to gamma radiation.

3.2 FT-IR analysis of NCs FT-IR spectroscopy was used to evaluate the effects of gamma irradiation on PCL and PCL/ZnO NCs films.

Figure 2 shows the spectra for PCL and PCL/ZnO NCs before irradiation. Pure PCL showed a peak at 1725 cm-1, relative to the stretching mode of C=O in the polymer, while the peaks at 2870 and 2850 cm-1 are attributed to the C-H bonds of the polymer[28].

The same spectrum for PCL was observed in PCL irradiated at 25 KGy (Figure 3). FT-IR analysis of PCL/ZnO NCs before and after exposure to gamma irradiation showed the same spectrum for pure polymer, with no shift in the peak positions (Figures 2 and 3, respectively). This contrasts with the work of Augustine and co-workers, who observed a shift in peak position of C=O to a lower energy level as the nanoparticles of ZnO content increased[28]. These findings suggest that gamma irradiation at 25 KGy did not cause structural changes in the polymer chains.

Figure 2. FT-IR spectra of (a) pure PCL; (b) PCL/ZnO 0.5%; (c) PCL/ZnO 1%; (d) PCL/ZnO 2%; (e) PCL/ZnO 5%.

3.3 X-ray diffraction patterns The crystallinity of NCs films before and after exposure to gamma irradiation was analyzed by XRD. The diffractograms for films are shown in Figures 4 and 5.

Pure PCL has three reflection angles near 21.5°, 22° and 23.8°, relative to (110), (111) and (200) planes of the orthorhombic structure of the polymer, respectively[28].

The diffractogram for ZnO NPs shows the characteristic crystalline behavior of ZnO[28]. NCs with up 0.5% of ZnO showed only the semi-crystaliine behavior of the polymer; this may be due to the low concentration of NPs, resulting in the suppression of the ZnO diffraction[12] (Figure 4). With the addition of more than 1% of nanoparticles, however, it was possible to observe the peaks related to ZnO. These peaks were more intense until the content increased to 5% of the nanoparticles.

The diffractograms for the irradiated samples at 25 KGy were very similar to samples before irradiation, indicating that the gamma irradiation at 25 KGy did not create new crystalline symmetries(Figure 5)[20].

3.4 Thermogravimetric analysis The TGA results of PCL and of PCL and PCL/ZnO NCs films before and after exposure to gamma irradiation are shown in Figures 6 and 7. PCL and PCL/ZnO NCs exhibited Polímeros, 29(1), e2019014, 2019

Figure 3. FT-IR spectra of (a) pure PCL-25 kGy; (b) PCL/ZnO 0.5%-25 kGy; (c) PCL/ZnO 1%-25 kGy; (d) PCL/ZnO 2%-25 kGy; (e) PCL/ZnO 5%-25 kGy.

the same thermal degradation profile, showing a mass loss step in the range of 200 to 500 ° C. This event was ascribed to polymer chain decomposition[12,29] (Figure 6). The Thermal stability of the all samples was analyzed by comparing the 10% (T10) and 50% (T50) weight loss temperature of 3/7

Paula, M., Diego, I., Dionisio, R., Vinhas, G., & Alves, S.

Figure 4. XRD pattern of (a) pure PCL; (b) PCL/ZnO 0.5%; (c) PCL/ZnO 1%; (d) PCL/ZnO 2%; (e) PCL/ZnO 5%; (f) ZnO NPs.

Figure 6. Thermogravimetric analysis of (a) pure PCL; (b) PCL/ ZnO 0.5%; (c) PCL/ZnO 1%; (d) PCL/ZnO 2%; (e) PCL/ZnO 5%.

Figure 5. XRD pattern of (a) pure PCL-25 kGy; (b) PCL/ZnO 0.5%-25 kGy; (c) PCL/ZnO 1%-25 kGy; (d) PCL/ZnO 2%-25 kGy; (e) PCL/ZnO 5%-25 kGy.

Figure 7. Thermogravimetric analysis of pure (a) PCL-25 kGy; (b) PCL/ZnO 0.5%-25 kGy; (c) PCL/ZnO 1%-25 kGy; (d) PCL/ZnO 2%-25 kGy; (e) PCL/ZnO 5%-25 kGy.

the films (Table1). In addition, residual mass is reported.

Table 1 shows that T10 and T50 of PCL start at 370 °C and 400 °C, respectively, while all NCs showed lower values of T10 and T50. The addition of ZnO NPs decreased the thermal stability of the PCL, higher loadings of NPs decreased the T10 and T50 of the NCs. This behavior can be attributed to the interaction established between the interface of the ZnO NPs and the polymer, which decreased the interaction between semi-crystalline polymer chains. Also, the ZnO promoted an oxidative catalysis of NCs, which was in agreement with previous studies reported in the literature[15,30].

Figure 7 shows that the NCs irradiated at 25 KGy had the same thermal degradation profile, but with lower T10 and T50, than the NCs films before irradiation. These findings can be described by the creation of reactive centers during radiation exposure, which anticipate the thermal degradation of the polymer[21].

Table 1. Thermal properties of PCL and PCL/ZnO NCs films.

3.5 Differential scanning calorimetry DSC curves were obtained to quantify the crystallinity of PCL and PCL/ZnO NCs films. Table 2 demonstrates the results for Tm (melting temperature), Tc (temperature of crystallization), ∆Hm (enthalpy of melting) and Xc (degree 4/7

Sample

T10 (°C)

T50 (°C)

PCL PCL/ZnO 0.5% PCL/ZnO 1% PCL/ZnO 2% PCL/ZnO 5% PCL-25 kGy PCL/ZnO 0.5%-25 kGy PCL/ZnO 1%-25 kGy PCL/ZnO 2%-25 kGy PCL/ZnO 5%-25 kGy

370 326 331 333 326 376 293 282 274 285

400 389 378 372 358 407 344 326 318 330

Residual mass (%) at 600 °C 1.94 1.90 1.77 2.20 4.86 1.55 1.51 1.46 2.28 5.18

of crystallinity) for all samples. Melting temperature of the polymer increased with the addition of up to 1% of ZnO NPs. However, the addition of more than this caused a decrease in the melting point temperature. The increase in the melting point temperature for NCs with a low NPs concentration may be due to the formation of thick crystals, shifting the melting point to a higher temperature[17]. Nevertheless, the Polímeros, 29(1), e2019014, 2019

Gamma irradiation effects on polycaprolactone/zinc oxide nanocomposite films Table 2. Melting point temperature, crystallization temperature, enthalpy of melting and percentage of crystallinity for PCL and PCL/ZnO NCs films.

Sample

Tm(°C)

Tc (°C)

∆Hm (J/g)

Xc (%)

PCL PCL/ZnO 0.5% PCL/ZnO 1% PCL/ZnO 2% PCL/ZnO 5% PCL-25 KGy PCL/ZnO 0.5%-25 kGy PCL/ZnO 1%-25 kGy PCL/ZnO 2%-25 kGy PCL/ZnO 5%-25 kGy

61.33 63.47 65.46 63.67 61.72 61.91 62.41 62.07 61.82 63.91

20.46 22.58 21.15 23.26 20.73 23.46 22.56 23.10 23.46 21.06

61.14 57.34 55.02 55.40 55.15 44.82 58.08 67.62 56.90 58.19

43.89 41.11 39.50 39.80 39.59 32.18 41.70 48.54 40.85 41.77

addition of more than 1% of NPs produced agglomerates, which hindered the crystallization of polymer, producing crystals with a thinner lamellar, decreasing the melting point temperature, corroborating with the results obtained

by Augustine and co-workers[17]. PCl and PCL/ZnO NCs irradiated at 25 kGy exhibited marginal variations in the melting point temperature. The temperature of crystallization for films before exposure to gamma irradiation increased with the addition of 0.5% and 1% of NPs, revealing that NPs, in low concentration, had a nucleating effect on the PCL17.

Irradiated PCL/ZnO NCs films had a small decrease in the temperature needed for crystallization when compared to the irradiated PCL film. This was attributed to an enhancement in the crystallinity of the PCL being exposed to radiation.

The percentage of crystallinity for PCL (43.89%) is in accordance with the literature[8]. The crystallinity of the PCL/ZnO NCs showed a comparable crystallinity, around 40%. PCL irradiated film decreased in crystallinity by 32.18%.

Neverthelles irradiated NCs films that had almost the same crystallinity, with the exception of PCL/ZnO 1%-25 kGy, which exhibited 48.54% for Xc. These findings suggest that gamma irradiation at 25 kGy, did not cause changes in the temperature of the melting point, crystallization temperature or percentage of crystallinity in all the samples analyzed.

3.6 Scanning electron microscopy The distribution of ZnO NPs in the polymer matrix was evaluated by scanning electron microscopy. Figure 8 shows the surface morphology of the PCL/ZnO as 5% before exposure to irradiation. Figure 8 demonstrates the presence of some nanoparticle aggregates dispersed randomly in the polymer matrix. NPs can form aggregates due to their high surface energy and the strong interactions between particles, which hinders their dispersion in the polymer matrix[18].

3.7 Transmission electron microscopy Figure 9 shows TEM image of the PCL/ZnO at 5% before exposure to radiation. The TEM image reveals a distribution of ZnO NPs aggregates, randomly dispersed in the PCL matrix.

3.8 Mechanical properties The effects of gamma irradiation on the mechanical properties of PCL and PCL/ZnO films were investigated and are summarized in Table 3. The properties analyzed Polímeros, 29(1), e2019014, 2019

Figure 8. SEM image of PCL/ZnO 5%.

Figure 9. TEM image of PCL/ZnO 5%.

Table 3. Mechanical properties obtained for PCL and PCL/ZnO NCs films.

Sample PCL PCL/ZnO 0.5% PCL/ZnO 1% PCL/ZnO 2% PCL/ZnO 5% PCL-25 kGy PCL/ZnO 0.5%-25 kGy PCL/ZnO 1%-25 kGy PCL/ZnO 2%-25 kGy PCL/ZnO 5%-25 kGy

σ (MPa) 6.65±1.23a 10.69±3.33a,b 9.22±1.43a,b 13.15±3.81b 7.84±0.25a 9.97±0.84a,b 8.15±1.10a 8.50±1.58a 7.00±1.77a 6.70±4.13a

ε (MPa) 168.83±16.35b,c,e,f 144.23±22.25a 162.30±7.96a,b,c 173.60±4.38c,e,f 183.33±11.24e,f 165.06±14.51a,b,c 155.13±5.09a,b,c 148.10±11.10a,b 165.66±6.81a,b,c 188.5±9.00f

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

were tensile strength (σ) and elastic modulus (ε). The mean values of the mechanical properties were compared with the Duncan test with a significance level of 5%. The Duncan test verified that the tensile strength increased by 2.0% for the PCL/ZnO and decreased by 5.0% for the PCL/ZnO. 5/7

Paula, M., Diego, I., Dionisio, R., Vinhas, G., & Alves, S. This can be explained by the formation of agglomerates in higher concentrations, preventing stress transfer to NPs, reducing σ28. Nevertheless, the irradiated samples did not present significant statistical variations as compared with pure PCL. There was an increase in ε for PCL/ZnO 2%; PCL/ZnO 5%; and PCL/ZnO, 5%-25 kGy. The highest value for ε indicated that the NCs films were more rigid as a consequence of the high loadings of the NPs.

4. Conclusions PCL/ZnO NCs films were obtained by solvent casting method with different contents of ZnO NPs and irradiated at 25 kGy. SEM images revealed that aggregates of ZnO NPs were dispersed randomly in the polymer matrix. FTIR gamma irradiation at 25 kGy did not cause structural changes in the polymer chains. XRD showed that gamma irradiation did not create new crystalline symmetries in the NCs. The addition of ZnO NPs into the PCL Matrix decreased the thermal stability in irradiated and non-irradiated samples. DSC measurements revealed that the crystallinity of PCL/ZnO NCs films had small variations in all samples before and after exposure to gamma irradiation. The same behavior was verified for mechanical properties. This suggests that gamma irradiation at 25 kGy is a successful method for the sterilization of PCL/ZnO NCs films without changes in the properties of the materials.

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

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Received: Aug. 01, 2018 Revised: Dec. 29, 2018 Accepted: Jan. 12, 2019

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ISSN 1678-5169 (Online)

https://doi.org/10.1590/0104-1428.04418

Synthesis and characterization of isoprene oligomers to compare different production chemical processes Renata Vieira Pires1, Larissa Mota Barros Pessoa1, Monica de Almeida de Sant’Anna1, Alexander Fainleib1,2, Rita de Cassia Pessanha Nunes3 and Elizabete Fernandes Lucas1,3*  1 Laboratório de Macromoléculas e Colóides na Indústria de Petróleo – LMCP, Instituto de Macromoléculas Professora Eloisa Mano – IMA, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brasil 2 Institute of Macromolecular Chemistry – IMC, The National Academy of Sciences of Ukraine – NAS, Kyiv, Ukraine 3 Programa de Pós-graduação em Engenharia Metalúrgica e de Materiais – PEMM, Instituto Alberto Luiz Coimbra de Pós-graduação e Pesquisa de Engenharia – COPPE, Universidade Federal do Rio de Janeiro – UFRJ, Rio de Janeiro, RJ, Brasil

*elucas@metalmat.ufrj.br

Abstract Three methods to obtain isoprene oligomers were evaluated: chemical degradation of non-vulcanized coagulated natural rubber; chemical degradation of natural rubber latex; and oligomerization of the isoprene monomer. The products were characterized by infrared spectrometry (FTIR), nuclear magnetic resonance (NMR) and size-exclusion chromatography (SEC). All the three processes were efficient and can be controlled in order to obtain products with desired molar mass.

Among the degradation processes, the reaction with the non-vulcanized rubber led to the purest products, but this process has the disadvantage of relatively higher catalyst cost of the catalyst. Reactions of isoprene with free radical initiation produced oligomers under specific conditions: low isoprene concentration, low initiator concentration, and xylene as solvent. The results discussed here allows the readers to have a chemistry overview and experimental insights about different chemical routes to obtain isoprene oligomers, compiled together in the same work. It shall be helpful for applied chemistry researches.

Keywords: natural rubber, degradation, synthesis, isoprene oligomers.

How to cite: Pires, R. V., Pessoa, L. M. B., Sant’Anna, M. A., Fainleib, A., Nunes, R. C. P., & Lucas, E. F. (2019).

Synthesis and characterization of isoprene oligomers to compare different production chemical processes. Polímeros: Ciência e Tecnologia, 29(1), e2019015. https://doi.org/10.1590/0104-1428.04418

1. Introduction The use of products from the degradation of natural rubber (cis-polyisoprenes) has been widely studied for various application areas, because natural rubber is a renewable source resource and a potential raw material for various oligomers. These oligomers can be employed in formulations to modify polymers, to obtain polyelectrolytes, as artificial skin, among others[1-3].

Natural rubber, poly(cis-1,4-isoprene), is obtained from various trees that produce rubber latex, such as Hevea Brasiliensis. Latex is a polydisperse colloidal system of rubber particles in aqueous phase. The rubber produced from the latex has molar mass in the range of 50,000 to 3,000,000 (g/mol)[4-6]. The depolymerization and degradation of natural rubber include the following main methods: (i) metathesis reactions with chain transfer agents, a transition metal complex (organometallic catalysts); (ii) oxidative degradation through epoxidation with hydrogen peroxide, periodic acid, m-chloroperbenzoic acid, NaNO2, (NH4)2S2O8, K2S2O8/sodium dodecyl sulfate, microorganisms, iron (II) chloride/phenylhydrazine or iron(III) chloride/chloranil

Polímeros, 29(1), e2019015, 2019

systems, diallyl disulfide etc., and (iii) using sunlight, ultraviolet radiation or ultrasound[7-10]. All studies found in the literature, about chemical reactions and rubber and its derivate structures and properties, are more focused in describing separately natural rubber degradation methods and respective chemistry characterization.

The methods described enable obtaining isoprene oligomers from natural rubber. However, they can also be obtained from polymerization of isoprene. Isoprene is a conjugated diene, so it is able to produce a polymerization reaction through the use of different initiators, such as free radical, anionic, cationic and coordination initiators.

Distinct isomeric structures can be produced, depending on the initiation system used. The obtainment of the oligomer can be facilitated by the use of chain transfer agents such as dodecyl mercaptan[11,12].

To compare the isoprene oligomers obtained with different reagents and processes, in this work these oligomers were obtained by oligomerization of isoprene and chemical degradation of (i) natural rubber latex and

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

Pires, R. V., Pessoa, L. M. B., Sant’Anna, M. A., Fainleib, A., Nunes, R. C. P., & Lucas, E. F. (ii) coagulated non-vulcanized natural rubber. All the products were characterized by Fourier-transform infrared spectrometry (FTIR), nuclear magnetic resonance spectrometry (NMR) and size-exclusion chromatography (SEC). The aim of this work was somehow to facilitate the readers that may have interest in the obtention of isoprene oligomers, by providing some insights and comparison of three different chemical methods described together in the same academic source.

2. Materials and Methods 2.1 Methods to obtain the polyisoprene oligomers The polyisoprene oligomer samples were obtained by three chemical routes: chemical degradation of non-vulcanized coagulated natural rubber; chemical degradation of natural rubber latex; and oligomerization of isoprene.

2.1.1 Chemical degradation of coagulated non-vulcanized natural rubber The degradation of coagulated non-vulcanized natural rubber (NR) consisted of a metathesis reaction, i.e., a catalytically induced reaction in which the carbon‑carbon double bonds undergo cleavage and reformulation processes[13-16]. The reaction was performed with small pieces of NR (Braslatex Ind. Com. de Borrachas Ltda.), a second-generation catalyst based on Ru-alkylidene (second-generation Hoveyda-Grubbs, 97% - Sigma-Aldrich), dissolved in dichloromethane (Vetec Química Fina Ltda.), and β-pinene (Sigma-Aldrich) as transfer agent, at temperatures ranging from 30 to 45 °C, in a nitrogen atmosphere with magnetic stirring[14]. Molar ratios of NR:β-pinene varying from 3:1 to 1:1 and of NR:catalyst (NR:Grubbs) ranging from 50:1 to 1000:1 were used. After 48 hours, the reaction was stopped with ethyl vinyl ether (Sigma-Aldrich) and the final product was left for 72 hours in a vacuum oven at temperature of approximately 30 °C to remove the dichloromethane.

2.1.2 Chemical degradation of the natural rubber latex First, the natural rubber latex (Braslatex Ind. Com. de Borrachas Ltda.) was submitted to an epoxidation reaction followed by procedures described in the literature[17-19].

The latex was kept under magnetic stirring in an open system for about 4 hours at room temperature to eliminate ammonia.

Then, 5% of the stabilizer Ultranex (Oxiteno Nordeste S.A.) (% by mass, in relation to the dry mass of natural rubber) was added and the stirring was maintained for another 30 minutes, followed by 12 hours of rest. The epoxidation reaction was conducted through slowly addition of formic acid (Vetec Química Fina Ltda.) and hydrogen peroxide (Vetec Química Fina Ltda.), at NR:HCOOH or NR:H2O2 molar ratios of 1:0.2 or 1:0.4, at 60 °C. The degradation reaction of the epoxidized NR occurred in a closed system at temperatures of 30, 45 and 60 °C, in the presence of periodic acid (H5IO6) (Vetec Química Fina.) at a NRdry/epoxidized:H5IO6 molar ratio that varied from 1:1 to 1:3.75. After the reaction, the product was dried in a vaccum oven at 30 °C, over a period of 72 hours.

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2.1.3 Oligomerization of isoprene The oligomerization was carried out with isoprene (Lanxess Energizing Chemistry) distilled at 34 °C. The reactions were conducted with two types of initiation: anionic, using butyllithium (Sigma-Aldrich); and free radical, employing benzoyl peroxide (Vetec Química Fina Ltda.). The anionic polymerizations were carried out by the mass technique, using isoprene:butyllithium molar ratios of 1:0.04 and 1:0.16, varying the reaction temperature from 25 to 30 °C. The free radical polymerizations were performed in a toluene or xylene solution with isoprene:benzoyl peroxide molar ratios of 1:0.1, 1:0.3 and 1:0.5. The free radical reactions performed in the presence of the transfer agent utilized dodecyl-mercaptan (Merck), at a temperature of about 80 °C. After the reaction, the product was placed in a vacuum oven at approximately 35 °C to remove the residual monomer, until reaching constant weight (at least 72 hours).

2.2 Characterization of the oligomers The reaction products were characterized by Fouriertransform infrared spectroscopy (FTIR), nuclear magnetic resonance spectroscopy (NMR) and size-exclusion chromatography (SEC).

The FTIR analyses were carried out with a Varian Excalibur 3100 spectrometer, using casting film over a KBr window, with scan from 4000 to 600 cm-1 and 20 acquisitions with resolution of 4 cm-1.

Some reaction products were characterized by carbon nuclear magnetic resonance (C13NMR), with frequency of 75 MHz and sample concentration of approximately 5% in deuterated chloroform (Cambridge Isotope Laboratories, Inc.).

The epoxidation reaction products were characterized by hydrogen nuclear magnetic resonance (H1NMR) at concentration of 2% in deuterated chloroform. The molar percentage of the epoxidized natural rubber was calculated according to the method described in the literature[18], by relating the areas under the peaks referring to the olefinic protons present in the structure of the residual cis-1,4‑polyisoprene (1), at 5.1 ppm, and the protons adjacent to the oxygen of the epoxidized ring (2), at 2.7 ppm (Equation 1 and Figure 1).

Epoxidation = degree

Area2.7 /

( Area2.7 +

Area5.1 ) x 100 (1)

The product samples had the average molar mass measured by size-exclusion chromatography (SEC) using a Viscotek GPC Max VE 2005 chromatograph with a Shodex KF 806M column. The samples were dissolved in chromatographic-grade tetrahydrofuran (Vetec Química Fina Ltda.) at 10 mg/mL and filtered through a Chromafil PTFE membrane with pore size of 0.45 μm.

Figure 1. Polyisoprene expected structure after epoxidation reaction.

Polímeros, 29(1), e2019015, 2019

Synthesis and characterization of isoprene oligomers to compare different production chemical processes presence of polyisoprene (PI) and residual β-pinene. Total similarity can be observed with the natural rubber spectrum, indicating that the chemical degradation process led only to obtaining the structure of polyisoprene, since no chemical shifts were identified characteristic of the presence of the terminal terpene in the PI’s structure, as also found by other researchers for similar reactions[14]. Polyisoprene oligomers without terminal terpene structures can be more suitable for some specific applications. The signals identified from A to J refer to the residual β-pinene not incorporated in the PI’s structure[20], which was confirmed in the spectrum obtained with the ATP technique. The 1H-NMR spectrum presented in Figure 3b also indicates the presence of hydrogens characteristic of polyisoprene and β-pinene.

The molar percentage of β-pinene in relation to polyisoprene was calculated based on the proportion between the areas at 5.10 ppm and 4.60 ppm (doublet), related to the H of the secondary olefinic carbon of the PI and the 2H of the olefinic

3. Results and Discussion 3.1 Degradation products of the non-vulcanized natural rubber The products obtained from the degradation reaction of non-vulcanized natural rubber were identified as follows: R#-NR: reaction product (R) from degradation of coagulated non-vulcanized natural rubber (NR), where # corresponds to the reaction sequence number. The concentrations of the ruthenium-based catalyst and transfer agent, β-pinene, were varied, aiming to reduce the molar mass of the products.

Table 1 identifies the products and the average molar mass results.

Unexpectedly, the average molar mass values obtained for the degradation products of coagulated NR were not coherent with the variation of the reaction conditions, i.e., they were not in line with the NR:catalyst and NR:transfer agent ratios, representing a difficult reaction control.

The FTIR spectra of the coagulated natural rubber degradation products were similar with each other and also similar to the spectrum of the natural rubber before degradation. This result was expected, since all the products were obtained using the same degradation methods, only varying the concentrations of the catalyst and transfer agent.

The FTIR spectrum referring to the product R12‑NR, presented as an example in Figure 2a, indicates the characteristic absorption frequencies of the organic groups present in natural rubber, including the vibration due to the C-H out‑of-plane angular deformation of the polyisoprene with cis isomerism, at approximately 840 cm-1. Only the products R5-NR to R9-NR presented two deformation bands in this region: one at 840 cm-1(A) and the other at 880 cm-1 (B), as shown in Figure 2b. The latter band can be attributed to the presence of residual β-pinene, which also was identified in the spectrum of β-pinene.

Figures 3a and 3b present, respectively, the 13C-NMR and 1H-NMR spectra of the product R5-NR. All the products characterized presented similar spectra, with the predominant

Figure 2. Infrared spectra of the non-vulcanized coagulated NR degradation products: (a) R12-NR and (b) R5-NR.

Table 1. Average molar masses of non-vulcanized coagulated NR degradation products, obtained after 48 h of reaction at 30 °C. Product R2-NR* R3-NR R4-NR R5-NR R6-NR R9-NR** R8-NR** R7-NR** R12-NR R11-NR R10-NR R13-NR R14-NR R15-NR

Ratio NR: β-pinene mol:mol 1:1 2:1 3:1 1:1 2:1 1:1 2:1 3:1 1:1 2:1 3:1 1:1 2:1 1:1

Ratio NR:Grubbs mol:mol 1000:1 1000:1 1000:1 500:1 500:1 200:1 200:1 200:1 200:1 200:1 200:1 100:1 100:1 50:1

Mp

Mn

Mw

(g/mol) 5,990 2,382 3,140 11,942 2,772 1,311 1,052 46,156 6,127 7,980 6,861 12,505 8,577

3,090 1,034 2,179 29,500 7,075 1,341 979 1,324 28,064 2,230 3,094 2,647 9,383 6,326

12,850 6,641 10,149 64,228 13,526 2,886 1,505 1,689 51,265 6,601 8,313 7,072 14,245 9,591

Mw

Mn

4.16 6.42 4.66 2.17 1.35 2.14 1.54 1.28 1.83 2.96 2.69 2.67 1.52 1.52

*Reaction of 72 hours; **Reaction at 45 °C; M = molar mass at the peak of the molar mass distribution curve; M n = number average molar p mass; M w = weight average molar mass.

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Pires, R. V., Pessoa, L. M. B., Sant’Anna, M. A., Fainleib, A., Nunes, R. C. P., & Lucas, E. F. CH2 of the β-pinene, respectively. The molar percentages calculated for the free b-pinene were 30, 9, 9, 7 and 32%, respectively, for samples R5-NR, R6-NR, R7-NR, R11-NR and R12-NR. These values are coherent with the proportion of NR:β-pinene utilized in the reaction.

The more subtle peaks, indicated by the arrow in Figure 3b, suggest the presence of other different compounds of PI and β-pinene, probably the result of side reactions specific to self-metathesis or isomerization of the β-pinene. Another study reported that the terminal olefin metathesis is followed by the methyl groups exchange and that this reaction occurs substantially faster than the metathesis reactions of olefins with double internal bonds[14]. These peaks were identified in the spectra obtained for some other NR degradation products, but with different intensities than those observed in the spectrum of R5-NR.

3.2 Degradation products of the natural rubber latex The NR latex was first epoxidized and the products of this reaction were characterized by FTIR and 1H-NMR. The FTIR spectra revealed the formation of epoxidized isoprene units, by the presence of the bands in the 800-900 cm-1 and 1200-1300 cm-1 regions, which are characteristic of the asymmetric and symmetric axial deformation of the

Figure 3. 13C-NMR (a) and 1H-NMR (b) spectra of R5-NR product, obtained from non-vulcanized coagulated NR degradation.

epoxide ring, respectively. The continuing presence of the band near 1670 cm-1, related to the axial deformation of the C=C bond, representative of the isoprene unit, indicates the epoxidation reaction did not attain 100% yield.

The content of the epoxy groups in the latex, which served as the base for calculating the quantity of the oxidative degradation agent (periodic acid) used in the degradation reactions, was determined by 1H-NMR. Table 2 presents the epoxidation reaction conditions and epoxidation degree values obtained from Equation 1 (Figure 1).

In general, the epoxidized materials had epoxidation levels between 10 and 20% for the same reaction conditions (1:0.2 for NR:HCOOH and NR:H2O2, 8 hours of reaction).

The epoxidation level increased with storage time before using the samples in the degradation reaction. For example, the epoxidation level of reaction #4 increased from 5 to 11% (4E) and for reaction #5 it rose from 12 to 19% (5E). A kinetic study showed that the epoxidation degree increased until a reaction time of 48 yours, after which it remained practically constant. The reactions conducted for equal times (8E, 9E and 10E) or longer (6E) than 48 hours reached levels of about 28% (Table 2).

All the epoxidized products were submitted to the degradation reaction and identified as R#D(#E), where R#D represents the product of the degradation reaction after the epoxidation reaction of the latex (#E). Table 3 reports the degradation reaction of the epoxidized products and average molar masses obtained. The average molar mass results were coherent with the reaction conditions used. Lower molar masses were obtained in the following conditions: (i) higher epoxidation degrees of the samples (e.g., comparing R5D to R9D and R15D to R17D); (ii) greater amount in excess of H5IO6 in relation to the epoxidized NR (e.g., comparing R1D to R3D and R4D, and R5D and R14D to R15D); (iii) longer degradation times (e.g., comparing R3D to R4D, R6D to R7D, and R14D to R11D). The degradation reactions conducted at 60 °C seemed to be less efficient.

Among the products obtained under comparable conditions (R5D, R6D, R7D, R8D and R9D), the temperature appeared to have a stronger influence on molar mass than did the degradation time.

In all the FTIR spectra, the relation between the bands referring to C=C (~1670 cm-1), the epoxide ring (~850 cm-1) and C=O (~1720 cm-1) showed reduced intensities of the oxirane and C=C bands in the degraded product, in respecting to an increase of the C=O band, which confirms the degradation

Table 2. Reaction conditions and respective epoxidation degrees.

Sample # 1E 2E 3E 4E 5E 6E 8E 9E 10E

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NRdry:HCOOH (mol/mol) 1:0.2 1:0.2 1:0.2 1:0.2 1:0.2 1:0.2 1:0.4 1:0.2 1:0.4

NRdry:H2O2 (mol/mol) 1:0.2 1:0.2 1:0.2 1:0.2 1:0.2 1:0.2 1:0.2 1:0.4 1:0.4

Temperature (°C) 60 60 60 60 60 60 60 60 60

Time (h) 8 8 8 8 8 72 48 48 48

Epoxidation degree (%) 14 16 10 11 19 31 29 26 29

Polímeros, 29(1), e2019015, 2019

Synthesis and characterization of isoprene oligomers to compare different production chemical processes reaction. The intensity reduction of the band refered to the epoxide ring depends on the epoxidation condition, resulting in different degradation degrees, as exemplified by FTIR spectra of the degraded samples R1D(1E) and R6D(3E) depicted in Figure 4.

The comparison of R1D(1E) spectrum of Figure 4 with the spectrum of the epoxidized latex 1E shows the continued presence of the epoxide group (~850 and 1250 cm-1), even after the degradation reaction, although with lower intensity, as expected. The appearance of a pronounced band at 1720 cm-1 and the significant increase in the intensity and better definition of the band at 3430 cm-1 indicate the presence of carbonyl groups and hydroxyl, respectively, which were present at the end of the degraded polyisoprene chains.

The various absorption bands present at lower wavelengths (1100 to 1200 cm-1 and 600 to 800 cm-1) can be attributed to the residual periodic acid. The products R6D(3E) and R7D(3E), both from the degradation reaction of the latex with 10% epoxidation conducted at 60 °C, presented practically no difference between the intensities corresponding to the carbonyl and terminal hydroxyl groups in relation to the C=C group of the isoprene unit, and in relation to the characteristic bands of the epoxide ring, disregardless of the concentration of periodic acid utilized. This result suggests low efficiency of the degradation reaction in the conditions analyzed, and is concordant with the much higher molar mass values obtained for these products (Table 3).

The FTIR spectra of the products R5D(2E), R14D(8E) and R16D(9E) (Figure 5), obtained under the same degradation conditions, i.e., NR:H5IO6 ratio of 1:1.25, temperature of 30 °C and 24 hours of reaction, revealed that the ratio between the intensities related to the carbonyl/hydroxyl ratio and the C=C group of the isoprene unit did not follow

Figure 4. FTIR spectra of R1D(1E) and R6D(3E) products.

Figure 5. Infrared spectra of (a) R5D(2E), (b) R14D(8E) and (c) R16D(9E) products.

Table 3. Relative molar mass values of the epoxidized latex degradation products, obtained from 30 (*) and 60 (**) °C reactions.

Sample

Razão NRepoxidada:H5IO6

Tempo (h)

Epoxidação (%)

Mp

R1D(1E)* R2D(1E)* R3D(2E)* R4D(2E)* R5D(2E)* R6D(3E)** R7D(3E)** R8D(4E)* R9D(4E)* R10D(5E)* R11D(6E)* R14D(8E)* R15D(8E)* R16D(9E)* R17D(9E)* R18D(10E)* R19D(10E)** R20D(10E)* R21D(10E)** R22D(10E)*

1:1 1:1.25 1:1.25 1:1.25 1:1.25 1:1.25 1:1.25 1:1.25 1:1.25 1:2.5 1:1.25 1:1.25 1:2.5 1:1.25 1:2.5 1:1.25 1:1.25 1:2.25 1:2.25 1:3.75

12 24 12 24 24 24 12 12 24 12 12 24 24 24 24 48 6 48 6 48

50 14 25Ϯ 25Ϯ 16 10 10 11 11 19 32 28 29 26 26 29 29 29 29 29

2,745 3,068 941 1,969 10,647 15,217 3,179 4,000 2,665 5,339 2,487 3,850 7,266 6,735 6,962 1,760 6,455 14,454 2,478

Ϯ

M n (g/mol) 2,229 2,739 851 2,117 4,840 14,060 1,931 3,558 2,640 3,453 2,189 2,636 4,133 2,565 2,475 3,818 2,183 14,553 2,586

Mw 8.058 11,215 2,682 7,292 24,169 39,059 7,274 4,084 17,493 16,530 8,494 7,573 28,473 13,632 27,695 83,329 25,196 19,504 11,439

Mw

Mn

3.62 4.09 3.15 3.45 4.99 2.78 3.77 1.15 6.63 4.79 3.88 2.88 6.89 5.31 11.19 21.82 11.54 1.34 4.42

M p = molar mass at the peak of molar mass distribution curve; M n = number average molar mass; M w = weight average molar mass;

Ϯ

= estimated value.

Polímeros, 29(1), e2019015, 2019

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Pires, R. V., Pessoa, L. M. B., Sant’Anna, M. A., Fainleib, A., Nunes, R. C. P., & Lucas, E. F. the epoxidation level. In other words, the product R5D(2E), obtained from the latex with 16% epoxidation, presented a higher ratio between the intensities of these groups than did the products R14D(8E) and R16D(9E), obtained from latex with 29% and 26% epoxidation, respectively, suggesting greater efficiency of the degradation reaction with the latex having 16% epoxidation. This was not expected, but it can be related to the larger quantity of formic acid or hydrogen peroxide remaining in the 16% epoxidized latex.

The epoxidation reactions of these materials differed regarding the NRepoxidized:HCOOH and NRepoxidized:H2O2 ratios, which were respectively 1:0.2 and 1:0.2 in reaction 2E, 1:0.4 and 1:0.2 in reaction 8E, and 1:0.2 and 1:0.4 in reaction 9E. In general, the products from degradation with a larger quantity of H5IO6 presented band intensities characteristic of the carbonyl function, at around 1720 cm-1, greater in comparison with the C=C absorption band of the isoprene unit at 1660 cm-1, a finding that is in accordance with the literature[18] and with the molar mass results (Table 3).

This difference was more pronounced in the products of the degradation reactions conducted for longer intervals.

The degradation products of the latex were also characterized by 13C-NMR and 1H-NMR. Figure 6 shows the spectra of R7D(3E), with the same characteristic peaks of PI from the original latex. The low-intensity peaks at 3.65 ppm (A), 2.69 ppm (B) and in the region between 2.11 NS 2.0 ppm (C) in Figure 6a can be attributed to, respectively, hydrogens of the hydroxyl group, hydrogens adjacent to the epoxide ring, and hydrogens of the methyl and methylene groups adjacent to the carbonyl, present in the structure after the degradation reaction. This indicates that the chemical structure of the degraded material is composed of isoprene units, epoxidized units and terminal ketone and aldehyde groups, although there was no evidence of H related to the carbonyl of the aldehyde group with greater chemical shift.

The 13C-NMR spectrum (Figure 6b) also did not indicate the presence of carbonyl, suggesting that the products R6D(3E) and R7D(3E) degraded inefficiently. This is in agreement with the FTIR spectra, which contained a carbonyl band with intensity less than or equal to that of the isoprene band.

Some peaks were present at approximately 1.60 ppm and 1.30 ppm (D) (Figure 6a), which can be attributed to the hydrogens of the secondary, tertiary and primary carbons of the hydrocarbon chain, probably arising from reactive impurities, as also observed in the carbon spectrum (peaks with low intensity in the range from 25 to 35 ppm).

3.3 Products from oligomerization of isoprene The reactions using butyllithium (BuLi) were conducted by the mass technique (R#-M) and the reactions using benzoyl peroxide were conducted in a solution of toluene (R#-T) or xylene (R#-X), seeking to favor reduction of the molar mass due to the chain transfer by the solvent. In some cases, dodecyl mercaptan (C12SH) was used as chain transfer agent.

Table 4 presents the molar masses of the samples obtained under different reaction conditions. In general, higher polydispersion ( M w / M n ) and higher values of M w were observed in comparison with the products from degradation of NR.

Figure 6. 1H-NMR (a) and 13C-NMR (b) spectra of R7D(3E), degraded latex products.

Table 4. Reaction conditions and respective molar masses of the products obtained from isoprene polimerization at 78 °C, after 48 hours.

[Iso] (mol/L)

Sample RP01-T RP02-T RP03-T RP04-X RP05-X RP06-X RP07-T RP08-T RP09-X RP10-X RP11-T RP12-X R8-M (30 °C; 1:40h) R14-M (30 °C; 8h)

6.7 5.7 5.0 6.7 5.7 5.0 1.3 1.3 1.3 1.3 6.7 6.7 (a) (a)

Iso:Per 1:0.1 1:0.3 1:0.5 1:0.1 1:0.3 1:0.5 1:0.5 1:0.1 1:0.5 1:0.1 1:0.1 1:0.1 1:0.16(b) 1:0.16(b)

Iso: C12SH

Mp

Mn

Mw

1:0.025 1:0.025 1:0.025

3,072 2,116 15,488 3,346 2,926 2,701 2,327 2,317 2,636 1,617 3,119 1,889 1,971 2,074

(g/mol) 2,857 3,109 18,512 3,722 2,712 3,430 2,109 2,917 2,550 1,797 2,431 1,396 1,659 1,766

10,179 29,183 90,069 78,369 10,797 25,158 13,490 11,637 8,786 9,865 9,968 8,730 1,867 2,260

Mw

Mn

3.56 9.39 4.86 21.05 3.98 7.33 6.39 3.99 3.45 5.49 4.10 6.25 1.13 1.28

Mass polymerization; Carried out with butil-Li instead peroxide; Iso = isoprene; Per = peroxide; RP#T = carried out in toluene; RP#X = carried out in xylene. (a)

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(b)

Polímeros, 29(1), e2019015, 2019

Synthesis and characterization of isoprene oligomers to compare different production chemical processes The reactions conducted with higher concentrations of isoprene ([Iso]) tended to generate products with larger molar masses. The smallest molar masses were obtained at an isoprene concentration of 1.3 mol/L. The monomer:initiator molar ratio that best promoted obtaining products with low molar masses was 1:0.1. This behavior was expected, because the presence of excess initiator can cause parallel reactions, which was revealed by FTIR for unreacted peroxide. With respect to the influence of the reaction solvents, the expected correlation was observed between molar mass and type of solvent: the lowest average molar masses were obtained from the reaction conducted with xylene, due to the higher chain transfer constant of this solvent in relation to toluene[11].

The FTIR spectra FTIR (Figure 7) confirmed the reaction of isoprene by opening of the double bond, but due to the different reaction conditions, some absorption bands different to the bands of PI were identified. Figures 8a, 8b and 8c show the FTIR spectra obtained, respectively, for the products R14-M, RP01-T and RP06-X, which represent the other products obtained in the same conditions, respectively, of anionic, free radical in toluene and free radical in xylene initiation.

The spectrum of the product R14-M (Figure 7a) was very similar to that of natural rubber, observing as well a band at approximately 840 cm-1 (A), which indicates the possibility that the product possessmolecules with cis isomerism. The absorption band at 3,700 cm-1 (B) can be related to the presence of isopropyl alcohol, utilized to stop the polymerization reaction. All the products obtained from the polymerization of isoprene with benzoyl peroxide presented spectra similar to those of the products RP01-T (Figure 7b) and RP06-X (Figure 7c). Particularly noteworthy was the presence of absorption bands in regions of lower frequency, generally attributed to the C=C and C-H vibrations

of aromatic compounds, indicating the presence of residual solvent, besides the bands related to the chain’s terminal peroxide group. The well-defined band at 1720 cm-1 can be attributed to the carbonyl derived from the peroxide.

In RP06-X, there were two adjacent bands with equal intensity at 1789 cm-1 and 1767 cm-1, which can be related to the presence of residual benzoyl peroxide. The intensity of these bands increased as the peroxide concentration in the reaction increased, both in the reactions conducted in toluene and xylene. In other words, the concentration of benzoyl peroxide rose in the orders RP01-T<RP02-T<RP03-T and RP04-X<RP05-X<RP06-X, as did the intensities of the absorption bands referring to the C=O coupled axial deformations In the 1H-NMR spectra (Figure 8), characteristic peaks of hydrocarbon chains were observed in regions of low chemical shift (~1.0 to 3.0 ppm) (A), as well as two doublets corresponding to the hydrogens of terminal olefinic CH2, adjacent to the tertiary carbon probably present in the isoprene oligomers with low molar mass. These two doublets (B) were more evident in the samples obtained by mass polymerization. The samples obtained by solution polymerization contained residual solvent, revealed by the peaks in the chemical shift region between 7 and 9 ppm, which is characteristic of aromatic hydrocarbons.

A 3-factorial analysis was performed to evaluate the variables relevance and the confidence interval of actual molar mass from the three processes to obtain isoprene oligomers studied. Method (categorical) and, Reaction Temperature and Time (continuous) were considered the factors and M p, M n, M w and M w/ M n (PD) the responses.

The less significant interactions were removed from the effect summary. A log mathematical treatment could also be applied in molar mass data in order to improve the residual by predict plot. In general, the RSquare and RMSE values indicate an acceptable fit to the data and that the model was a good predictor of the response. The low PValues represent a higher confidence interval (95%) of the prediction model.

The greater significant factor to predict the response was Time and Temperature Interactions, while Method and its Interactions are less relevant as showed in Figure 9.

The mean and standard error values suggest that both methods of degradation were similar and provide polyisoprene with low molar mass as proposed. Greater standard error values were observed for the isoprene polymerization method, in comparison to the degradation methods, as showed in Table 5.

Figure 7. FTIR spectra of (a) R14-M, (b) RP01-T and (c) RP06-X products.

Polímeros, 29(1), e2019015, 2019

Figure 8. 1H-NMR spectrum of R10-M product.

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Pires, R. V., Pessoa, L. M. B., Sant’Anna, M. A., Fainleib, A., Nunes, R. C. P., & Lucas, E. F.

Figure 9. Effect summary obtained from JMP software.

Table 5. Standard error and mean of molar masses and polydispersion.

Method Coagulated NR degradation Latex degradation Isoprene olygomerization

Standard error

Mp 2,205 2,007 8,689

Mn

Mw

1,625 1,479 6,403

6,239 5,681 24,585

Mean PD

Mp

Mn

Mw

PD

1.92 1.75 7.55

5,886 5,481 2,396

3,391 3,976 2,394

7,922 14,787 11,826

2.74 4.65 4.70

M p = molar mass at the peak of molar mass distribution curve; M n = number average molar mass; M w = weight average molar mass; PD = polydispersion.

4. Conclusions

5. Acknowledgements

The isoprene oligomerization, degradation of non‑vulcanized coagulated natural rubber and degradation of natural rubber latex processes have satisfactorily produced isoprene oligomers, even that the reactions have been performed with different chemicals and reagents, and isoprene source.

The molar mass can be somehow controlled in each process, for specific purposes. Reactions of isoprene with free radical initiation produced oligomers under specific conditions: low isoprene concentration, low initiator concentration, and xylene as solvent. However, the polydispersity of the products was relatively high. Isoprene oligomers with narrow molar mass distribution were easily obtained by anionic initiation. The mean and standard error values suggest that both methods of degradation (coagulated non-vulcanized and latex degradation reactions) were similar and provide polyisoprene with low molar mass as proposed. Greater standard error values were observed for the isoprene polymerization method, in comparison to the degradation methods. In this case, different experiment conditions were treated as a single factor, that is, the different mechanisms and technics of polymerization used, and also the solvent, are not distinctly evaluated in the categorical factor that represents the isoprene polymerization method. The degradation processes of natural rubber has as an advantage since this product comes from a renewable source, despite the relatively higher cost of the catalyst. However, the reaction with the coagulated non-vulcanized material led to purer products than those obtained from latex. Although the products have a similar main molecular structure, that is, polyisoprene based-chain, the functional end groups and lateral group from the three methods applied will differentiate the final products, providing varied hydrophilicity/hydrophobicity characteristic. This observation should be taken into account for material compatibility intended purpose, to use the isoprene oligomers.

The authors thank CNPq (307193/2016-0), CAPES (PVE #1315110), FAPERJ (E-26/201.233/2014), ANP and Petrobras (0050.0069420.11.9) for financial support.

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6. References 1. Magioli, M., Sirqueira, A. S., & Soares, B. G. (2010). The effect of dynamic vulcanization on the mechanical, dynamic mechanical and fatigue properties of TPV based on polypropylene and ground tire rubber. Polymer Testing, 29(7), 840-848. http:// dx.doi.org/10.1016/j.polymertesting.2010.07.008.

2. Fainleib, A., Grigoryeva, O., Youssef, B., & Saiter, J. M. (2011). Utilization of tire rubber and recycled polyolefins into thermoplastic elastomers. In A. Fainleib, & O. Grigoryeva (Eds.), Recent developments in polymer recycling (pp. 1-46).

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8. Hesham, A. E.-L., Mohamed, N. H., Ismail, M. A., & Shoreit, A. A. M. (2015). Degradation of natural rubber latex by new Streptomyces labedae strain ASU-03 isolated from Egyptian Polímeros, 29(1), e2019015, 2019

Synthesis and characterization of isoprene oligomers to compare different production chemical processes soil. Microbiology, 84(3), 351-358. http://dx.doi.org/10.1134/ S0026261715030078.

9. Polymer Properties Database. (2015). Thermal-oxidative degradation of rubber. Retrieved in 2018, May 30, from http:// polymerdatabase.com/polymer%20chemistry/Thermal%20 Degradation%20Elastomers.html 10. Ibrahim, S., Othman, N., & Ismail, H. (2016). Degradation of natural rubber latex. In J. L. Hamilton (Ed.), Natural rubber: properties, behavior and applications (pp. 105-136). New York: Nova Science Publishers.

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Computational study of metathesis degradation of rubber. distributions of products for the ethenolysis of 1,4-polyisoprene.

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16. Grubbs, R. H. (2007). Realizing the promise of olefin metathesis.

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Received: June 14, 2018 Revised: Aug. 31, 2018 Accepted: Jan. 30, 2019

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