**Fiber Reinforced Polymer Composites: Mechanical Properties and Applications**

Editors

**R.A. Ilyas S.M. Sapuan Emin Bayraktar Shukur Abu Hassan Nabil Hayeemasae Khubab Shaker**

MDPI • Basel • Beijing • Wuhan • Barcelona • Belgrade • Manchester • Tokyo • Cluj • Tianjin

*Editors* R.A. Ilyas Department of Chemical Engineering Universiti Teknologi Malaysia Skudai Malaysia S.M. Sapuan Department of Mechanical and Manufacturing Engineering Universiti Putra Malaysia Serdang Malaysia Emin Bayraktar School of Mechanical and Manufacturing Engineering ISAE-SUPMECA Paris France Shukur Abu Hassan Centre for Advanced Composite Materials (CACM) Universiti Teknologi Malaysia Skudai Malaysia Nabil Hayeemasae Department of Rubber Technology and Polymer Science Prince of Songkla University Pattani Thailand Khubab Shaker Department of Materials National Textile University Faisalabad Pakistan

*Editorial Office* MDPI St. Alban-Anlage 66 4052 Basel, Switzerland

This is a reprint of articles from the Special Issue published online in the open access journal *Polymers* (ISSN 2073-4360) (available at: www.mdpi.com/journal/polymers/special issues/HK8P7953V9).

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## **Contents**



## **About the Editors**

### **R.A. Ilyas**

R.A. Ilyas is a senior lecturer in the Faculty of Chemical and Energy Engineering, Universiti Teknologi Malaysia, Malaysia. He is also a Fellow of the International Association of Advanced Materials (IAAM), Sweden, Fellow of the International Society for Development and Sustainability (ISDS), Japan, a member of Royal Society of Chemistry, UK, and the Institute of Chemical Engineers (IChemE), UK. R.A. Ilyas was the recipient of the MVP Doctor of Philosophy Gold Medal Award UPM in 2019 for Best Ph.D. Thesis and Top Student Award, INTROP, UPM. He was awarded with Outstanding Reviewer by Carbohydrate Polymers, Elsevier, United Kingdom, Top Cited Article 2020-2021 Journal Polymer Composite, Wiley, 2022, and Best Paper Award at various International Conferences. R.A. Ilyas was also listed and awarded Among World's Top 2% Scientist (Subject-Wise) Citation Impact during the Single Calendar Years of 2019, 2020, and 2021 by Stanford University, US, PERINTIS Publication Award 2021 and 2022 by Persatuan Saintis Muslim Malaysia, Emerging Scholar Award by Automotive and Autonomous Systems 2021, Belgium, Young Scientists Network - Academy of Sciences Malaysia (YSN-ASM) 2021, UTM Young Research Award 2021, UTM Publication Award 2021, and UTM Highly Cited Researcher Award 2021. In 2021, he won the Gold Award and Special Award (Kreso Glavac (The Republic of Croatia) at the Malaysia Technology Expo (MTE2022), Gold Award dan Special Award at International Borneo Innovation, Exhibition and Competition 2022 (IBIEC2022), and a Gold Award at New Academia Learning Innovation (NALI2022). His main research interests are (1) polymer engineering (biodegradable polymers, biopolymers, polymer composites, polymer-gels) and (2) material engineering (natural fiber-reinforced polymer composites, biocomposites, cellulose materials, nano-composites). To date, he has authored or co-authored more than 480 publications.

### **S.M. Sapuan**

S.M. Sapuan is Professor (A grade) of composite materials in the Department of Mechanical and Manufacturing, Universiti Putra Malaysia (UPM), Head of Advanced Engineering Materials and Composite Research Centre, UPM, and Chief Executive Editor of Pertanika journals in the office of Deputy Vice Chancellor, UPM. He earned his BEng in Mechanical Engineering, University of Newcastle, Australia, MSc in Engineering Design, Loughborough University, UK, and PhD in Material Engineering, De Montfort University, UK. He is a Professional Engineer and Fellow of Society of Automotive Engineers, Academy of Science Malaysia, International Society for Development and Sustainability, Plastic and Rubber Institute Malaysia, Malaysian Scientific Association, International Biographical Association and Institute of Material Malaysia. He is an Honorary Member of Asian Polymer Association and Founding Chairman and Honorary Member of Society of Sugar Palm Development and Industry, Malaysia. He was World Class Visiting Professor, Andalas University Indonesia. He is Editor-in-Chief of the *Journal of Natural Fibre Polymer Composites*, co-Editor-in-Chief of *Functional Composites and Structures*, Associate Editor-in-Chief of *Defence Technology*, Elsevier, and Editorial Board Member of 30 journals. He has produced over 2000 publications, including over 900 journal papers, 55 books, and 220 chapters in book. He has delivered 60 plenary and keynote lectures, and 150 invited lectures. He has organized 31 journal Special Issues as Guest Editor, reviewed 1500 journal papers, and has 8 patents. He successfully supervised 98 PhD and 75 MSc students and 16 postdoctoral researchers. His h-index is 107 with 41710 citations (Google Scholar).

He received the ISESCO Science Award, Khwarizimi International Award, Kuala Lumpur Royal Rotary Gold Medal Research Award, National Book Award, Endeavour Research Promotion Award, TMU/IEEE India, Citation of Excellence Award, Emerald, UK, and Malaysia's Research Star Award.

#### **Emin Bayraktar**

Emin Bayraktar (Habil., Dr (PhD), DSc- Dr es Science) is an academic and research staff-member at Mechanical and Manufacturing engineering at SUPMECA/Paris, France. His research areas include manufacturing techniques of new materials (basically composites –hybrid), metal forming of thin sheets (design + test + FEM), static and dynamic behaviour and optimisation of materials (experimental and FEM –utilisation and design of composites-based metallic and non-metallic, powder metallurgy, energetic materials aeronautical applications)-based metallic, non-metallic, powder metallurgy and metallurgy of steels, welding, heat treatment, as well as the processing of new composites, sintering techniques, sinter –forging, thixoforming, etc. He has authored more than 200 publications in international journals and international conference proceedings, and has also authored more than 90 research reports (European = Steel Committee projects, Test + Simulation). He has already advised 32 PhDs and 120 MSc Theses, and 7 are ongoing. He is a Fellow of WAMME (World Academy of Science in Materials and Manufacturing Engineering), Editorial Board - Member of JAMME (International Journal of Achievement in Materials and Manufacturing Engineering), Advisory board member of AMPT-2009 (Advanced Materials Processing technologies), APCMP-2008, and APCMP-2010. He was a Visiting Professor at Nanyang Technology University, Singapore in 2012, Xi'an Northwestern Technical University, Aeronautical Engineering, in 2016, University of Campinas, UNICAMP-Brazil in 2013 until 2023. He is recipient of Silesian University Prix pour "FREDERIK STAUB Golden Medal-2009" by Academy of WAMME, "World Academy of Science"-Poland, materials science section, and recipient of the William Johnson International Gold Medal - 2014, AMPT academic association.

#### **Shukur Abu Hassan**

Shukur Abu Hassan, PhD (UTM,Malaysia), MSc (Warwick University, UK), BSc (UTM, Malaysia) is an Engineer, Associate Professor and the current Director of the Centre for Advanced Composite Materials (CACM), Universiti Teknologi Malaysia (UTM), Malaysia. In his teaching career of 33 years, he has worked as the Centre Director, Head of Composite Lab and Committee Member at the Faculty and University level. He is also involved as a Consultant and on the Editorial Advisory Board for various international journals. Assoc. Prof. Shukur research focuses on (1) composites structure durability, (2) recycling composites, (3) biomimetics design, (4) composites testing, (5) natural fibre composites, (6) hybrid composites, (7) composites bonding, and (8) polymer mortar composites. He is registered with the Board of Engineers Malaysia (BEM) and Member of Malaysian Board of Technologist (MBOT). To date, he has been a consultant to 16 projects amounting to RM 480,000.00 and Principal Investigator and member to 49 research grants worth RM 3.9 million. He has also supervised and co-supervised 16 postgraduates. He has also authored and co-authored 106 publications (published/accepted): 22 ISI Article/Journal, 9 Scopus Article/Journal, 10 Non-Indexed Article/Journal, 2 Indexed Conference Proceedings, 43 Non-Indexed Conference Proceedings, 7 Book Chapters, and 13 Classified/Technical/Expert Reports. He also has one patent granted, one patent filed, and seven registered industrial designs.

#### **Nabil Hayeemasae**

Nabil Hayeemasae is currently an Assistant Professor at the Department of Rubber Technology and Polymer Science, Faculty of Science and Technology, Prince of Songkla University, Pattani Campus. He obtained a Ph.D. in Polymer Engineering at the School of Materials and Mineral Resources Engineering, Universiti Sains Malaysia. He has conducted research in several research areas, particularly on the re-utilization of polymer and solid wastes into potential materials, rubber composites, and rubber recycling. So far, his research outputs have been selected to be published in over 80 papers internationally.

#### **Khubab Shaker**

Khubab Shaker currently serves as an Assistant Professor and Chairman of the Department of Materials at National Textile University, Pakistan. He is an active Textile Composite Materials Research Group member at NTU, with 73 publications (impact factor = 200), 6 books, and 34 conference presentations. He has more than 13 years of industrial and teaching experience in textiles and composite materials. His research areas include composites joining, natural fibers, biopolymers, green composites, and their applications as potential replacements for synthetic composite materials.

## **Preface to "Fiber Reinforced Polymer Composites: Mechanical Properties and Applications"**

Composites are the most promising material available in the twenty-first century. Composites reinforced with synthetic or natural fibers are becoming extremely prevalent as the demand for lightweight materials with high strength for specific applications grows. Fiber-reinforced polymer composite has a high strength-to-weight ratio and excellent properties, such as high durability, stiffness, damping property, flexural strength, and resistance to corrosion, wear, impact, and fire. Various properties of composite materials have led to applications in construction, aerospace, automobile, biomedical, marine, and many other industries. Because their constituent elements and manufacturing techniques primarily determine the performance of composite materials, the functional properties of various fibers available worldwide, their classifications, and the manufacturing techniques used to fabricate the composite materials must be investigated.

## **R.A. Ilyas, S.M. Sapuan, Emin Bayraktar, Shukur Abu Hassan, Nabil Hayeemasae, and Khubab**

**Shaker** *Editors*

## *Editorial* **Fibre-Reinforced Polymer Composites: Mechanical Properties and Applications**

**R. A. Ilyas 1,2,3,\* , S. M. Sapuan <sup>4</sup> , Emin Bayraktar <sup>5</sup> , Shukur Abu Hassan <sup>2</sup> , Nabil Hayeemasae <sup>6</sup> , M. S. N. Atikah <sup>7</sup> and Khubab Shaker <sup>8</sup>**


*"Fibre-Reinforced Polymer Composites: Mechanical Properties and Applications"* is a newly open Special Issue of *Polymers*, which aims to publish original and review papers on new scientific and applied research and make boundless contributions to the finding and understanding of the reinforcing effects of various synthetic and natural fibres on the performance of biopolymer composites. This Special Issue also covers the fundamentals, characterisations, and applications of synthetic and natural fibre-reinforced biopolymer composites.

Rapid growth in the manufacturing industries has necessitated the improvement in materials in terms of density, stiffness, strength, and cost-effectiveness with increased sustainability. Composite materials have been developed as one of the materials with such improvements in these qualities that serve their promise in various applications. Composite materials are made up of two or more elements, one of which is present in the matrix phase (synthetic or biopolymer [1–6]) and the other in particle or fibre form. Composites have been discovered to be the most promising material available in the twenty-first century. Composites reinforced with synthetic or natural fibres are becoming extremely prevalent as the market grows in demand for lightweight materials with high strength for specific applications. The matrix, which serves primarily to hold the reinforcement together, is also regarded as resin, particularly in the case of polymers.

In order to create eco-friendly composites, natural fillers such as natural fibres [7–9], nanocrystalline cellulose, nanofibrillated cellulose [10,11], bacterial nanocellulose [12], and chitosan have been added to the polymer matrix. This increased material qualities while minimising the problem of residue formation. Many researchers have reported the benefits of cellulosic fibres, including the fact that they are abundant in nature, renewable, cost-effective, and non-toxic, as well as providing necessary bonding with the cementbased matrix for significant improvements in material properties such as flexural capacity, toughness, ductility, and impact resistance.

Several studies that investigated the fibre-loading effect on polymer composites found that it had a good relationship with tensile strength. Studies on the fibre loading effect that

**Citation:** Ilyas, R.A.; Sapuan, S.M.; Bayraktar, E.; Hassan, S.A.; Hayeemasae, N.; Atikah, M.S.N.; Shaker, K. Fibre-Reinforced Polymer Composites: Mechanical Properties and Applications. *Polymers* **2022**, *14*, 3732. https://doi.org/10.3390/ polym14183732

Received: 19 August 2022 Accepted: 26 August 2022 Published: 7 September 2022

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2022 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

led to the tensile strength were observed [13,14]. It was demonstrated that the optimum fibre loading for kenaf/thermoplastic polyurethane composites was 30% [15]. Other studies regarding kenaf fibre and phenol-formaldehyde (KF/PF) composites reported that kenaf fibre loading up to 43% showed the best tensile strength for the composites [16].

A study on a kenaf-fibre-reinforced, corn-starch-based biocomposite film investigated the fibre loading effect for the tensile properties. It was found that at 6% of the kenaf fibre loading, the optimum tensile properties (17.74 MPa) were observed [17]. Studies on *Cymbopogan citratus* fibre with cassava starch showed higher tensile properties (19.27 MPa) with the use of 50% fibre loading [18]. Another study on arrowroot (*Maranta arundinacea*) fibre-reinforced arrowroot starch biopolymer composites conducted by Tarique et al. [19] shows that the mechanical properties were enhanced up to 15.22 MPa with the optimum filler content was 10%. Previous studies on the correlation of fibre length to tensile strength were conducted by several researchers. Studies on wheat husk length reinforced rubber composites showed that the highest tensile strength achieved was the medium length of fibre (125–250µm), where the fibre was arranged longitudinally [20]. Jihua Zhu [21] conducted a study on glass-fibre-reinforced composites under acid–base and salt environments. The results indicated that the tensile strength of the GFRP decreased by 22%, 71%, and 87% after 56 d of exposure to 5% NaOH solutions at 20 ◦C, 50 ◦C, and 80 ◦C, respectively.

Not only do fibre-reinforced polymer composites have a high strength-to-weight ratio, but they also have excellent properties such as high durability, stiffness, damping property, flexural strength, and resistance to corrosion, wear, impact, and fire. Various properties of composites materials have led to applications in construction, aerospace, packaging [22–24], electronic, electrical, structural, energy storage [25], automotive [26], filter, coating, bone tissue engineering, and drug delivery [27] and many other industries. Because the performance of composite materials is primarily determined by their constituent elements and manufacturing techniques, the functional properties of various fibres available worldwide, their classifications, and the manufacturing techniques used to fabricate the composite materials must be investigated.

This Special Issue will also cover recent advances in composite processing, mechanical characterisation, and potential applications. Both synthetic and natural fibre-reinforced polymer composites are welcome. Moreover, we welcome approaches to this issue from several vital directions, such as the production of fibres, surface and interfacial characterisation of its properties, economic feasibility, challenges, and future perspectives in the field of polymer composites. As a result of this Special Issue, current and future literature data can be enriched.

**Author Contributions:** Conceptualization, R.A.I.; writing—original draft preparation, R.A.I., S.M.S., E.B., S.A.H., N.H., M.S.N.A. and K.S.; writing—review and editing, R.A.I., S.M.S., E.B., S.A.H., N.H., M.S.N.A. and K.S.; project administration, R.A.I., S.M.S., E.B., S.A.H., N.H., M.S.N.A. and K.S.; funding acquisition, R.A.I. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by Universiti Teknologi Malaysia grant number PY/2022/02318— Q.J130000.3851.21H99, and Ministry of Higher Education Malaysia (MOHE) grant number JPT (BPKI) 1000/016/018/25 (57).

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** Not applicable.

**Acknowledgments:** The authors would like express gratitude for the financial support received from the Universiti Teknologi Malaysia, the project "The impact of Malaysian bamboos' chemical and fibre characteristics on their pulp and paper properties, grant number PY/2022/02318— Q.J130000.3851.21H99". The research has been carried out under the program Research Excellence Consortium (JPT (BPKI) 1000/016/018/25 (57)) provided by the Ministry of Higher Education Malaysia (MOHE).

**Conflicts of Interest:** The authors declare no conflict of interest.

## **References**


24. Nazrin, A.; Sapuan, S.M.; Zuhri, M.Y.M.; Tawakkal, I.S.M.A.; Ilyas, R.A. Flammability and physical stability of sugar palm crystalline nanocellulose reinforced thermoplastic sugar palm starch/poly(lactic acid) blend bionanocomposites. *Nanotechnol. Rev.* **2022**, *11*, 86–95. [CrossRef] 25. Nurazzi, N.M.; Sabaruddin, F.A.; Harussani, M.M.; Kamarudin, S.H.; Rayung, M.; Asyraf, M.R.M.; Aisyah, H.A.; Norrrahim, M.N.F.; Ilyas, R.A.; Abdullah, N.; et al. Mechanical Performance and Applications of CNTs Reinforced Polymer Composites— A Review. *Nanomaterials* **2021**, *11*, 2186. https://doi.org/10.3390/nano11092186.

crystalline nanocellulose reinforced thermoplastic sugar palm starch/poly(lactic acid) blend bionanocomposites. *Nanotechnol.* 

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20. Sobhy, M.S.; Tammam, M.T. The Influence of Fiber Length and Concentration on the Physical Properties of Wheat Husk Fibers

21. Zhu, J.; Deng, Y.; Chen, P.; Wang, G.; Min, H.; Fang, W. Prediction of Long-Term Tensile Properties of Glass Fiber Reinforced Composites under Acid-Base and Salt Environments. *Polymers* **2022**, *14*, 3031. https://doi.org/10.3390/polym14153031. 22. Rozilah, A.; Jaafar, C.N.A.; Sapuan, S.M.; Zainol, I.; Ilyas, R.A. The Effects of Silver Nanoparticles Compositions on the Mechanical, Physiochemical, Antibacterial, and Morphology Properties of Sugar Palm Starch Biocomposites for Antibacterial

23. Alias, A.H.; Norizan, M.N.; Sabaruddin, F.A.; Asyraf, M.R.M.; Norrrahim, M.N.F.; Ilyas, A.R.; Kuzmin, A.M.; Rayung, M.; Shazleen, S.S.; Nazrin, A.; et al. Hybridization of MMT/Lignocellulosic Fiber Reinforced Polymer Nanocomposites for Structural

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#### **Short Biography of Authors R. A. Ilyas** is a senior lecturer at the School of Chemical and Energy Engineering, Faculty of Engi-

neering, Universiti Teknologi Malaysia (UTM), Malaysia. He received his Diploma in Forestry at Universiti Putra Malaysia, Bintulu Campus (UPMKB), Sarawak, Malaysia, from May 2009 to April 2012. In 2012, he was awarded the Public Service Department (JPA) scholarship to pursue his bachelor's degree (BSc) in Chemical Engineering at Universiti Putra Malaysia (UPM). Upon completing his BSc. programme in 2016, he was again awarded the Graduate Research Fellowship (GRF) by the Universiti Putra Malaysia (UPM) to undertake a PhD degree in the field of Biocomposite Technology and Design at Institute of Tropical Forestry and Forest Products (INTROP) UPM. R.A. Ilyas was the recipient of MVP Doctor of Philosophy Gold Medal Award UPM 2019, for Best Ph.D. Thesis and Top Student Award, INTROP, UPM. He was awarded Outstanding Reviewer by Carbohydrate Polymers, Elsevier United Kingdom, Best Paper Award (11th AUN/SEED-Net Regional Conference on Energy Engineering), and National Book Award 2018, Best Paper Award (Seminar Enau Kebangsaan 2019, Persatuan Pembangunan dan Industri Enau Malaysia) and Top Cited Article 2020-2021 Journal Polymer Composite, Wiley, 2022. R.A. Ilyas also was listed and awarded among the World's Top 2% Scientists (Subject-Wise) citation impact during the single calendar year 2019 and 2020 by Stanford University, US, PERINTIS Publication Award 2021 and 2022 by Persatuan Saintis Muslim Malaysia, Emerging Scholar Award by Automotive and Autonomous Systems 2021, Belgium, Young Scientists Network—Academy of Sciences Malaysia (YSN-ASM) 2021, UTM Young Research Award 2021, UTM Publication Award 2021, and UTM Highly Cited Researcher Award 2021. His main research interests are: (1) polymer engineering (biodegradable polymers, biopolymers, polymer composites, polymer gels) and (2) material engineering (natural fibre-reinforced polymer composites, biocomposites, cellulose materials, nano-composites). To date, he has authored or co-authored more than 404 publications (published/accepted): 164 Journals Indexed in JCR/Scopus, 2 non-index Journal, 15 books, 104 book chapters, 78 conference proceedings/seminars, 4 research bulletins, 10 conference papers (abstract published in book of abstract), 17 Guest Editor of Journal Special Issues and 10 Editor/Co-Editor of Conference/Seminar Proceedings on green materials related subjects. **R. A. Ilyas** is a senior lecturer at the School of Chemical and Energy Engineering, Faculty of Engineering, Universiti Teknologi Malaysia (UTM), Malaysia. He received his Diploma in Forestry at Universiti Putra Malaysia, Bintulu Campus (UPMKB), Sarawak, Malaysia, from May 2009 to April 2012. In 2012, he was awarded the Public Service Department (JPA) scholarship to pursue his bachelor's degree (BSc) in Chemical Engineering at Universiti Putra Malaysia (UPM). Upon completing his BSc. programme in 2016, he was again awarded the Graduate Research Fellowship (GRF) by the Universiti Putra Malaysia (UPM) to undertake a PhD degree in the field of Biocomposite Technology and Design at Institute of Tropical Forestry and Forest Products (INTROP) UPM. R.A. Ilyas was the recipient of MVP Doctor of Philosophy Gold Medal Award UPM 2019, for Best Ph.D. Thesis and Top Student Award, INTROP, UPM. He was awarded Outstanding Reviewer by Carbohydrate Polymers, Elsevier United Kingdom, Best Paper Award (11th AUN/SEED-Net Regional Conference on Energy Engineering), and National Book Award 2018, Best Paper Award (Seminar Enau Kebangsaan 2019, Persatuan Pembangunan dan Industri Enau Malaysia) and Top Cited Article 2020-2021 Journal Polymer Composite, Wiley, 2022. R.A. Ilyas also was listed and awarded among the World's Top 2% Scientists (Subject-Wise) citation impact during the single calendar year 2019 and 2020 by Stanford University, US, PERINTIS Publication Award 2021 and 2022 by Persatuan Saintis Muslim Malaysia, Emerging Scholar Award by Automotive and Autonomous Systems 2021, Belgium, Young Scientists Network—Academy of Sciences Malaysia (YSN-ASM) 2021, UTM Young Research Award 2021, UTM Publication Award 2021, and UTM Highly Cited Researcher Award 2021. His main research interests are: (1) polymer engineering (biodegradable polymers, biopolymers, polymer composites, polymer gels) and (2) material engineering (natural fibre-reinforced polymer composites, biocomposites, cellulose materials, nano-composites). To date, he has authored or co-authored more than 404 publications (published/accepted): 164 Journals Indexed in JCR/Scopus, 2 non-index Journal, 15 books, 104 book chapters, 78 conference proceedings/seminars, 4 research bulletins, 10 conference papers (abstract published in book of abstract), 17 Guest Editor of Journal Special Issues and 10 Editor/Co-Editor of Conference/Seminar Proceedings on green materials related subjects.

**Mohd Sapuan Salit** is an "A" Grade Professor of composite materials at Department of Mechanical and Manufacturing, Universiti Putra Malaysia (UPM) and a Head of Laboratory of Biocomposite Technology, INTROP, UPM. He has a BEng in Mechanical Engineering from University of Newcastle, Australia, an MSc in Engineering Design from Loughborough University, UK and PhD in Material Engineering from De Montfort University, UK. He is a Professional Engineer, a Society of Automotive Engineers Fellow, an Academy of Science Malaysia Fellow, a Plastic and Rubber and Institute Malaysia Fellow, a Malaysian Scientific Association Fellow, an International Biographical Association Fellow and an Institute of Material Malaysia Fellow. He is an Honorary Member and immediate past Vice President of Asian Polymer Association based in IIT Delhi and Founding Chairman and Honorary Member of Society of Sugar Palm Development and Industry, Malaysia. He is the co-editor-in-chief of Functional Composites and Structures, and member of editorial boards of more than two dozen journals. To date he has produced more than 1800 publications including over 860 journal papers, 50 books, and 175 chapters in book. He has delivered over 50 plenary and keynote lectures, and over 150 invited lectures. He organized 30 journal special issues as a guest editor, presented over 650 technical articles in conferences and seminars, reviewed over 1300 journal papers and has 8 patents. He successfully supervised 91 PhD and 70 MSc students and 15 postdoctoral researchers. His current h-index is 93 and the number of citations is 31,647 (Google Scholar). He received nine Outstanding Researcher Awards from UPM, ISESCO Science Award (Gold Medal), Plastic and Rubber Institute Malaysia Fellowship Award and Forest Research Institute Malaysia First Prize Publication Award. He also received Khwarizimi International Award, SEARCA Regional Professorial Chair award, Kuala Lumpur Royal Rotary Gold Medal Research Award and two National Book Awards. He received the Endeavour Research Promotion Award by TMU/IEEE India, Citation of Excellence Award, Emerald, UK, Malaysia's Research Star Award, Elsevier/Ministry of Education Malaysia, Publons Peer Review Award, Publons, USA, Professor of Eminence Award from Aligarh Muslim University, India, Top Research Scientists' Malaysia Award, Academy of Science Malaysia, Gold in Invention and Innovation Awards, Malaysia Technology Expo and PERINTIS Publication Award, PERINTIS, Malaysia. He was listed among the World Top 2% Scientists by Stanford University, USA. He is the finalist of IET Achievements Award, IET, UK and 2021 SAE Subir Chowdhury Medal of Quality Leadership, SAE, USA. **Emin Bayraktar** (Prof. Emeritus, Habil., Dr (Ph.D.), DSc—Doctor of Science) is an academic and research staff member in Mechanical and Manufacturing engineering at SUPMECA/Paris, France. **Mohd Sapuan Salit** is an "A" Grade Professor of composite materials at Department of Mechanical and Manufacturing, Universiti Putra Malaysia (UPM) and a Head of Laboratory of Biocomposite Technology, INTROP, UPM. He has a BEng in Mechanical Engineering from University of Newcastle, Australia, an MSc in Engineering Design from Loughborough University, UK and PhD in Material Engineering from De Montfort University, UK. He is a Professional Engineer, a Society of Automotive Engineers Fellow, an Academy of Science Malaysia Fellow, a Plastic and Rubber and Institute Malaysia Fellow, a Malaysian Scientific Association Fellow, an International Biographical Association Fellow and an Institute of Material Malaysia Fellow. He is an Honorary Member and immediate past Vice President of Asian Polymer Association based in IIT Delhi and Founding Chairman and Honorary Member of Society of Sugar Palm Development and Industry, Malaysia. He is the co-editor-inchief of Functional Composites and Structures, and member of editorial boards of more than two dozen journals. To date he has produced more than 1800 publications including over 860 journal papers, 50 books, and 175 chapters in book. He has delivered over 50 plenary and keynote lectures, and over 150 invited lectures. He organized 30 journal special issues as a guest editor, presented over 650 technical articles in conferences and seminars, reviewed over 1300 journal papers and has 8 patents. He successfully supervised 91 PhD and 70 MSc students and 15 postdoctoral researchers. His current h-index is 93 and the number of citations is 31,647 (Google Scholar). He received nine Outstanding Researcher Awards from UPM, ISESCO Science Award (Gold Medal), Plastic and Rubber Institute Malaysia Fellowship Award and Forest Research Institute Malaysia First Prize Publication Award. He also received Khwarizimi International Award, SEARCA Regional Professorial Chair award, Kuala Lumpur Royal Rotary Gold Medal Research Award and two National Book Awards. He received the Endeavour Research Promotion Award by TMU/IEEE India, Citation of Excellence Award, Emerald, UK, Malaysia's Research Star Award, Elsevier/Ministry of Education Malaysia, Publons Peer Review Award, Publons, USA, Professor of Eminence Award from Aligarh Muslim University, India, Top Research Scientists' Malaysia Award, Academy of Science Malaysia, Gold in Invention and Innovation Awards, Malaysia Technology Expo and PERINTIS Publication Award, PERINTIS, Malaysia. He was listed among the World Top 2% Scientists by Stanford University, USA. He is the finalist of IET Achievements Award, IET, UK and 2021 SAE Subir Chowdhury Medal of Quality Leadership, SAE, USA. **Mohd Sapuan Salit** is an "A" Grade Professor of composite materials at Department of Mechanical and Manufacturing, Universiti Putra Malaysia (UPM) and a Head of Laboratory of Biocomposite Technology, INTROP, UPM. He has a BEng in Mechanical Engineering from University of Newcastle, Australia, an MSc in Engineering Design from Loughborough University, UK and PhD in Material Engineering from De Montfort University, UK. He is a Professional Engineer, a Society of Automotive Engineers Fellow, an Academy of Science Malaysia Fellow, a Plastic and Rubber and Institute Malaysia Fellow, a Malaysian Scientific Association Fellow, an International Biographical Association Fellow and an Institute of Material Malaysia Fellow. He is an Honorary Member and immediate past Vice President of Asian Polymer Association based in IIT Delhi and Founding Chairman and Honorary Member of Society of Sugar Palm Development and Industry, Malaysia. He is the co-editor-in-chief of Functional Composites and Structures, and member of editorial boards of more than two dozen journals. To date he has produced more than 1800 publications including over 860 journal papers, 50 books, and 175 chapters in book. He has delivered over 50 plenary and keynote lectures, and over 150 invited lectures. He organized 30 journal special issues as a guest editor, presented over 650 technical articles in conferences and seminars, reviewed over 1300 journal papers and has 8 patents. He successfully supervised 91 PhD and 70 MSc students and 15 postdoctoral researchers. His current h-index is 93 and the number of citations is 31,647 (Google Scholar). He received nine Outstanding Researcher Awards from UPM, ISESCO Science Award (Gold Medal), Plastic and Rubber Institute Malaysia Fellowship Award and Forest Research Institute Malaysia First Prize Publication Award. He also received Khwarizimi International Award, SEARCA Regional Professorial Chair award, Kuala Lumpur Royal Rotary Gold Medal Research Award and two National Book Awards. He received the Endeavour Research Promotion Award by TMU/IEEE India, Citation of Excellence Award, Emerald, UK, Malaysia's Research Star Award, Elsevier/Ministry of Education Malaysia, Publons Peer Review Award, Publons, USA, Professor of Eminence Award from Aligarh Muslim University, India, Top Research Scientists' Malaysia Award, Academy of Science Malaysia, Gold in Invention and Innovation Awards, Malaysia Technology Expo and PERINTIS Publication Award, PERINTIS, Malaysia. He was listed among the World Top 2% Scientists by Stanford University, USA. He is the finalist of IET Achievements Award, IET, UK and 2021 SAE Subir Chowdhury Medal of Quality Leadership, SAE, USA.

metal Forming of thin sheets (Design + test + FEM), static and dynamic behavior and optimization of materials (experimental and FEM—utilization and design of composite-based metallic and nonmetallic, powder metallurgy, and energetic material aeronautical applications), metallic-based and non-metallic materials, powder metallurgy and metallurgy of steels, welding, and heat treatment, as well as the processing of new composites, sintering techniques, sinter–forging, thixoforming, etc. He has authored more than 200 publications in the International Journals and International Conference Proceedings, and has also authored more than 90 research reports (European = Steel Committee projects, Test + Simulation). He already advised 32 Ph.D. and 120 MSc theses, and is currently advising 7. He is a Fellow of WAMME (World Academy of Science in Materials and Manufacturing Engineering), Editorial Board—Member of JAMME (*International Journal of Achievement in Materials and Manufacturing Engineering*), Advisory board member of AMPT—2009 (*Advanced Materials Processing technologies*), APCMP—2008 and APCMP—2010. He was Visiting Professor at Nanyang Technology University, Singapore in 2012, Xi'an Northwestern Technical University, Aeronautical Engineering, in 2016, University of Campinas, UNICAMP-Brazil in 2013 until 2023. He is a recipient of the Silesian University Prix pour "FREDERIK STAUB Golden Medal-2009" by the Academy of WAMME, "World Academy of Science"—Poland, materials science section, and a recipient of the William Johnson International Gold Medal—2014, AMPT academic association. **Emin Bayraktar** (Prof. Emeritus, Habil., Dr (Ph.D.), DSc—Doctor of Science) is an academic and research staff member in Mechanical and Manufacturing engineering at SUPMECA/Paris, France. His research areas include manufacturing techniques of new materials (basic composites—hybrid), metal Forming of thin sheets (Design + test + FEM), static and dynamic behavior and optimization of materials (experimental and FEM—utilization and design of composite-based metallic and nonmetallic, powder metallurgy, and energetic material aeronautical applications), metallic-based and non-metallic materials, powder metallurgy and metallurgy of steels, welding, and heat treatment, as well as the processing of new composites, sintering techniques, sinter–forging, thixoforming, etc. He has authored more than 200 publications in the International Journals and International Conference Proceedings, and has also authored more than 90 research reports (European = Steel Committee projects, Test + Simulation). He already advised 32 Ph.D. and 120 MSc theses, and is currently advising 7. He is a Fellow of WAMME (World Academy of Science in Materials and Manufacturing Engineering), Editorial Board—Member of JAMME (*International Journal of Achievement in Materials and Manufacturing Engineering*), Advisory board member of AMPT—2009 (*Advanced Materials Processing technologies*), APCMP—2008 and APCMP—2010. He was Visiting Professor at Nanyang Technology University, Singapore in 2012, Xi'an Northwestern Technical University, Aeronautical Engineering, in 2016, University of Campinas, UNICAMP-Brazil in 2013 until 2023. He is a recipient of the Silesian University Prix pour "FREDERIK STAUB Golden Medal-2009" by the Academy of WAMME, "World Academy of Science"—Poland, materials science section, and a recipient of the William Johnson International Gold Medal—2014, AMPT academic association. **Emin Bayraktar** (Prof. Emeritus, Habil., Dr (Ph.D.), DSc—Doctor of Science) is an academic and research staff member in Mechanical and Manufacturing engineering at SUPMECA/Paris, France. His research areas include manufacturing techniques of new materials (basic composites—hybrid), metal Forming of thin sheets (Design + test + FEM), static and dynamic behavior and optimization of materials (experimental and FEM—utilization and design of composite-based metallic and nonmetallic, powder metallurgy, and energetic material aeronautical applications), metallic-based and non-metallic materials, powder metallurgy and metallurgy of steels, welding, and heat treatment, as well as the processing of new composites, sintering techniques, sinter–forging, thixoforming, etc. He has authored more than 200 publications in the International Journals and International Conference Proceedings, and has also authored more than 90 research reports (European = Steel Committee projects, Test + Simulation). He already advised 32 Ph.D. and 120 MSc theses, and is currently advising 7. He is a Fellow of WAMME (World Academy of Science in Materials and Manufacturing Engineering), Editorial Board—Member of JAMME (*International Journal of Achievement in Materials and Manufacturing Engineering*), Advisory board member of AMPT—2009 (*Advanced Materials Processing technologies*), APCMP—2008 and APCMP—2010. He was Visiting Professor at Nanyang Technology University, Singapore in 2012, Xi'an Northwestern Technical University, Aeronautical Engineering, in 2016, University of Campinas, UNICAMP-Brazil in 2013 until 2023. He is a recipient of the Silesian University Prix pour "FREDERIK STAUB Golden Medal-2009" by the Academy of WAMME, "World Academy of Science"—Poland, materials science section, and a recipient of the William Johnson International Gold Medal—2014, AMPT academic association.

His research areas include manufacturing techniques of new materials (basic composites—hybrid),

is an Engineer, associate professor and current Director of Centre for Advanced Composite Materials (CACM), Universiti Teknologi Malaysia (UTM), Malaysia. In his teaching career of 33 years, he has been working as Centre Director, Head of Composite Lab and Committee Member at the Faculty and University level. He is also involved as Consultant and on the Editorial Advisory Board for various international journals. Assoc. Prof. Shukur research focuses on (1) composites structure durability, (2) recycling composites, (3) biomimetics design, (4) composites testing, (5) natural fibre composites, (6) hybrid composites, (7) composites bonding and (8) polymer mortar composites. He is registered with Board of Engineers Malaysia (BEM) and Member of Malaysian Board of Technologist (MBOT). To date he has been a consultant to 16 projects amounting to RM 480,000. 00 and Principal Investigator and member to 49 research grants worth RM 3.9 million. He has also supervised and co-supervised 16 postgraduates. He has also authored and co-authored 106 publications (published/accepted); 22 ISI Article/Journal, 9 Scopus Article/Journal, 10 Non-Indexed Article/Jour-**Shukur Abu Hassan**, PhD (UTM, Malaysia), MSc (Warwick University, UK), BSc (UTM, Malaysia) is an Engineer, associate professor and current Director of Centre for Advanced Composite Materials (CACM), Universiti Teknologi Malaysia (UTM), Malaysia. In his teaching career of 33 years, he has been working as Centre Director, Head of Composite Lab and Committee Member at the Faculty and University level. He is also involved as Consultant and on the Editorial Advisory Board for various international journals. Assoc. Prof. Shukur research focuses on (1) composites structure durability, (2) recycling composites, (3) biomimetics design, (4) composites testing, (5) natural fibre composites, (6) hybrid composites, (7) composites bonding and (8) polymer mortar composites. He is registered with Board of Engineers Malaysia (BEM) and Member of Malaysian Board of Technologist (MBOT). To date he has been a consultant to 16 projects amounting to RM 480,000. 00 and Principal Investigator and member to 49 research grants worth RM 3.9 million. He has also supervised and co-supervised 16 postgraduates. He has also authored and co-authored 106 publications (published/accepted); 22 ISI Article/Journal, 9 Scopus Article/Journal, 10 Non-Indexed Article/Journal, 2 Indexed Conference Proceedings, 43 Non-Indexed Conference Proceedings, 7 Book Chapters, 13 Classified/Technical /Expert Report. He also has 1 patent granted, 1 patent filed and 7 registered industrial design. has been working as Centre Director, Head of Composite Lab and Committee Member at the Faculty and University level. He is also involved as Consultant and on the Editorial Advisory Board for various international journals. Assoc. Prof. Shukur research focuses on (1) composites structure durability, (2) recycling composites, (3) biomimetics design, (4) composites testing, (5) natural fibre composites, (6) hybrid composites, (7) composites bonding and (8) polymer mortar composites. He is registered with Board of Engineers Malaysia (BEM) and Member of Malaysian Board of Technologist (MBOT). To date he has been a consultant to 16 projects amounting to RM 480,000. 00 and Principal Investigator and member to 49 research grants worth RM 3.9 million. He has also supervised and co-supervised 16 postgraduates. He has also authored and co-authored 106 publications (published/accepted); 22 ISI Article/Journal, 9 Scopus Article/Journal, 10 Non-Indexed Article/Journal, 2 Indexed Conference Proceedings, 43 Non-Indexed Conference Proceedings, 7 Book Chapters, 13 Classified/Technical /Expert Report. He also has 1 patent granted, 1 patent filed and 7 registered industrial design. and University level. He is also involved as Consultant and on the Editorial Advisory Board for various international journals. Assoc. Prof. Shukur research focuses on (1) composites structure durability, (2) recycling composites, (3) biomimetics design, (4) composites testing, (5) natural fibre composites, (6) hybrid composites, (7) composites bonding and (8) polymer mortar composites. He is registered with Board of Engineers Malaysia (BEM) and Member of Malaysian Board of Technologist (MBOT). To date he has been a consultant to 16 projects amounting to RM 480,000. 00 and Principal Investigator and member to 49 research grants worth RM 3.9 million. He has also supervised and co-supervised 16 postgraduates. He has also authored and co-authored 106 publications (published/accepted); 22 ISI Article/Journal, 9 Scopus Article/Journal, 10 Non-Indexed Article/Journal, 2 Indexed Conference Proceedings, 43 Non-Indexed Conference Proceedings, 7 Book Chapters, 13 Classified/Technical /Expert Report. He also has 1 patent granted, 1 patent filed and 7 registered industrial design. **Nabil Hayeemasae** is currently an Assistant Professor at the Department of Rubber Technology and

**Shukur Abu Hassan**, PhD (UTM, Malaysia), MSc (Warwick University, UK), BSc (UTM, Malaysia) is an Engineer, associate professor and current Director of Centre for Advanced Composite Materials (CACM), Universiti Teknologi Malaysia (UTM), Malaysia. In his teaching career of 33 years, he

**Shukur Abu Hassan**, PhD (UTM, Malaysia), MSc (Warwick University, UK), BSc (UTM, Malaysia) is an Engineer, associate professor and current Director of Centre for Advanced Composite Materials (CACM), Universiti Teknologi Malaysia (UTM), Malaysia. In his teaching career of 33 years, he has been working as Centre Director, Head of Composite Lab and Committee Member at the Faculty

*Polymers* **2022**, *14*, x FOR PEER REVIEW 6 of 7

*Polymers* **2022**, *14*, x FOR PEER REVIEW 6 of 7

**Shukur Abu Hassan**, PhD (UTM, Malaysia), MSc (Warwick University, UK), BSc (UTM, Malaysia)

**Nabil Hayeemasae** is currently an Assistant Professor at the Department of Rubber Technology and Polymer Science, Faculty of Science and Technology, Prince of Songkla University, Pattani Campus.

**M.S.N. Atikah** is a PhD candidate in the field of Chemical Engineering at Universiti Putra Malaysia.

nal, 2 Indexed Conference Proceedings, 43 Non-Indexed Conference Proceedings, 7 Book Chapters, 13 Classified/Technical /Expert Report. He also has 1 patent granted, 1 patent filed and 7 registered industrial design. **Nabil Hayeemasae** is currently an Assistant Professor at the Department of Rubber Technology and Polymer Science, Faculty of Science and Technology, Prince of Songkla University, Pattani Campus. He obtained a Ph.D. in Polymer Engineering at the School of Materials and Mineral Resources Engineering, Universiti Sains Malaysia. He has been conducting several research areas, particularly on the re-utilization of polymer and solid wastes into potential materials, rubber composites, and rubber recycling. So far, his research outputs have been selected to be published in over 80 papers internationally. **Nabil Hayeemasae** is currently an Assistant Professor at the Department of Rubber Technology and Polymer Science, Faculty of Science and Technology, Prince of Songkla University, Pattani Campus. He obtained a Ph.D. in Polymer Engineering at the School of Materials and Mineral Resources Engineering, Universiti Sains Malaysia. He has been conducting several research areas, particularly on the re-utilization of polymer and solid wastes into potential materials, rubber composites, and rubber recycling. So far, his research outputs have been selected to be published in over 80 papers internationally. Polymer Science, Faculty of Science and Technology, Prince of Songkla University, Pattani Campus. He obtained a Ph.D. in Polymer Engineering at the School of Materials and Mineral Resources Engineering, Universiti Sains Malaysia. He has been conducting several research areas, particularly on the re-utilization of polymer and solid wastes into potential materials, rubber composites, and rubber recycling. So far, his research outputs have been selected to be published in over 80 papers internationally.

ternationally.

He obtained a Ph.D. in Polymer Engineering at the School of Materials and Mineral Resources Engineering, Universiti Sains Malaysia. He has been conducting several research areas, particularly on the re-utilization of polymer and solid wastes into potential materials, rubber composites, and rubber recycling. So far, his research outputs have been selected to be published in over 80 papers in-**M.S.N. Atikah** is a PhD candidate in the field of Chemical Engineering at Universiti Putra Malaysia. She graduated with a Bachelor of Engineering (Chemical) and a Master of Science (Chemical) from the same university. She has authored and co-authored 54 publications from 2016-2022 and has an h-index of 20. She was awarded scholarships for her bachelor's degree and doctor of philosophy programs from Jabatan Perkhidmatan Awam (JPA), Malaysia. **M.S.N. Atikah** is a PhD candidate in the field of Chemical Engineering at Universiti Putra Malaysia. She graduated with a Bachelor of Engineering (Chemical) and a Master of Science (Chemical) from the same university. She has authored and co-authored 54 publications from 2016-2022 and has an h-index of 20. She was awarded scholarships for her bachelor's degree and doctor of philosophy programs from Jabatan Perkhidmatan Awam (JPA), Malaysia. **M.S.N. Atikah** is a PhD candidate in the field of Chemical Engineering at Universiti Putra Malaysia. She graduated with a Bachelor of Engineering (Chemical) and a Master of Science (Chemical) from the same university. She has authored and co-authored 54 publications from 2016-2022 and has an h-index of 20. She was awarded scholarships for her bachelor's degree and doctor of philosophy programs from Jabatan Perkhidmatan Awam (JPA), Malaysia.

She graduated with a Bachelor of Engineering (Chemical) and a Master of Science (Chemical) from the same university. She has authored and co-authored 54 publications from 2016-2022 and has an h-index of 20. She was awarded scholarships for her bachelor's degree and doctor of philosophy programs from Jabatan Perkhidmatan Awam (JPA), Malaysia. **Khubab Shaker** is currently serving as Assistant Professor and Chairman, Department of Materials at National Textile University, Pakistan. He is an active member of the Textile Composite Materials Research Group at NTU, with 63 publications (impact factor=166), 3 books, and 32 conference presentations. He has more than 12 years of industrial and teaching experience in textiles and composite materials. His areas of research include composites joining, natural fibres, biopolymers, green composites, and their applications as a **Khubab Shaker** is currently serving as Assistant Professor and Chairman, Department of Materials at National Textile University, Pakistan. He is an active member of the Textile Composite Materials Research Group at NTU, with 63 publications (impact factor=166), 3 books, and 32 conference presentations. He has more than 12 years of industrial and teaching experience in textiles and composite materials. His areas of research include composites joining, natural fibres, biopolymers, green composites, and their applications as a potential replacement for synthetic composite materials.

potential replacement for synthetic composite materials.

## *Article* **Effect of** *Cymbopogan citratus* **Fibre on Physical and Impact Properties of Thermoplastic Cassava Starch/Palm Wax Composites**

**Zatil Hafila Kamaruddin 1,2,\*, Ridhwan Jumaidin 3,\* , Zatil Hazrati Kamaruddin <sup>2</sup> , Muhammad Rizal Muhammad Asyraf 4,5,\* , Muhammad Rizal Razman <sup>6</sup> and Tabrej Khan <sup>7</sup>**


**Abstract:** *Cymbopogan citratus* fibre (CCF) is an agricultural waste plant derived from a natural cellulosic source of fibre that can be used in various bio-material applications. This paper beneficially prepared thermoplastic cassava starch/palm wax blends incorporated with *Cymbopogan citratus* fibre (TCPS/PW/CCF) bio-composites at different CCF concentrations of 0, 10, 20, 30, 40, 50 and 60 wt%. In contrast, palm wax loading remained constant at 5 wt% concentration using the hot moulding compression method. TCPS/PW/CCF bio-composites were characterised in the present paper via their physical and impact properties. The addition of CCF significantly improved impact strength by 50.65% until 50 wt% CCF loading. Furthermore, it was observed that the inclusion of CCF resulted in a little decrement in biocomposite solubility compared to neat TPCS/PW biocomposite from 28.68% to 16.76%. Water absorption showed higher water resistance in the composites incorporating 60 wt.% fibre loading. The TPCS/PW/CCF biocomposites with different fibre contents had 11.04–5.65% moisture content, which was lower than the control biocomposite. The thickness of all samples decreased gradually with increasing fibre content. Overall, these findings provide evidence that CCF waste can be utilised as a high-quality filler in biocomposites due to its diverse characteristics, including improving the properties of biocomposites and strengthening their structural integrity.

**Keywords:** starch; *Cymbopogan citratus* fibre (CCF); thermoplastic cassava starch

## **1. Introduction**

Over the past two decades, demand for agriculturally based products has grown due to environmental concerns and the awareness that petroleum resources are limited [1,2]. Polymers derived from biological sources have often been raised as the most promising solutions since they are renewable and biodegradable [3–6]. The development of organic, renewable, and biodegradable materials has been driven by the escalating amount of plastic trash that threatens the environment [7]. Sustainable development and the absence of damaging carbon emissions during processing and after destruction are desirable properties of

**Citation:** Kamaruddin, Z.H.; Jumaidin, R.; Kamaruddin, Z.H.; Asyraf, M.R.M.; Razman, M.R.; Khan, T. Effect of *Cymbopogan citratus* Fibre on Physical and Impact Properties of Thermoplastic Cassava Starch/Palm Wax Composites. *Polymers* **2023**, *15*, 2364. https://doi.org/10.3390/ polym15102364

Academic Editor: Alexey Iordanskii

Received: 15 February 2023 Revised: 5 May 2023 Accepted: 14 May 2023 Published: 18 May 2023

**Copyright:** © 2023 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

prospective materials [8–10]. Excess plastic garbage is too much for recycling programmes because it is difficult to distinguish between polymers and eco-friendly plastics [11,12]. It is urgent to provide eco-friendly packaging items at comparable prices to preserve our planet, since synthetic plastic packaging pollution has exacerbated the issue of environmental destruction. In order to solve this recurring issue brought on by non-biodegradable plastics, natural biopolymers such as starch-based composites are being investigated as potential alternatives to conventional plastics [13–15].

Similar traits to conventional synthetic fibre-reinforced composites can be seen in natural fibre-reinforced composites. Due to considerable environmental concerns, they have been considered an alternative to organic and inorganic fillers and fibres [16,17]. These natural-source polymers can be found in various plant parts, such as grass [18], stalks [19], and leaves [20]. Hemicellulose, cellulose, and lignin are all examples of natural fibres [21]. Cellulose and hemicellulose components are all hydrophilic, whereas lignin is slightly hydrophobic [22,23]. Natural fibres are desirable filler materials for polymer composites because of their low manufacturing costs, particular acceptable characteristics, energy consumption, and environmentally friendly nature [24–26]. Natural fibres were incorporated to reinforce starch-based polymers, significantly impacting the materials' physical and impact qualities. These were primarily attributable to cellulose and starch's close structural resemblance and connection [27,28].

Starch, a form of carbohydrate, is one of the natural polymers that has received attention for use in the production of biodegradable polymers. Starch is generally composed of linear and branched chains of glucose molecules known as amylose and amylopectin, respectively [29]. Starch is combined with a plasticiser and fibre reinforcement and then subjected to high-temperature compression moulding to produce thermoplastic starch. Native starch can be transformed into a thermoplastic material by adding plasticising agents such as sorbitol, xylitol and glycerol. Incorporating plasticisers results in reduced intermolecular forces, increased mobility of polymer chains, and decreased glass transition temperature [30]. In our prior studies, we discovered that combining thermoplastic starch and palm wax had better properties than biopolymer blends [31,32]. As well, palm wax is considered an excellent matrix because it may reduce starch's hydrophobicity, improving processability.

The use of starch in both food and non-food products has been widespread. Since starch-based composites have poor mechanical efficiency, particularly impact strength, improving their qualities is a big challenge [33,34]. Numerous blending and compositing techniques have been devised to enhance these mechanical properties, such as mixing with other polymers or strengthening natural fillers [35–37]. Various additional agents support the interaction between starch and other components. In order to build polymer blend composites and biocomposites, filler reinforcement added to starch-based polymers has recently received more attention [38,39]. It was established that using natural fillers to reinforce materials with specific capabilities and add new features is a successful strategy. Although pure thermoplastic starch has many benefits, it also has some drawbacks that restrict the range of its potential uses. These drawbacks include limited water barriers, mechanical strength, and long-term stability [40–42]. Thus, modifications are typically needed to make thermoplastic starch useful in actual applications. The utilization of natural fibers as reinforcement for thermoplastic starch is an intriguing approach that could be utilized to solve these shortcomings. The mechanical properties of natural fibers, when combined with thermoplastic starch, are obviously improved due to the chemical similarity of starch and plant fibers [43]. In addition, thermoplastic starch can be used in combination with other natural polymers to reduce the downsides of this biopolymer while maintaining the material's biodegradability [15]. Results showed that adding palm wax significantly improved the mechanical characteristics of a cassava starch/palm wax thermoplastic blend [31] and improved the biopolymer's performance. Additionally, it contributed to the biopolymer's distinctiveness as an edible food coating and food packaging material.

Lemongrass, also called *Cymbopogan citratus,* is one of the Poaceae family's aromatic plants. These natural crops are grown in the tropical and semitropical areas of Africa, Asia, South America, and India [32]. In Malaysia, *Cymbopogon citratus* fibre (CCF) is becoming more well-liked as a natural resource that can help create ecologically friendly resources. The leaves of *Cymbopogan citratus* grow directly from the earth and have short rhizomes. All parts of the *Cymbopogan citratus* plant, including the leaves, are abundantly produced since the plant is simple to grow. The leaves have a glabrous surface and a greenish interior and can reach a length and width of about 50 and 1.5 cm, respectively [44]. The leaves of the *Cymbopogon citratus* plant are thrown away as trash in Malaysia, where the stalks are most frequently utilised as a constituent to flavour cuisine [44]. Due to the significant volume of garbage created, various trash reduction attempts have been conducted by preventing disposal and reusing waste. CCF has recently been employed as an adsorbent from aqueous solutions to remove methylene blue dye, as a raw material for making paper and pulp, and broadly in medicinal activities [45,46]. In addition, CCF is one of the substitute materials used to manufacture polymer composites as a reinforcing agent. The fibre of *Cymbopogan citratus* contains a significant quantity of cellulose, making it a source of cellulose fibre, according to research [47].

Although research has been reported on using Cymbopogan species from China in synthetic polymers, none was discovered on the characterisation of Cymbopogan citratus fibre from Malaysia in a thermoplastic cassava starch/palm wax matrix. As a result, the primary objective of this research topic is to investigate the effects of Malaysian-sourced Cymbopogan citratus fibre on the impact resistance and physical properties of thermoplastic cassava starch/palm wax composites. This study aims to produce decomposable materials that reduce pollution and are more eco-friendly.

### **2. Materials and Methodology**

#### *2.1. Materials*

CCF, with contents of cellulose (37.56%), hemicellulose (29.29%), lignin (11.14%), and ash (4.28%), was collected from a rural area located in Beranang, Selangor (West Malaysia). The extraction of CCF was performed according to our previous research [48] via the water-retting process. Antik Sempurna Sdn. Bhd., based in Selangor, Malaysia, provided the food-grade cassava starch used in this investigation. The palm wax in the commercial grade was provided by Green & Natural Industries Sdn. Bhd and the glycerol (99.5% purity) was purchased from QRec Chemicals Sdn, Selangor, Malaysia.

### *2.2. Sample Preparation of Biocomposites*

The thermoplastic cassava starch (TPCS) preparation was made by combining 5% palm wax with 65% starch and 30% glycerol. The mixture was then blended at room temperature for 5 min at 1200 rpm using a Panasonic Dry Mixer MX-GM1011. The materials were subjected to thermo-pressing at 150 ◦C for 30 min using a Malaysian-made Technopress-40HC-B Plastic Hydraulic Moulding Press under a weight of 10 tonnes to produce plates with a thickness of about 3 mm. Similar techniques were applied to create TPCS/PW/CCF composites. The matrix's characteristics were altered by introducing various CCF ratios ranging from 0 to 60 wt.%. The specimens were immediately put in a desiccator with silica gel to avoid unforeseen water absorption before conditioning.

#### *2.3. Density*

The ASTM D1895 [49] standard was followed to determine the biocomposite material density of the samples (10 mm × 10 mm × 3 mm), which were fabricated and dried for 24 h at 105 ◦C in the oven. The specimens were then positioned in a desiccator with silica gel in granulated form, and their weights and volumes were calculated using an electronic densimeter and weighing balance. Equation (1) was used to determine the density value.

$$\text{Density } (\text{g/cm}^3) = \frac{\text{Mass (g)}}{\text{Volume (cm}^3)} \tag{1}$$

#### *2.4. Moisture Content of Composite Materials*

Moisture content analysis was performed for five samples that were prepared and heated in the oven for 24 h at 105 ◦C. The initial weight (W<sup>i</sup> , grammes) and the final weight (W<sup>f</sup> , grammes) after heating are needed to calculate the moisture content [50]. Equation (2) was utilised to compute the moisture content.

$$\text{Moisture content (\%)} = \frac{\text{W}\_{\text{i}} - \text{W}\_{\text{f}}}{\text{W}\_{\text{i}}} \times 100\tag{2}$$

#### *2.5. Water Absorption*

Five samples (10 mm × 10 mm × 3 mm) were dried for 24 h at 105 ± 2 ◦C in an air-circulating oven to remove existing moisture. The test samples were fully submerged in water for 2 h at ambient temperature (23 ± 1 ◦C). The initial weight (W<sup>i</sup> ) and final immersion weight (W<sup>f</sup> ) were calculated for determining the water absorption rate using Equation (3).

$$\text{Water absorption (\%)} = \frac{\text{W}\_{\text{i}} - \text{W}\_{\text{f}}}{\text{W}\_{\text{i}}} \times 100\tag{3}$$

#### *2.6. Thickness Swelling*

Analysis of the samples' swelling was conducted based on the amended procedure by Jawaid et al. [51]. Five samples with measurements of 10 mm × 10 mm × 3 mm were prepared for thickness swelling calculation. Specimens were then dried in a dry laboratory oven at 105 ◦C for 24 h. Each sample was measured for thickness before testing (T<sup>i</sup> ). After that, each sample was submerged in 30 mL of distilled water for 2 h at room temperature (23 ± 1 ◦C). During the immersion process, the starting and final thicknesses, T<sup>i</sup> and T<sup>f</sup> , were measured using a Mitutoyo brand digital vernier calliper with a precision of 0.01 cm. This allowed for more accurate reading. Equation (4) summarizes the steps used to calculate the samples' thickness swelling ratio:

$$\text{Thickness swelling} \left( \% \right) = \frac{\text{T}\_{\text{i}} - \text{T}\_{\text{f}}}{\text{T}\_{\text{i}}} \times 100 \tag{4}$$

#### *2.7. Impact Testing of Composites Material*

Analysis via Izod impact tests was performed using ASTM D256 [52] with 50 ± 5% relative humidity and temperature of 23 ± 1 ◦C, and five replicates of each sample with dimensions of 13 mm (W) × 3 mm (T) × 60 mm (L) were set up. Victor Equipment Resources Sdn. Bhd. (Subang Jaya, Malaysia) provided testing using a digital pendulum impact tester. Before testing, all samples were preconditioned for two days and then processed at 53% RH. The results of the impact properties were obtained by taking the average of the data using Equation (5).

$$\text{Impact strength} = \text{Impact energy (l)} / \text{area (mm}^2) \tag{5}$$

#### *2.8. Water Solubility of Composite Materials*

The procedure described by Zhang et al. [53] was used to measure the samples' water solubility with a few minor adjustments. Five samples (10 mm × 10 mm × 3 mm) were cut and dried in a dry laboratory oven for 24 h at 105 ◦C 2 and the initial dry matter of each sample was noted as W<sup>i</sup> . Each sample was immersed in 30 mL of distilled water while stirring vigorously. The undissolved portions of the sample were then taken out of the laboratory cup after a 24-h immersion, and any remaining water on the sample's surface was wiped away using filter paper. The samples were then once again dried for 24 h at 105 ◦C to produce the final dried sample, indicated as W<sup>f</sup> . Equation (6) was used to calculate the samples' water solubility.

$$\text{Water solubility (\%)} = \frac{\text{W}\_{\text{i}} - \text{W}\_{\text{f}}}{\text{W}\_{\text{i}}} \times 100\tag{6}$$

#### *2.9. Statistical Analysis of Composites Material*

Analysis of variance (ANOVA) was used for statistical analysis of the experimental results in SPSS. Duncan's test was used to compare the means with a significance level of 0.05 (*p* < 0.05).

#### **3. Results and Discussions**

### *3.1. Density*

Biocomposites of TPCS/PW/CCF were produced by varying the CCF loading from 0% to 60% of fibre. There was a slightly significant difference in TPCS/PW/CCF biocomposite density. However, when fibre was introduced, it was discovered that the density value of biocomposite reduced as the fibre loading increased, as illustrated in Figure 1, ranging from 1.31 <sup>±</sup> 0.01 to 1.22 <sup>±</sup> 0.01 g/cm<sup>3</sup> . This finding might be related to the low density of the added CCF, which is known to be 0.25 <sup>±</sup> 0.002 g/cm<sup>3</sup> , relative to the TPCS/PW control sample (CCF0), which had a density value of 1.31 <sup>±</sup> 0.01 g/cm<sup>3</sup> . This affected the overall density of the bio-composite that was produced. Considering the density value of TPCS/PW/CCF, it can be noticed that the density of samples decreased by 6.87%. This specifies that the decreasing trend in density value might be associated with the reduction in the proportion of biocomposite mass when CCF is added, as the percentage of the TPCS/PW matrix is reduced while the volume remains constant [28]. As a result, reduction in biocomposite mass had a corresponding effect on the density values of the samples. This finding agreed with a previous study on developing Dioscorea hispida fibre-reinforced Dioscorea hispida starch, where the biocomposites' density values decreased as fibre loading was added [29]. Meanwhile, a similar result was reported when developing wood apple shell-reinforced epoxy composites. The density of the composites decreasing as the amount of filler used increased might be due to the lighter density of the filler material [54]. In terms of practicality, the decrease in bio-composite density when fibre loading is increased makes a substantial contribution by providing lightweight materials [55]. *Polymers* **2023**, *14*, x FOR PEER REVIEW 6 of 15

**Figure 1.** Density of TPCS/PW/CCF biocomposites. **Figure 1.** Density of TPCS/PW/CCF biocomposites.

*3.2. Moisture Content*

of accessible hydroxyl groups.

hydrophobicity, which improves processability [31,60].

Thermoplastic starch exhibits a high moisture content, which is caused by the inherent hydrophilic nature of native cassava starch, and results in poor moisture barrier

Consequently, the biocomposites have increased resistance to moisture. This result agrees with the previous study's finding on the effect on moisture content of cassava starch foam with the addition of natural fibre and chitosan [57]. Likewise, Soykeabkaew et al. [58] reported that similar results occurwhen jute and flax are added to tapioca starch foam, decreasing the amount of moisture content. This could be related to the hydrophobic nature of CCF, especially when compared to the CCF 0% sample. Although all-natural fibre is classified as hydrophilic, it is not as hygroscopic as TPCS. This might be due to the presence of lignin and wax in CCF composition, which makes it comparatively hydrophobic and gives it better water-resistant properties than TPCS [59]. In contrast, palm wax is considered a superior matrix since it can reduce starch

the inclusion of components that are either hydrophilic or hydrophobic. This study investigated the effect of Cymbopogan citratus fibre on thermoplastic cassava starch/palm wax composites' moisture content. The moisture content of the thermoplastic starch/palm wax composite without *Cymbopogan citratus* fibre was the highest, with absorption rates of 11.04%, which might be due to the strong hydrogen interactions between the hydroxyl groups in starch and the free water molecules [56]. The TPCS/PW/CCF biocomposites with different fibre contents had moisture contents of 11.04–5.65%, which was lower than that of the control biocomposite sample (11.04%), as illustrated in Figure 2. However, the increase in CCF content from 10 to 60 wt.% resulted in a slight reduction in the moisture content of the TPCS/PW/CCF samples. This could be explained by the strong hydrogen bonds between the Cymbopogan citratus fibre and the starch matrix, reducing the number

#### *3.2. Moisture Content*

Thermoplastic starch exhibits a high moisture content, which is caused by the inherent hydrophilic nature of native cassava starch, and results in poor moisture barrier capabilities. In spite of this, thermoplastic starch's resistance to water can be increased by the inclusion of components that are either hydrophilic or hydrophobic. This study investigated the effect of Cymbopogan citratus fibre on thermoplastic cassava starch/palm wax composites' moisture content. The moisture content of the thermoplastic starch/palm wax composite without *Cymbopogan citratus* fibre was the highest, with absorption rates of 11.04%, which might be due to the strong hydrogen interactions between the hydroxyl groups in starch and the free water molecules [56]. The TPCS/PW/CCF biocomposites with different fibre contents had moisture contents of 11.04–5.65%, which was lower than that of the control biocomposite sample (11.04%), as illustrated in Figure 2. However, the increase in CCF content from 10 to 60 wt.% resulted in a slight reduction in the moisture content of the TPCS/PW/CCF samples. This could be explained by the strong hydrogen bonds between the Cymbopogan citratus fibre and the starch matrix, reducing the number of accessible hydroxyl groups. *Polymers* **2023**, *14*, x FOR PEER REVIEW 7 of 15

**Figure 2.** Moisture content of TPCS/PW/CCF biocomposites. **Figure 2.** Moisture content of TPCS/PW/CCF biocomposites.

*3.3. Water Absorption of Composites* Water absorption is crucial for many applications of TPS products. The results of the water absorption tests performed on TPCS/PPW/CCF biocomposites are shown in Figure 3. In general, water absorption followed the same trend as moisture content, which exhibited decreasing values when fibre content is increased. In this study, the neat TPCS/PW biocomposite showed the highest water absorption at 33.15%. However, the absorption rate appeared to be reduced with CCF reinforcement, which reached 10.90% with 60 wt.% loading. The ability of fibres to produce strong interfacial bonding, which prevents water from penetrating through the matrix, may play a role in the decrease in absorption rate in the presence of CCF [61]. In addition, the presence of lignin as the main component, in Consequently, the biocomposites have increased resistance to moisture. This result agrees with the previous study's finding on the effect on moisture content of cassava starch foam with the addition of natural fibre and chitosan [57]. Likewise, Soykeabkaew et al. [58] reported that similar results occurwhen jute and flax are added to tapioca starch foam, decreasing the amount of moisture content. This could be related to the hydrophobic nature of CCF, especially when compared to the CCF 0% sample. Although all-natural fibre is classified as hydrophilic, it is not as hygroscopic as TPCS. This might be due to the presence of lignin and wax in CCF composition, which makes it comparatively hydrophobic and gives it better water-resistant properties than TPCS [59]. In contrast, palm wax is considered a superior matrix since it can reduce starch hydrophobicity, which improves processability [31,60].

#### conjunction with wax and fatty substances, in fibre enhances the ability of composites to *3.3. Water Absorption of Composites*

withstand the effects of water [61,62]. A prior study also reported that the gradual loading of sugar palm fibre into TPS-based hybrid composites resulted in a considerable reduction in the water absorption rate [63]. Meanwhile, a previous study by Sarifuddin et al. [64] on developing thermoplastic sago starch reinforced kenaf core fibre biocomposites reported low water absorption when adding kenaf core fibre into the samples. This was associated Water absorption is crucial for many applications of TPS products. The results of the water absorption tests performed on TPCS/PPW/CCF biocomposites are shown in Figure 3. In general, water absorption followed the same trend as moisture content, which exhibited decreasing values when fibre content is increased. In this study, the neat

with the fibre having a less hydrophilic nature than the starch. The higher the fibre loading, the lower the amount of absorbed water. The strong hydrogen bonding between

biocomposite is less water-resistant than the control CCF samples.

TPCS/PW biocomposite showed the highest water absorption at 33.15%. However, the absorption rate appeared to be reduced with CCF reinforcement, which reached 10.90% with 60 wt.% loading. *Polymers* **2023**, *14*, x FOR PEER REVIEW 8 of 15

**Figure 3.** TPCS/PW/CCF biocomposites' water absorption. **Figure 3.** TPCS/PW/CCF biocomposites' water absorption.

*3.4. Thickness Swelling* A material's dimensional stability can be evaluated by determining the degree to which it swells or contracts as a direct result of the movement of moisture in the material. The thickness swelling ratio of TPCS/PW/CCF composites was evaluated to determine changes in TPCS/PW dimensional stability following the introduction of CCF. Figure 4 illustrates the thickness swelling percentage for TPCS/PW with the inclusion of CCF after immersion for 2 h. The thickness of all samples decreased gradually with increasing fibre content. In comparison to the TPCS/PW control sample, the sample containing 60% CCF content exhibited the lowest water uptake and thickness swelling after 2 h of immersion, decreasing by 54.24%. This finding can be attributed to the presence of fibre in the composites, which possesses a more rigid structure than starch, providing higher dimensional stability to the composites [65]. This study's findings agree with those of a prior investigation into composites made of cassava starch and green coir fibres [66]. Additionally, similar findings were reported when developing a cassava/sugar palm fibre-The ability of fibres to produce strong interfacial bonding, which prevents water from penetrating through the matrix, may play a role in the decrease in absorption rate in the presence of CCF [61]. In addition, the presence of lignin as the main component, in conjunction with wax and fatty substances, in fibre enhances the ability of composites to withstand the effects of water [61,62]. A prior study also reported that the gradual loading of sugar palm fibre into TPS-based hybrid composites resulted in a considerable reduction in the water absorption rate [63]. Meanwhile, a previous study by Sarifuddin et al. [64] on developing thermoplastic sago starch reinforced kenaf core fibre biocomposites reported low water absorption when adding kenaf core fibre into the samples. This was associated with the fibre having a less hydrophilic nature than the starch. The higher the fibre loading, the lower the amount of absorbed water. The strong hydrogen bonding between the matrix phase and the fibre phase can be attributed to the overall lower water absorption of the material. However, it is clear that the CCF-reinforced starch biocomposite is less water-resistant than the control CCF samples.

#### reinforced cassava starch hybrid. In this study, the incorporation of sugar palm fibre made *3.4. Thickness Swelling*

the sample swell less and influenced the interaction between the polymer chains, which could reduce the polymer's ability to swell. In each sample, increasing the fibre concentration resulted in a decrease in swelling behaviour. The sample that contained a low fibre loading was more swollen than the ones that contained a high fibre loading [67]. A material's dimensional stability can be evaluated by determining the degree to which it swells or contracts as a direct result of the movement of moisture in the material. The thickness swelling ratio of TPCS/PW/CCF composites was evaluated to determine changes in TPCS/PW dimensional stability following the introduction of CCF. Figure 4 illustrates the thickness swelling percentage for TPCS/PW with the inclusion of CCF after immersion for 2 h. The thickness of all samples decreased gradually with increasing fibre content. In comparison to the TPCS/PW control sample, the sample containing 60% CCF content exhibited the lowest water uptake and thickness swelling after 2 h of immersion, decreasing by 54.24%. This finding can be attributed to the presence of fibre in the composites, which possesses a more rigid structure than starch, providing higher dimensional stability to the composites [65]. This study's findings agree with those of a prior investigation into composites made of cassava starch and green coir fibres [66]. Additionally, similar findings were reported when developing a cassava/sugar palm fibre-reinforced cassava starch hybrid. In this study, the incorporation of sugar palm fibre made the sample swell less and

influenced the interaction between the polymer chains, which could reduce the polymer's ability to swell. In each sample, increasing the fibre concentration resulted in a decrease in swelling behaviour. The sample that contained a low fibre loading was more swollen than the ones that contained a high fibre loading [67]. *Polymers* **2023**, *14*, x FOR PEER REVIEW 9 of 15

**Figure 4.** Thickness swelling of TPCS/PW/CCF biocomposites. **Figure 4.** Thickness swelling of TPCS/PW/CCF biocomposites.

#### *3.5. Impact Testing of Composites 3.5. Impact Testing of Composites*

The impact strength of the biocomposites was also determined to measure the material's impact resistance, indicating the capability to endure an unexpectedly applied load. Figure 5 demonstrates the TPCS/PW/CCF's impact strength with different fibre contents. Analysis of variance (ANOVA) of the impact characteristics is presented in Table 1. This indicates that there is a statistically significant difference (*p* < 0.05) between the mean impact strength from one level of composites to another. In general, the findings of the impact test results demonstrate the highest impact strength at 50 wt.% fibre loading. Significant improvement (*p* < 0.05) was evidenced in the impact properties of the composite when incorporated with *Cymbopogan citratus* fibre. Failure of a fibre composite to withstand impact loading can be attributed to several other factors, including fibre pullout, matrix fracture, fibre breakage, and debonding of the fibres from the matrix [68]. It could be observed that the TPCS/PW/CCF biocomposite had significantly increased impact strength by 50.65% (*p* < 0.05) with increasing fibre content from (0 to 50 wt.%). This finding might be associated with the strong interfacial bonding at fibre–matrix interfaces, which improved energy absorption during impact loading [69]. However, impact resistance declined after increasing reinforcement from 50 to 60 wt.%. The reduction in impact strength at higher fibre contents might be attributed to the higher rigidity of the materials, which led to more brittle properties and, as a result, a decreased capacity to absorb the impact energy [31]. In our prior studies, we discovered that the results revealed that elongation at break values increased from 40 to 50 wt.% as a result of the addition of *Cymbopogan citratus* fibre loading; however, an increase in the amount of *Cymbopogan citratus* fibre up to 60 wt.% caused a decrease in elongation at break values. Overall, the inclusion of *Cymbopogan citratus* fibre content resulted in a decrease in the molecular mobility of the TPCS matrix, which produced biocomposite materials with a greater degree of rigidity [48]. A similar result was observed in a study on the development of date palmreinforced epoxy composites. A significant improvement in impact strength was found at 50% date palm fibre reinforced epoxy composites in comparison to loadings of 40% and 60%, which can be primarily attributed to better adhesion of the date palm fibre with the The impact strength of the biocomposites was also determined to measure the material's impact resistance, indicating the capability to endure an unexpectedly applied load. Figure 5 demonstrates the TPCS/PW/CCF's impact strength with different fibre contents. Analysis of variance (ANOVA) of the impact characteristics is presented in Table 1. This indicates that there is a statistically significant difference (*p* < 0.05) between the mean impact strength from one level of composites to another. In general, the findings of the impact test results demonstrate the highest impact strength at 50 wt.% fibre loading. Significant improvement (*p* < 0.05) was evidenced in the impact properties of the composite when incorporated with *Cymbopogan citratus* fibre. Failure of a fibre composite to withstand impact loading can be attributed to several other factors, including fibre pull-out, matrix fracture, fibre breakage, and debonding of the fibres from the matrix [68]. It could be observed that the TPCS/PW/CCF biocomposite had significantly increased impact strength by 50.65% (*p* < 0.05) with increasing fibre content from (0 to 50 wt.%). This finding might be associated with the strong interfacial bonding at fibre–matrix interfaces, which improved energy absorption during impact loading [69]. However, impact resistance declined after increasing reinforcement from 50 to 60 wt.%. The reduction in impact strength at higher fibre contents might be attributed to the higher rigidity of the materials, which led to more brittle properties and, as a result, a decreased capacity to absorb the impact energy [31]. In our prior studies, we discovered that the results revealed that elongation at break values increased from 40 to 50 wt.% as a result of the addition of *Cymbopogan citratus* fibre loading; however, an increase in the amount of *Cymbopogan citratus* fibre up to 60 wt.% caused a decrease in elongation at break values. Overall, the inclusion of *Cymbopogan citratus* fibre content resulted in a decrease in the molecular mobility of the TPCS matrix, which produced biocomposite materials with a greater degree of rigidity [48]. A similar result was observed in a study on the development of date palm-reinforced epoxy composites. A significant improvement in impact strength was found at 50% date palm fibre reinforced epoxy composites in comparison to loadings of 40% and 60%, which can be primarily attributed to better adhesion of the date palm fibre with the epoxy matrix in order to overcome the high impact stress/load. Furthermore, the increase in impact strength might be associated with

epoxy matrix in order to overcome the high impact stress/load. Furthermore, the increase in impact strength might be associated with enhanced stress capability, which would limit

enhanced stress capability, which would limit the contribution of fibre-related mechanisms such as fibre pull out [70]. *Polymers* **2023**, *14*, x FOR PEER REVIEW 10 of 15

**Figure 5.** Impact strength of TPCS/PW/CCF biocomposites. **Figure 5.** Impact strength of TPCS/PW/CCF biocomposites.

**Table 1.** Analysis of variance (ANOVA) summary of TPCS/PW/CCF composites. **Table 1.** Analysis of variance (ANOVA) summary of TPCS/PW/CCF composites.


Note: \* Significantly different at *p* < 0.05.

#### *3.6. Water Solubility of Composites 3.6. Water Solubility of Composites*

thickness swelling findings.

The solubility of samples in water is one of the essential properties worthy of consideration in many applications. Applications that require moisture and water loss protection must have low water solubility. Figure 6 shows the TPCS/PW/CCF biocomposites' water solubility. This solubility demonstrates the effect of water immersion with continuous stirring on the composite samples. The inclusion of CCF resulted in a little decrement in biocomposite solubility compared to neat TPCS/PW biocomposite, from 28.68% to 16.76%. This phenomenon could be attributed to the role of the fibre, which produces a network that tightly keeps composites together and prevents the dissolution of composites by lowering the solubility of the samples [55]. Additionally, the solubility of samples reduced with increasing fibre content, which might be due to fibre's capacity to resist water diffusion and enhance composite integrity by limiting water penetration, thus reducing sample solubility [63]. This observation agrees with previous findings on the development of a cornhusk/sugar palm fibre reinforced corn starch-based hybrid, which reported that adding sugar palm fibre decreased the solubility of the samples [61]. The result supports this study's water absorption, moisture content, and The solubility of samples in water is one of the essential properties worthy of consideration in many applications. Applications that require moisture and water loss protection must have low water solubility. Figure 6 shows the TPCS/PW/CCF biocomposites' water solubility. This solubility demonstrates the effect of water immersion with continuous stirring on the composite samples. The inclusion of CCF resulted in a little decrement in biocomposite solubility compared to neat TPCS/PW biocomposite, from 28.68% to 16.76%. This phenomenon could be attributed to the role of the fibre, which produces a network that tightly keeps composites together and prevents the dissolution of composites by lowering the solubility of the samples [55]. Additionally, the solubility of samples reduced with increasing fibre content, which might be due to fibre's capacity to resist water diffusion and enhance composite integrity by limiting water penetration, thus reducing sample solubility [63]. This observation agrees with previous findings on the development of a cornhusk/sugar palm fibre reinforced corn starch-based hybrid, which reported that adding sugar palm fibre decreased the solubility of the samples [61]. The result supports this study's water absorption, moisture content, and thickness swelling findings.

**Figure 6.** Water solubility of TPCS/PW/CCF biocomposites. **Figure 6.** Water solubility of TPCS/PW/CCF biocomposites.

#### **4. Conclusions 4. Conclusions**

Utilising hot pressing, a novel biocomposite made of thermoplastic cassava starch and varying concentrations of *Cymbopogan citratus* fibre (CCF) was created. Its water barrier behaviour and impact qualities were studied. The experimental results revealed that the inclusion of CCF resulted in a little decrement in biocomposite solubility compared to neat TPCS/PW biocomposite, from 28.68% to 16.76%, and a reduction in the absorption rate, which reached 10.90% with 60 wt.% loading. TPCS/PW/CCF biocomposites with different fibre contents had moisture contents of 11.04–5.65%, which was lower than the control biocomposite sample, and a slightly significant difference in biocomposite density. TPCS-reinforced CCF composites show promise as an alternative to non-environmentally friendly polymers and composites, with better water barrier and impact performance. Utilising hot pressing, a novel biocomposite made of thermoplastic cassava starch and varying concentrations of *Cymbopogan citratus* fibre (CCF) was created. Its water barrier behaviour and impact qualities were studied. The experimental results revealed that the inclusion of CCF resulted in a little decrement in biocomposite solubility compared to neat TPCS/PW biocomposite, from 28.68% to 16.76%, and a reduction in the absorption rate, which reached 10.90% with 60 wt.% loading. TPCS/PW/CCF biocomposites with different fibre contents had moisture contents of 11.04–5.65%, which was lower than the control biocomposite sample, and a slightly significant difference in biocomposite density. TPCS-reinforced CCF composites show promise as an alternative to non-environmentally friendly polymers and composites, with better water barrier and impact performance.

**Author Contributions:** Conceptualization, Z.H.K. (Zatil Hafila Kamaruddin) and R.J.; methodology, Z.H.K. (Zatil Hafila Kamaruddin); formal analysis, Z.H.K. (Zatil Hazrati Kamaruddin) and R.J.; investigation, Z.H.K. (Zatil Hazrati Kamaruddin); resources, M.R.M.A and M.R.R.; data curation, Z.H.K. (Zatil Hazrati Kamaruddin) and M.R.M.A.; writing—original draft preparation, Z.H.K. (Zatil Hazrati Kamaruddin); writing—review and editing, Z.H.K. (Zatil Hazrati Kamaruddin); visualization, Z.H.K. (Zatil Hafila Kamaruddin) and M.R.M.A.; supervision, R.J.; project administration, Z.H.K. (Zatil Hafila Kamaruddin); funding acquisition, M.R.M.A and T.K. All authors have read and agreed to the published version of the manuscript. **Author Contributions:** Conceptualization, Z.H.K. (Zatil Hafila Kamaruddin) and R.J.; methodology, Z.H.K. (Zatil Hafila Kamaruddin); formal analysis, Z.H.K. (Zatil Hazrati Kamaruddin) and R.J.; investigation, Z.H.K. (Zatil Hazrati Kamaruddin); resources, M.R.M.A. and M.R.R.; data curation, Z.H.K. (Zatil Hazrati Kamaruddin) and M.R.M.A.; writing—original draft preparation, Z.H.K. (Zatil Hazrati Kamaruddin); writing—review and editing, Z.H.K. (Zatil Hazrati Kamaruddin); visualization, Z.H.K. (Zatil Hafila Kamaruddin) and M.R.M.A.; supervision, R.J.; project administration, Z.H.K. (Zatil Hafila Kamaruddin); funding acquisition, M.R.M.A. and T.K. All authors have read and agreed to the published version of the manuscript.

**Funding:** The financial support was received from the Universiti Teknikal Malaysia Melaka and the Ministry of Higher Education Malaysia for financial aid through research grant PJP/2021/FTKMP/S01815 and to publish these project findings. Additionally, the paper has received funding from Universiti Teknologi Malaysia through the project "Characterizations of Hybrid Kenaf Fibre/Fibreglass Meshes Reinforced Thermoplastic ABS Composites for Future Use in Aircraft Radome Applications" under grant number PY/2022/03758—Q.J130000.3824.31J25. **Funding:** The financial support was received from the Universiti Teknikal Malaysia Melaka and the Ministry of Higher Education Malaysia for financial aid through research grant PJP/2021/FTKMP/S01815 and to publish these project findings. Additionally, the paper has received funding from Universiti Teknologi Malaysia through the project "Characterizations of Hybrid Kenaf Fibre/Fibreglass Meshes Reinforced Thermoplastic ABS Composites for Future Use in Aircraft Radome Applications" under grant number PY/2022/03758—Q.J130000.3824.31J25.

**Institutional Review Board Statement:** Not applicable. **Institutional Review Board Statement:** Not applicable.

**Data Availability Statement:** The data is available within the article. **Data Availability Statement:** The data is available within the article.

**Acknowledgments:** The authors would like to acknowledge the support of Universiti Teknikal Malaysia Melaka and the Ministry of Higher Education Malaysia for financial aid through research grant PJP/2021/FTKMP/S01815 and to publish these project findings. Additionally, the authors would like to express gratitude for the financial support from the Universiti Teknologi Malaysia through the project "Characterizations of Hybrid Kenaf Fibre/Fibreglass Meshes Reinforced Thermoplastic ABS Composites for Future Use in Aircraft Radome Applications" under grant number PY/2022/03758—Q.J130000.3824.31J25. In addition, thank you to the German Malaysian Institute for providing a scholarship award to the principal author of this project. **Acknowledgments:** The authors would like to acknowledge the support of Universiti Teknikal Malaysia Melaka and the Ministry of Higher Education Malaysia for financial aid through research grant PJP/2021/FTKMP/S01815 and to publish these project findings. Additionally, the authors would like to express gratitude for the financial support from the Universiti Teknologi Malaysia through the project "Characterizations of Hybrid Kenaf Fibre/Fibreglass Meshes Reinforced Thermoplastic ABS Composites for Future Use in Aircraft Radome Applications" under grant number PY/2022/03758—Q.J130000.3824.31J25. In addition, thank you to the German Malaysian Institute for providing a scholarship award to the principal author of this project.

**Conflicts of Interest:** The authors declare no conflict of interest.

## **References**


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## *Article* **An Experimental Study on Drilling Behavior of Silane-Treated Cotton/Bamboo Woven Hybrid Fiber Reinforced Epoxy Polymer Composites**

**Karthik Aruchamy <sup>1</sup> , Sathish Kumar Palaniappan 2,\* , Rajeshkumar Lakshminarasimhan <sup>3</sup> , Bhuvaneshwaran Mylsamy <sup>4</sup> , Satish Kumar Dharmalingam <sup>5</sup> , Nimel Sworna Ross <sup>5</sup> and Sampath Pavayee Subramani <sup>6</sup>**

	- Coimbatore 641407, Tamil Nadu, India

**Abstract:** Machining is considered to be an important post-manufacturing process. Evaluation of machinability of natural-fiber-reinforced composites is important owing to its wide application spectrum. Current experiments focus on the drilling parameters of cotton/bamboo woven fabric reinforced epoxy composites laminates using a solid twist drill. Composites were manufactured with 45 wt.% cotton/bamboo woven fabric in epoxy resin using a compression molding method. Drilling experiments were carried out in pillar-type drilling machine and the drilling characteristics, such as thrust force, were analyzed using four process parameters like spindle speed, feed rate, drill diameter, and silane-treated fabric. Drilling experiments were carried out using the Box–Behnken Experimental Design, and the recommended drilling characteristics were analyzed using quadratic models based on response surface methodology. It was observed from the results that the thrust force is low with small drill-bit diameter, higher cutting speed, and lower feed rate, according to the response surface analysis. Surface morphology of the drilled hole suggested that a better quality of hole can be obtained at lower feed rates.

**Keywords:** cotton/bamboo woven fabric; silane treatment; thrust force; SEM analysis; RSM; machining operation

## **1. Introduction**

The study of composite material from the last two decades that focuses on investigating natural-fiber-reinforced composite is an important branch of materials science [1–5]. Cotton, bamboo, bark, wood, pulp, cereal straw, bark, bagasse, corncobs, nut shells, and vegetable (e.g., coir, ramie, sisal, jute, flax, sun hemp, banana, and pineapple) are the examples of plant-based natural fibers [6–8]. The major microconstituents of these fibers are hemicelluloses, cellulose, wax, and lignin, along with a minimum percentage of extractives. The composition of fiber varies depending on their origin [9–11]. The fibers used in engineering applications mainly include synthetic fibers and natural fibers. Synthetic fibers such as carbon fiber and glass fiber have been widely used in different fields because of their lightweight and high strength, excellent mechanical properties, corrosion resistance,

**Citation:** Aruchamy, K.; Palaniappan, S.K.; Lakshminarasimhan, R.; Mylsamy, B.; Dharmalingam, S.K.; Ross, N.S.; Pavayee Subramani, S. An Experimental Study on Drilling Behavior of Silane-Treated Cotton/Bamboo Woven Hybrid Fiber Reinforced Epoxy Polymer Composites. *Polymers* **2023**, *15*, 3075. https://doi.org/10.3390/ polym15143075

Academic Editors: R.A. Ilyas, S.M. Sapuan, Emin Bayraktar, Shukur Abu Hassan, Nabil Hayeemasae, Khubab Shaker and Chenggao Li

Received: 15 May 2023 Revised: 10 July 2023 Accepted: 16 July 2023 Published: 18 July 2023

**Copyright:** © 2023 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

fatigue resistance, and so on. The above advantages can make synthetic fiber composites the main load-bearing components in engineering applications [12–14]. Also, natural fibers show several benefits compared to standard glass and carbon fibers. Natural fibers are of low cost, abundantly available, environment-friendly, easy to process, have low density, and possess reasonably good flexural and tensile modulus. By considering some key problems in synthetic fibers such as high cost, environmental problems during the preparation, and nonrenewable material, natural fibers overcome the above disadvantages and can be mixed with polymer matrix to create composites. Even if the strength of the fiber-reinforced composites is sometimes lower than the individual fiber itself, either hybridization or suitable treatments may eradicate this. However, the strength of composites, at most times, would be less than the individual fiber itself. Natural fibers with biodegradable and renewable properties make composting and incineration easier than all other synthetic and industrial fibers. In addition, the fibers contain stored atmospheric carbon dioxide and have a minimum embodied energy than industrially produced glass fibers [15–17].

Even though natural fibers have the possibility to complement synthetic fibers in polymer composite materials [15,16], natural fibers have limitations in terms of mechanical performance and absorption of humidity [17,18]. Regardless of whether thermoplastic or thermoset polymers are utilized as the matrix material, the same limitations exist. According to previous studies [4,19], the chemical incompatibility between the hydrophilic lignocellulosic molecules of the natural fiber and the hydrophobic thermoplastic molecules is the key factor in limiting the mechanical characteristics of natural-fiber-reinforced thermoplastic composites. Because of this incompatibility, achieving good fiber matrix interface bonding is challenging, resulting in inadequate load shift between the reinforced fibers and matrix. In order to obtain the requisite compatibility between the fiber and matrix, various solutions have been suggested, including chemical treatments of lignocelluloses fibers and the use of compatibilizers. Various authors stated that the different chemical processes can be used to enhance the lignocellulosic fiber compatibility with polymer matrices. The use of alkali treatment, isocyanate treatment, acetylation, permanganate treatment, peroxide treatment, benzoylation, and silane treatment were all studied. Treatments such as alkaline and silane were frequently reported among the various approaches [18,19].

The fibers were soaked in alkaline solution, most often NaOH, for a specified period of time during the alkaline treatment. It was expected to improve the mechanical bonding by enhancing the roughness of fiber surface. It increases the crystallinity of the fiber which enhances the possibility of chemical interaction with the matrix [20–22]. It also exposes more reactive functional groups of cellulose to be bonded with a polymer matrix. Alkaline treatment effects were studied by some authors on kenaf fiber [20]. Cleaning the fiber surface with an alkaline treatment using a 6% NaOH solution was found to be effective. On the other hand, a 9% NaOH solution was found to degrade the fiber surface and decreased its strength. The fibers are normally soaked in silane solution diluted in water/alcohol. As a result, silane hydrolyses itself into silanol from alcohol that contains water. The silanol combines with the OH groups of cellulose in natural fibers to generate stability in covalent linkages with cell walls that assimilate onto the fiber surface [21]. Silane treatment increases the cross-linking in the connected regions between fiber and the matrix, allowing for better interaction between the fiber and matrix [22]. Mechanical properties such as tensile strength and modulus of flax fiber reinforced with epoxy composite were found to increase, as stated by Van de Weyenburg et al. [23]. Results showed that improvement in the tensile strength and modulus was roughly 30% and 45%, respectively, when 40 vol.% of treated flax fibers were utilized as a reinforcement in epoxy matrix. This improvement was due to the combined chemical treatment of flax fibers, i.e., by using a combination of sodium hydroxide and silane coupling agent chemicals.

Natural-fiber-reinforced polymer composites are more difficult to machine because of lack of machining records and a higher level of parameter complexity [24]. Occurrence of voids, delamination, matrix, and fiber debonding are some of the examples for manufacturing and machining faults. Despite the fact that the majority of fiber-based hybrid

composites are made to actual shape by using the machining techniques such as drilling, grinding, and milling, which are completely unavoidable at the time of assembly process, drilling is considered the most important machining techniques in the overall process. When compared to drilling of conventional metals and alloys, the drilling procedure for fiber-reinforced composite samples is more complex because of the anisotropic nature of samples. Due to the presence of insoluble constituents in fiber-reinforced composites, the isotropy cannot be defined as in conventional metals and alloys, which makes them more anisotropic materials. This complexity is considered to be a significant one in the drilling of fiber-reinforced composites. It finds its application in structural automobile components, kitchen cabinets, doors, furniture, electronic components, and so on. The end product is assembled with sliced composite panels to form the desired shape and dimensions, and attaching them with bolts and screws in all of the above applications. As a result, before placing it into use, drilling is performed to complete the assembly process. Hence, drilling is considered to be extremely important, because a fault in the structural integrity of the drilling process influences the quality of the final product [25].

Drilling experiments in fiber-based polymer composites have sufficiently provided an indication of the importance and research direction. On a coir–polyester composite, the influence of drill diameter, speed, and feed on torque, thrust force, and tool wear was investigated. Drilling studies on roselle–sisal hybrid composites were performed and optimized using an artificial neural network model [26,27]. Few experimenters investigated initially with peel-up delamination and later with push-down delamination for hemp/flaxfiber-based composites and also discovered that the most important elements to consider throughout the drilling process are feed rate and cutting speed [28]. In the drilling of sisal polypropylene composites, drill geometry was determined to be the most influential element [29–31]. Drilling was performed on a cotton/bamboo (CB) woven fabric reinforced composite, and thrust force was calculated for various combinations of cutting speed, feed rate, and drill shape. The cutting parameters, as well as the input parameters, were optimized using a Box–Behnken Experimental Design (BBED). When the drill diameter was increased, the results revealed that the feed rate increases with the decrease in cutting speed. Drill diameter and rate of feed influence the increment of thrust force, according to Srinivasan et al. [32]. Using BBED and analysis of variance (ANOVA), the effect of thrust force was investigated, and they observed that low feed rate, high spindle speed, and smaller drill diameter resulted in the lowest thrust force.

In spite of various studies on drilling behavior of natural fiber composites, only few studies discuss the machining behavior of treated woven fiber reinforced composites. In the current study, silane-treated CB reinforced hybrid composites were analyzed for their machinability behavior through drilling process. Cutting thrust force was evaluated with various feed rates, cutting speeds, and drill geometries. The optimum parameter was determined in order to achieve the lowest cutting thrust force by using the response surface methodology (RSM) technique. The research experiments were controlled by BBED. The data were analyzed using 3D plots and influence graphs. The drilled surface morphology was examined using a scanning electron microscopy (SEM) and the optimal parameter settings for obtaining a good quality hole and better machinability were recommended.

### **2. Materials and Methodology**

## *2.1. Fabrication of CB Woven Fabric*

In this paper, the CB woven fabric composites were manufactured with cotton yarn and bamboo yarn packed with corresponding warp and weft directions. In this study, cotton yarn, bamboo yarn, and CB woven fabric were collected from Pallava Spinning Mills (Pvt.) Limited, Erode District, Tamil Nadu, India, and Ganapathi Chettiar Tex, Tirupur District, Tamil Nadu, India, with good-quality yarn that has bamboo fiber length of 36 mm, fiber fineness of 1.52 dtex, linear density of 0.155 tex, moisture regain value of 11.42%, tenacity of 22.84 g/tex, and elongation of 21.2% [33]. The yarn and fabric particulars are given in Tables 1 and 2, respectively.


**Table 1.** Particulars of yarns.

**Table 2.** Particulars of CB woven fabric [30,31].


#### *2.2. Silane Treatment of CB Fabrics*

CB fabrics were immersed in the 1.5% concentrations of 3-aminopropyl trimethoxy silanes solution (procured from M/s. Covai seenu, Coimbatore, Tamil Nadu, India) with water and acetone (50/50 volume) for 2 h. From various screening experiments and based on the literature, the concentration of silane was taken as 1.5% [23]. The fibers were then soaked in 1.5% silane mixed with 2% acetic acid in the form of acetic anhydride for pH adjustment between 4.5 and 5, and sundried for 24 h. Then, the fibers were oven dried at 80 ◦C for 4 h to maintain their stiffness. Silane coupling agents were expected to effectively modify the natural fiber–polymer matrix interface and increase their interfacial strength.

#### *2.3. Fabrication of Composite Laminates*

Composite laminates were manufactured using a compression molding technique. Initially, wax was applied to the surface of the mold for easy removal, and the size of the mold measured 270 mm × 270 mm × 3 mm. The CB woven fabrics (2 layers each) were laid and arranged, layer after layer, with epoxy adhesive to keep woven fabrics composed. For this experiment, Araldite LY 556 epoxy resin and HY951 hardener (procured from M/s. Covai seenu, Coimbatore, Tamil Nadu, India) were used. To assist gradual and homogeneous curing of the resin, a 10:1 mix of resin and hardener was used. A composite plate was created by stacking alternate layers of resin and fabric. It was then compressed at 30 MPa for 1 h at 160 ◦C. Although a greater temperature will speed up the process of curing, the natural fiber's existence limited high-temperature processing. A consistent weight fraction of 45% was used to make the composite laminates. The composite included isotropic character, which is the most important requirement in various commercial applications and is the only way to ensure uniform load-carrying capability of the material.

#### *2.4. Drilling Process Parameters and Optimization*

Drilling operations were performed in a pillar-type drilling machine (*Make: Kirloskar, Pune, India*) using HSS drill bits with pointed ends. All the operations were carried out at ambient atmospheric conditions. In this experimental work, data acquisition systems were used along with a drill dynamometer to obtain thrust force signals. Forces developed on the drill bit axis or spindle axis during the machining process are termed as thrust force, which induces damage during the drilling of silane-treated woven CB composite laminates, as shown in Figure 1.

the drill bit axis or spindle axis during the machining process are termed as thrust force, which induces damage during the drilling of silane-treated woven CB composite lami-

**Figure 1.** Drilling of CB composite laminate. **Figure 1.** Drilling of CB composite laminate.

nates, as shown in Figure 1.

The hole-making process in composite materials is influenced by a number of variables such as spindle speed, drill diameter, drill type feed rate, drill point angle drill material, type of matrix, type of fiber reinforcement, and fiber volume fraction. These drilling parameters define and influence the thrust force, surface quality of the drilled hole, delamination near the drilled hole, and other geometrical factors such as roundness and eccentricity [34]. Many investigations have discovered that the major influencing parameters in the drilling of fiber-reinforced composite panels include feed rate [35], drill diameter [36,37], and cutting speed [31]. Drilling operations were carried out by changing the drill bit rotation speed (rpm), drill diameter (mm), and rate of feed (mm/min). Table 3 shows the various process parameters used for the drilling of treated CB fabric woven The hole-making process in composite materials is influenced by a number of variables such as spindle speed, drill diameter, drill type feed rate, drill point angle drill material, type of matrix, type of fiber reinforcement, and fiber volume fraction. These drilling parameters define and influence the thrust force, surface quality of the drilled hole, delamination near the drilled hole, and other geometrical factors such as roundness and eccentricity [34]. Many investigations have discovered that the major influencing parameters in the drilling of fiber-reinforced composite panels include feed rate [35], drill diameter [36,37], and cutting speed [31]. Drilling operations were carried out by changing the drill bit rotation speed (rpm), drill diameter (mm), and rate of feed (mm/min). Table 3 shows the various process parameters used for the drilling of treated CB fabric woven hybrid composites.

A scanning electron microscope (model: VEGA 3 TE SCAN) was used to study the **Table 3.** Drilling process parameters [30,31].

hybrid composites.


**S. No. Process Parameter Unit Upper Limit Lower Limit** 1 Cutting speed (n) rpm 3000 1000 2 Drill diameter (d) mm 12 6 3 Feed rate (f) mm/min 60 20 A scanning electron microscope (model: VEGA 3 TE SCAN) was used to study the surface topography of the drilled holes in silane-treated CB composite laminates. A thin layer of sputtered gold was coated on the surface of all the tested composite specimens to improve electrical conductivity and to obtain better images.

RSM is the collection of mathematical and statistical data to develop modeling and analyses of engineering problems [38]. Nowadays, RSM is used in the design of experiments (DOE). Also, RSM helps in identifying the correlation between input parameters to obtain the response surface. Box and Draper [39] developed a model to suit physical experiments and later for machining problems. Further, RSM also has its objective to later RSM is the collection of mathematical and statistical data to develop modeling and analyses of engineering problems [38]. Nowadays, RSM is used in the design of experiments (DOE). Also, RSM helps in identifying the correlation between input parameters to obtain the response surface. Box and Draper [39] developed a model to suit physical experiments and later for machining problems. Further, RSM also has its objective to later mine the optimization of process parameters.

mine the optimization of process parameters. Douglas C. Montgomery [40] also developed a model relating the process variables with output response, which was utilized for process prediction and control. The experiment domain's boundary must be investigated in RSM. The process parameters for this analysis were feed rate, cutting speed, and diameter of the drill, and BBED was utilized to formulate the design of experiments. BBED, an alternate way to central composite design (CCD), is the common statistical tool employed to frame the number of experiments. This approach was used to analyze experiments comprising three levels of design. The BBED

can be rotated, and it is said to have no fractional factorial design, which makes it easy to understand the results [41]. In RSM, the general correlation of the control parameters and their responses is illustrated in Equation (1): where Y is the response (thrust force), β0, β1, β2,…β<sup>i</sup> are the regression coefficients, x1, x2,…x<sup>i</sup> denote the predictor variable, namely, feed rate, diameter, and cutting speed, and ε shows the error occurring in the thin model. Generally, in RSM, quadratic response func-

*Polymers* **2023**, *15*, x FOR PEER REVIEW 6 of 15

their responses is illustrated in Equation (1):

$$\mathbf{Y} = \boldsymbol{\beta}\_0 + \boldsymbol{\beta}\_1 \mathbf{x}\_1 + \boldsymbol{\beta}\_2 \mathbf{x}\_2 + \boldsymbol{\beta}\_i \mathbf{x}\_i + \dots \dots + \varepsilon \tag{1}$$

= <sup>0</sup> + 1<sup>1</sup> +2<sup>2</sup> + +. . . . .+ (1)

(2)

Douglas C. Montgomery [40] also developed a model relating the process variables

with output response, which was utilized for process prediction and control. The experiment domain's boundary must be investigated in RSM. The process parameters for this analysis were feed rate, cutting speed, and diameter of the drill, and BBED was utilized to formulate the design of experiments. BBED, an alternate way to central composite design (CCD), is the common statistical tool employed to frame the number of experiments. This approach was used to analyze experiments comprising three levels of design. The BBED can be rotated, and it is said to have no fractional factorial design, which makes it easy to understand the results [41]. In RSM, the general correlation of the control parameters and

where Y is the response (thrust force), β0, β1, β2, . . . β<sup>i</sup> are the regression coefficients, x1, x2, . . . x<sup>i</sup> denote the predictor variable, namely, feed rate, diameter, and cutting speed, and ε shows the error occurring in the thin model. Generally, in RSM, quadratic response functions (Equation (2)) are used and are represented as follows: = <sup>0</sup> +∑1<sup>1</sup> =1 + ∑ 2 =1 + ∑ + <

$$\mathbf{Y} = \boldsymbol{\beta}\_0 + \sum\_{i=1}^k \boldsymbol{\beta}\_1 \mathbf{x}\_1 + \sum\_{i=1}^k \boldsymbol{\beta}\_{li} \mathbf{x}\_i^2 + \sum\_{i$$

#### **3. Results and Discussion** Silane-treated CB woven fabric composites were drilled using twisted drill bits, and

#### *3.1. Thrust Force Analysis* the thrust force signals were obtained, as shown in Figure 2. The plot was obtained during

Silane-treated CB woven fabric composites were drilled using twisted drill bits, and the thrust force signals were obtained, as shown in Figure 2. The plot was obtained during drilling at a speed of 1800 rpm and at a feed of 60 mm/min. This plot depicts different phases, as shown in the figure: initially, the thrust force was found to be low before the entry of the drill bit and during the progress of the drill bit. Next, the thrust force increased to the maximum level during the process of making the hole, i.e., full contact of the drill bit and the specimen. Then, as the drill bit reached the bottom of the specimen, it gradually reduced the thrust force and marked the exit of the drill bit where the thrust force finally attained a zero value [37]. drilling at a speed of 1800 rpm and at a feed of 60 mm/min. This plot depicts different phases, as shown in the figure: initially, the thrust force was found to be low before the entry of the drill bit and during the progress of the drill bit. Next, the thrust force increased to the maximum level during the process of making the hole, i.e., full contact of the drill bit and the specimen. Then, as the drill bit reached the bottom of the specimen, it gradually reduced the thrust force and marked the exit of the drill bit where the thrust force finally attained a zero value [37].

**Figure 2.** Average cutting forces signal that used a twist drill at 1800 rpm and 60 mm/min. **Figure 2.** Average cutting forces signal that used a twist drill at 1800 rpm and 60 mm/min.

*3.2. Response Surface Methodology (RSM)*

By considering input variables, a quadratic equation was newly developed for determining thrust force in order to attain exact process parameters. The final model for thrust force is

$$\begin{array}{l} \text{Y}(N) = 19.9525 &+ (4.072667 \times d) - (0.30465 \times f) + (0.004042 \times n) - (0.03583 \times d \times f) \\ \text{ } & (0.00049 \times d \times n) + (0.000183 \times f \times n) - (0.06561 \times d^2) + (0.007761 \times f^2) - (3.5E \\ & -6 \times n^2) \end{array} \tag{3}$$

Table 4 shows the experimental results for the drilling of CB woven fiber reinforced hybrid epoxy composites using BBED.


**Table 4.** The experimental results for the BBED.

Table 5 shows the ANOVA table for thrust force. Based on F and *p*-value, the RSM model can be simultaneously analyzed, considering all the design factors and the value of F, which predicts the quality of the entire design model. The very low probability factors identified by *p*-value have an insignificant effect on the response [42]. The higher F value and lower *p*-value (*p*-value < 0.05) significance enhanced the fit for the predicted designed model with the experimental model. The experimental result was calculated with a confidence level of 95%. Any value of *p* less than 0.05 indicates that the developed thrust model will be significant. However, the F value of the lack of a fitted model is 0.372, which indicates that the model is insignificant. It proves that the planned model is very suitable for the prediction of thrust force during drilling of developed hybrid composites [43]. In Table 5, it is shown that the feed rate (f) and spindle speed (N) are significant. The coefficient correlation (R-Seq) and adjusted coefficient correlation (0.98438 and 0.9643) indicate that the model is sufficient, and so the procedures were successfully utilized to find the cutting thrust force during the drilling process in composite laminates.

Figure 3a depicts the normal percentage of probability plot with respect to internally studentized residuals and it is observed from the graph that the result is grouped to form a straight line; hence, it confirms the efficiency of the developed model. Similar results were obtained by some of the researchers previously [44]. The association between the actual and predicted thrust is depicted Figure 3b. This result shows that the predicted values are in close agreement with the experimental values [35].

Figure 4 shows the interrelationship between the considered process parameters and the cutting thrust force of CB-reinforced epoxy hybrid composite laminates. Figure 4a explains the influence of the drill diameter on the thrust force. It turns out that as the diameter of the drill bit increases, the thrust force also increases [42].

As the diameter of the drill bit increases, the contact area between the drill bit and the laminate becomes larger; hence, due to higher heat generated, the thrust force increases. The influence of feed speed on thrust is depicted in Figure 4b. The results show that whenever the rate of feed increases, the thrust force also increases. As the diameter of the drill (d) and the feed (f) increases, the cross-sectional area (A) increases because of the formed chip area, A = df/4, and the thrust force also increases, so the axial thrust and the resistance to the developed chip increase [45]. Figure 4c shows that the increase in spindle speed decreases the thrust force, and due to higher spindle speed, the chip formation

posites.

f

softens, thereby minimizing the heat of the material and increasing the thrust force. It could be stated from the study that the change in cutting speed and feed rate influenced the thrust force significantly, irrespective of the working environment [46]. d<sup>2</sup> 1 1.468 1.468 1.209 0.3079 <sup>2</sup> 1 40.581 40.581 33.422 0.0007

**Prob > F**


**Table 5.** Analysis of variance (ANOVA) table for thrust force. Residual 7 8.499 1.214

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**Source <sup>d</sup><sup>f</sup> Adj SS Adj MS F Value** *<sup>p</sup>***-Value**

d 1 23.256 23.256 19.154 0.0032 f 1 335.405 335.405 276.236 <0.0001 n 1 89.646 89.646 73.832 <0.0001

d<sup>n</sup> 1 3.168 3.168 2.609 0.1503 n<sup>f</sup> 1 19.360 19.360 15.945 0.0052

n<sup>2</sup> 1 6.584 6.584 5.423 0.0527

**Table 5.** Analysis of variance (ANOVA) table for thrust force.

Regression model 9 535.903 59.545 49.041 <0.0001

**Figure 3.** (**a**) Normal probability of residuals. (**b**) Predicated thrust force in the drilling of the com-**Figure 3.** (**a**) Normal probability of residuals. (**b**) Predicated thrust force in the drilling of the composites.

Figure 4 shows the interrelationship between the considered process parameters and the cutting thrust force of CB-reinforced epoxy hybrid composite laminates. Figure 4a explains the influence of the drill diameter on the thrust force. It turns out that as the diameter of the drill bit increases, the thrust force also increases [42]. Figure 5 shows the 3D model of the thrust force during the drilling of silane-treated CB composite laminate. The effect of the reaction among processing variables is discussed through 3D response plots. The 3D model has a graph that shows the influence of two changing variables while maintaining the constant third parameter throughout the experiments. Figure 5a indicates the variation of thrust force with the diameter of the drill and the rate of feed, keeping the cutting speed constant. The results show that the thrust

force increases as the diameter of the drill bit and feed rate increase. Figure 5b shows the impact of drill diameter and cutting speed on thrust force. As a result, it was found that as the spindle speed of the CB composite laminate increased, the thrust force decreased. Figure 5c explains the plot of cutting speed and feed rate on the thrust force of the CB composite laminate. The findings show that the enhancement in feed and speed increased the thrust force of the CB composite laminate. It can be seen from the experiment that the drill diameter, feed rate, and cutting speed affected the machining operation irrespective of the working environment. formed chip area, A = df/4, and the thrust force also increases, so the axial thrust and the resistance to the developed chip increase [45]. Figure 4c shows that the increase in spindle speed decreases the thrust force, and due to higher spindle speed, the chip formation softens, thereby minimizing the heat of the material and increasing the thrust force. It could be stated from the study that the change in cutting speed and feed rate influenced the thrust force significantly, irrespective of the working environment [46].

As the diameter of the drill bit increases, the contact area between the drill bit and

the laminate becomes larger; hence, due to higher heat generated, the thrust force increases. The influence of feed speed on thrust is depicted in Figure 4b. The results show that whenever the rate of feed increases, the thrust force also increases. As the diameter of the drill (d) and the feed (f) increases, the cross-sectional area (A) increases because of the

*Polymers* **2023**, *15*, x FOR PEER REVIEW 9 of 15

**Figure 4.** Influence of machining characterization on thrust force in CB fabric/epoxy composites. iments. Figure 5a indicates the variation of thrust force with the diameter of the drill and **Figure 4.** Influence of machining characterization on thrust force in CB fabric/epoxy composites.

Figure 5 shows the 3D model of the thrust force during the drilling of silane-treated

changing variables while maintaining the constant third parameter throughout the exper-

**Figure 5.** A 3D response plot for thrust force during drilling of silane-treated CB fabric/epoxy composites; (**a**) influence of drill diameter and feed rate on thrust force; (**b**) influence of drill diameter and spindle speed on thrust force; and (**c**) influence of feed rate and spindle speed on thrust force. **Figure 5.** A 3D response plot for thrust force during drilling of silane-treated CB fabric/epoxy composites; (**a**) influence of drill diameter and feed rate on thrust force; (**b**) influence of drill diameter and spindle speed on thrust force; and (**c**) influence of feed rate and spindle speed on thrust force.

the rate of feed, keeping the cutting speed constant. The results show that the thrust force increases as the diameter of the drill bit and feed rate increase. Figure 5b shows the impact of drill diameter and cutting speed on thrust force. As a result, it was found that as the spindle speed of the CB composite laminate increased, the thrust force decreased. Figure 5c explains the plot of cutting speed and feed rate on the thrust force of the CB composite laminate. The findings show that the enhancement in feed and speed increased the thrust force of the CB composite laminate. It can be seen from the experiment that the drill diameter, feed rate, and cutting speed affected the machining operation irrespective of the

#### *3.3. Comparison of Experimental and Thrust Force Values 3.3. Comparison of Experimental and Thrust Force Values*

In this experiment, the R-squared value of the model is greater than 90%, and so the modeling system can be effectively used to predict the different properties of composite drilling. The comparison between the RSM value and the experimental value is shown in Figure 6. It is proven that the relevant statistics between the experimental model result and the RSM predicted result can be obtained through the RSM technique. In this experiment, the R-squared value of the model is greater than 90%, and so the modeling system can be effectively used to predict the different properties of composite drilling. The comparison between the RSM value and the experimental value is shown in Figure 6. It is proven that the relevant statistics between the experimental model result and the RSM predicted result can be obtained through the RSM technique.

## *3.4. SEM Observation*

working environment.

The morphology of the drilled holes in silane-treated CB composite laminates was examined using a scanning electron microscope at various locations. The SEM micrograph shows the microstructure of the drilled surface. It could be observed that debonding between fiber and matrix, surface roughness, and fiber pull-out exist in the drilled silanetreated CB fabric epoxy composite laminate using a solid twist drill. Figure 7 shows that minimum cutting speed and maximum feed rate damage the drill walls and cause large fiber pull-out and surface roughness. This is due to the effect of high thrust force applied on the drill zone of the composite laminate [47]. Figure 8 shows the morphology of the composite specimen under better cutting behavior with less drill damage during drilling when compared to high-speed cutting and minimum feed rate conditions. This is due to less contact of the drill bit with the CB composite laminates. The result indicates that the thrust force is low and has a better quality of hole at lower feed rates [33,48].

**Figure 6.** Comparison of experimental and RMS thrust force values. **Figure 6.** Comparison of experimental and RMS thrust force values.

**Figure 7.** SEM image of drilled hole drill (1800 rpm and 60 mm/min). **Figure 7.** SEM image of drilled hole drill (1800 rpm and 60 mm/min).

**Figure 8.** SEM image of drilled hole drill (2400 rpm and 20 mm/min).

force on the composite material.

ysis. The following conclusions were obtained from the experimental study.

was during the complete contact of drill bit with the composites.

Cotton/bamboo woven fabric reinforced epoxy composite laminates were manufactured using a compression molding technique and were subjected to machinability anal-

icant phases of thrust force could be obtained. The maximum value of thrust force

posite can be examined by RSM. The optimum cutting parameters are identified at a minimum rate of feed (20 mm/min), maximum spindle speed (2400 rpm), and drill diameter (6 mm). The adequacy of this model is proven to be fitted for finding thrust

• From the thrust force analysis, it could be observed that during drilling, three signif-

• The influence of various process parameters on the thrust force on fabric/epoxy com-

**4. Conclusions**

**Figure 7.** SEM image of drilled hole drill (1800 rpm and 60 mm/min).

**Figure 8.** SEM image of drilled hole drill (2400 rpm and 20 mm/min). **Figure 8.** SEM image of drilled hole drill (2400 rpm and 20 mm/min).

## **4. Conclusions**

**4. Conclusions** Cotton/bamboo woven fabric reinforced epoxy composite laminates were manufactured using a compression molding technique and were subjected to machinability anal-Cotton/bamboo woven fabric reinforced epoxy composite laminates were manufactured using a compression molding technique and were subjected to machinability analysis. The following conclusions were obtained from the experimental study.


**Author Contributions:** Conceptualization, K.A., S.K.P. and B.M.; methodology, R.L., S.K.D. and N.S.R.; software, K.A., S.K.P. and B.M.; validation, K.A., S.K.P., R.L., B.M., S.K.D., N.S.R. and S.P.S.; formal analysis, K.A., S.K.P. and B.M.; investigation, K.A. and B.M.; resources, K.A. and B.M.; data curation, K.A.; S.K.P. and B.M.; writing—original draft preparation, K.A. and S.K.P.; writing—review and editing, S.K.P. and R.L.; visualization, S.K.D., N.S.R. and S.P.S.; supervision, S.K.P. and S.P.S.; project administration, S.K.P. and S.P.S. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research received no external funding.

**Institutional Review Board Statement:** Not applicable.

**Data Availability Statement:** The data presented in this study are available on request from the corresponding author. The data are not publicly available due to privacy.

**Acknowledgments:** The authors would like to thank Akshaya College of Engineering, K. S. R. College of Engineering and K. S. Rangasamy College of Technology for providing the necessary research facilities and technical support to carry out this research work successfully.

**Conflicts of Interest:** The authors declare no conflict of interest.

### **References**


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## *Article* **Evaluation of Physico-Mechanical Properties on Oil Extracted Ground Coffee Waste Reinforced Polyethylene Composite**

**Hoo Tien Nicholas Kuan 1,\* , Ming Yee Tan <sup>1</sup> , Mohamad Zaki Hassan <sup>2</sup> and Mohamed Yusoff Mohd Zuhri 3,4**


**Abstract:** The current work discusses ground coffee waste (GCW) reinforced high-density polyethylene (HDPE) composite. GCW underwent two types of treatment (oil extraction, and oil extraction followed by mercerization). The composites were prepared using stacking HDPE film and GCW, followed by hot compression molding with different GCW particle loadings (5%, 10%, 15% and 20%). Particle loadings of 5% and 10% of the treated GCW composites exhibited the optimum level for this particular type of composite, whereby their mechanical and thermal properties were improved compared to untreated GCW composite (UGC). SEM fracture analysis showed better adhesion between HDPE and treated GCW. The FTIR conducted proved the removal of unwanted impurities and reduction in water absorption after the treatment. Specific tensile modulus improved for OGC at 5 vol% particle loading. The highest impact energy absorbed was obtained by OGC with a 16% increment. This lightweight and environmentally friendly composite has potential in high-end packaging, internal automotive parts, lightweight furniture, and other composite engineering applications.

**Keywords:** polyethylene; composite materials; ground coffee waste; impact; TGA

## **1. Introduction**

The escalating demand for synthetic fiber has led to over exploration of natural resources causing environmental pollution as they are resistant to degradation. The production and disposal treatment processes of synthetic fiber contribute to environmental pollution as well. In order to protect the earth, natural fiber-reinforced biodegradable polymer matrix and bio-composite have gained attention from researchers worldwide as they are green and eco-friendly materials. Natural fiber is an eco-friendly resource obtained from wood [1], bamboo [2], and agricultural waste [3–6] such as straw [7] that comprises cellulose, hemicellulose, lignin and aromatics, waxes and other lipids, ash, and water-soluble compounds [8]. It is biodegradable, sustainable, economically viable, low density and possesses good specific strengths and moduli [9–14]. Both the natural filler and polymer matrix used in this study are fully biodegradable at the end of their life cycles.

Coffee beans are seeds of the genus flowering plant, Coffea, a member of the botanical family of *Rubiaceace* which comprises almost 500 genera and more than 6500 species [8]. They are normally distributed in tropical and sub-tropical regions. Coffee is the world's second largest traded commodity, and the top five world leading countries in coffee production are Brazil, Vietnam, Indonesia, Colombia, and Ethiopia [15]. In accordance with the International Coffee Organization, world coffee consumption was about 9 million tons or 149.82 million bags (60 kg per bag) in 2014, an increment of 1.8% since 2013 [16].

**Citation:** Kuan, H.T.N.; Tan, M.Y.; Hassan, M.Z.; Zuhri, M.Y.M. Evaluation of Physico-Mechanical Properties on Oil Extracted Ground Coffee Waste Reinforced Polyethylene Composite. *Polymers* **2022**, *14*, 4678. https://doi.org/ 10.3390/polym14214678

Academic Editor: Swarup Roy

Received: 28 September 2022 Accepted: 12 October 2022 Published: 2 November 2022

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2022 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

Meanwhile, the solid waste generated by coffee production is about one ton annually [17]. Concerted efforts are underway to cultivate sustainable practices for the disposal of residues from processing, such as GCW.

Oxo-biodegradable plastic is a commodity plastic containing highly degradable plastic additive (TDPA) and 2–3% transition metals. The metal salts catalyze the degradation process while speeding it up to degrade abiotically at the end of its useful life in the presence of oxygen faster than ordinary plastic [18–20]. The European Standards Organization, CEN, defines oxo-biodegradation as degradation resulting from oxidative and cell-mediated phenomena, either simultaneously or successively [21]. In addition, the useful life of product made using biodegradable plastic could be programmed at manufacture, and the existing technology, machinery and workforce in factories can manufacture it at no extra cost. Biodegradable plastic retains all the advantages of conventional plastic but displays the required additional characteristic of faster degradation after use and disposal by natural mechanisms and with no unwanted environmental consequences [22–24].

Fermoso et al. [25] used thermogravimetric analysis (TGA) and mass spectrometry (MS) to carry out pyrolysis tests on spent coffee grounds in a nitrogen atmosphere with heating rates ranging from 5 to 100 ◦C/min. From the test, they found that the main gases produced during ground coffee pyrolysis are those containing oxygen: H2O, CO and CO<sup>2</sup> within a temperature range between 250 and 425 ◦C. When a zeolite is used as a catalyst, they found a significant increase in the amount of low molecular weight hydrocarbons such as CH4, C2H<sup>4</sup> and C2H<sup>6</sup> [25]. Cai et al. [26] used the thermogravimetric technique together with Fourier transform infrared spectroscopy to analyze the thermal decomposition of some plant residues used in the preparation of tea. They used two mixed N2/O<sup>2</sup> and CO2/O<sup>2</sup> working atmospheres and found that the main gases resulting from the combustion process are: O2, H2O, CH4, CO, CO2, NH<sup>3</sup> and HCN [26]. GCW has a structure that is very similar to components often utilized as filler particles in wood polymer composites. This by-product stands out due to its high protein content (17–18%), which may improve the characteristics of polymer composites. Moreover, because of the presence of these functional groups, it may provide extra opportunities for interfacial bonding within the polar polymeric matrix. Table 1 summarizes the outcomes of previous studies on the mechanical and thermal properties of coffee waste element reinforcement in polymer composites. Even though the hydrophilic nature of GCW may reduce the composite performance due to their poor matrix compatibility and tendency to absorb water, surface treatments may increase the compatibility of the matrix and GCW reinforcement. For example, the treated waste coffee fiber using steam explosion treatment could enhance the thermostability of the composite, increase the fiber crystallinity index and finally be suitable for use as a green composite product [27]. Also, according to Zhao et al. [28], treating the natural fiber using sodium hydroxide (NaOH) increased the interfacial interaction between the matrix and the fiber by up to 300%. Goncalves et al. [29] stated that a NaOH concentration of less than 5 wt% could improve the interfacial bonding between coffee waste and the polyurethane matrix. However, increasing NaOH concentrations up to 10 wt% worsens coffee husk fiber by eroding and deteriorating the structure. Generally, the NaOH solutions ionize the hydroxyl groups in the substrate and chemically alter the fiber structure by dissolving hydrogen bonds. The effectiveness of this process is determined by elements such as response time, heat applied, and solution concentration.

Furthermore, the proportion of GCW in the structure is an essential component that influences the mechanical performance of composites. Mechanical properties increase if high-strength GCW content increases. However, if the quantity of this GCW exceeds its ideal level, the mechanical capabilities may deteriorate. This is owing to the belief that as fiber volume increases, matrix composition decreases, resulting in poorer interfacial interaction. Table 1 reveals that almost all the investigations used a filler ratio of 1–20 wt% and, in some cases, up to 30 wt% of fiber loading. As can be seen from the table, an improvement of the elastic modulus of coffee silverskin-reinforced PP was obtained at the higher loading of fiber, as reported by Dominici et al. [30]. Similar findings were also reported

by Sarasini et al. [31] and Hejna et al. [32]. When coffee silverskin was added to a green polymer, such as poly (butylene succinate) or poly (butylene adipate-co-terephthalate)/poly (3-hydroxybutyrate-co-3-hydroxyvalerate), the tensile strength of the composite structure increased. However, this effect was not significantly improved for stronger poly (lactic acid) matrices [33]. Nevertheless, a great deal of research solely examined the effects of coffee silverskin on morphological, mechanical, and thermal efficiency in composite technology. It was noted that the stiffness and crystallinity of polymer matrices were improved by adding coffee silverskin. However, the effects of GCW on the physico-mechanical properties and thermal stability of composites have not been reported.

**Table 1.** Previous reported work on the mechanical and thermal characteristics of coffee wastereinforced composites.


PP: polypropylene, PBAT: poly(butylene adipate-co-terephthalate), PHBV: poly(3-hydroxybutyrate-co-3 hydroxyvalerate), PLA: poly(lactic acid), PBS: poly(butylene succinate), PE: polyethylene, HDPE: high-density polyethylene, PU: polyurethane.

The aim of this study is to reuse the waste from the coffee beverage industry while reinforcing it with HDPE to produce green and eco-friendly lightweight composites. Oil extracted from GCW using the Soxhlet method can be used to produce biodiesel and for the synthesis of bio-plastic or other potential applications, while the extracted GCW can be used as reinforcement in composite fabrication. The GCW underwent oil extraction and a combination of oil extraction followed by alkaline treatment. The composites were prepared using different particle loadings and were characterized by tensile, impact, FTIR, SEM, TGA, DSC and water absorption.

#### **2. Materials and Methods**

### *2.1. Materials*

Ground coffee waste (GCW) (Arabica) was obtained from a local cafe. The GCW was washed with distilled water, dried in an oven at 105 ◦C for 24 h to 1–2% moisture content, and then sieved to a particle size of 850 µm. Table 2 shows the chemical composition of GCW from previous work conducted by other researchers. Commercial bio-degradable high-density polyethylene (HDPE) thin film (density of 0.96 g/mL at 23 ◦C, melt flow index of 0.35 g/10 min, 190 ◦C/5 kg) was used as the polymer matrix. GCW was extracted using the Soxhlet method in a solvent, n-hexane. Approximately 200 g of dried GCW was extracted in a Soxhlet extractor with 800 ml of n-hexane at a reaction temperature between 80 ◦C and 95 ◦C for 3 h. The oil-extracted GCW (OG) was dried at 105 ◦C for 24 h to evaporate n-hexane. After that, the OG was soaked in 1% concentration of NaOH aqueous solution (0.25 molarity) at room temperature for 24 h. Then, the NaOH-treated OG (ONG) was rinsed with distilled water until it reached pH 7. Both the OG and ONG were dried in an oven at 105 ◦C for 24 h before use.


**Table 2.** Chemical composition of GCW.

#### *2.2. Methods*

2.2.1. Compounding and Compression Molding

The dimensions of HDPE thin film used were 230 mm × 230 mm. Each laminate was prepared by sieving GCW in between plies of HDPE films. The compounding of GCW and HDPE were placed in a mold and covered. The composite was prepared using a hot compression molding method at 3 MPa for approximately 15 min until it reached 150 ◦C. The hot compression mold was then left to cool to room temperature for approximately 2 h before the laminate was taken out. Table 3 shows a summary of all the composites investigated in this research. The results of HDPE and UGC from the previous study were compared, whereby the earlier study performed on modified ground coffee waste showed an enhancement of the properties as compared to untreated ground coffee waste composite [39,40]. OGC is oil-extracted GCW-reinforced HDPE composite (Figure 1) and ONGC is oil-extracted + NaOH GCW-reinforced HDPE composite with different particle loadings.


#### **Table 3.** Summary of type of composites investigated in this study.

UG = Untreated ground coffee waste, OG = Oil-extracted ground coffee waste, ONG = Oil-extracted + NaOH ground coffee waste. 1) and ONGC is oil-extracted + NaOH GCW-reinforced HDPE composite with different particle loadings.

**Figure 1.** OGC, oil-extracted GCW-reinforced HDPE composite laminate. **Figure 1.** OGC, oil-extracted GCW-reinforced HDPE composite laminate.

**Table 3.** Summary of type of composites investigated in this study. 2.2.2. Composite Property Testing

**Composite Sample Stacking**  A. Fourier Transform Infrared Spectroscopy (FT-IR)

**Sequence GCW (%) HDPE (%) References**  HDPE HDPE 150PE 0 100 [39] UGC5 Untreated GCW5/HDPE 150PE + 5UG 5 95 [39] FTIR was analyzed using an IRAffinity-1 from Shimadzu, Kyoto, Japan. Approximately 20 mg of GCW each were put on the ATR (Attenuated Total Reflectance). Each sample was subjected to 20 scans in the range of 600 cm−1–4000 cm−<sup>1</sup> .

**Volume of** 

15OG 15 85

20OG 20 80

5ONG 5 95

10ONG 10 90

15ONG 15 85

20ONG 20 80

**Volume of** 

UGC10 Untreated GCW10/HDPE 150PE + B. Scanning Electron Microscopy (SEM)

OGC15 Oil Extraction

ONGC5 Oil Extraction +

ONGC10 Oil Extraction +

ONGC15 Oil Extraction +

ONGC20 Oil Extraction +

10UG 10 90 [39] UGC15 Untreated GCW15/HDPE 150PE + 15UG 15 85 [39] SEM (Hitachi TM3030) with 15 kV electrons was used to investigate all the samples. Before observation, the samples were coated with a thin layer of gold using a JEOL JFC-600 auto fine coater.

	- 20UG 20 80 [39] OGC5 Oil Extraction GCW5/HDPE 150PE + 5OG 5 95 OGC10 Oil Extraction 150PE + 10OG 10 90 I. Thermogravimetric Analysis (TGA) TGA was performed via Netzsch TG 200 F3 Tarsus (Selb, Germany). Approximately 20 mg of the composite was prepared and put into a crucible. A temperature range from 30 ◦C to 600 ◦C was used with a heating rate of 10 ◦C min−<sup>1</sup> in a nitrogen atmosphere. The

150PE +

150PE +

150PE +

150PE +

150PE +

150PE +

GCW10/HDPE

GCW15/HDPE

GCW20/HDPE

NaOH5/HDPE

NaOH10/HDPE

NaOH15/HDPE

NaOH20/HDPE

mass loss, initial degradation temperature (Tonset), maximum degradation temperature (Tmax), and final degradation temperature (Tfinal) were determined.

	- i. Tensile Testing Tensile properties were examined using a Shimadzu universal testing machine, model AG-300K IS MS (Kyoto, Japan). Tensile tests were carried out in accordance with ASTM using a load cell of 50 kN at a crosshead rate displacement of 1 mm/min [41]. Tensile strength and tensile modulus were based on initial sample dimensions and the results were averaged over five measurements.
	- ii. Impact Testing Dart drop impact properties were characterized using Instron 9250 HV (U.S.) in accordance with ASTM D3763-15. A 12.88 mm hemisphere head of 3.4727 kg was employed and a free-falling initial drop height was set at 0.50 m. A total of five different samples were subjected to impact testing. The composites can be tested to assess their resistance to falling weight [42].

A water absorption test was performed according to ASTM D570-98. Test specimens were prepared at a size of 76.2 mm × 25.4 mm by the thickness of the composite. The specimens were immersed entirely in a container of distilled water maintained at a temperature of 23 ± 1 ◦C for the test. The specimens were weighed after 2 h and then repeatedly every 24 h after that. The specimens were removed from distilled water one at a time, wiped with a dry cloth and weighed to the nearest 0.001 g. The percentage of water absorption was measured by [43]:

$$percentage\text{ of water absorption, }\%=\frac{W\_{final}-W\_{initial}}{W\_{initial}}\times\ 100\% \tag{1}$$

### **3. Results and Discussion**

### *3.1. FT-IR Spectroscopy*

The FT-IR spectroscopy in Figure 2 shows exemplary peaks of untreated GCW (UG), oil-extracted GCW (OG) and oil-extracted + NaOH GCW (ONG), respectively, at 3351 cm−<sup>1</sup> , 2923 cm−<sup>1</sup> , 2859 cm−<sup>1</sup> , 1733 cm−<sup>1</sup> , 1444 cm−<sup>1</sup> , 1367 cm−<sup>1</sup> , 1238 cm−<sup>1</sup> , 1120 cm−<sup>1</sup> and 1031 cm−<sup>1</sup> . The absorption band of 3351 cm−<sup>1</sup> belongs to the absorption of –OH stretching vibrations on the GCW. The –OH compound consisted of the presence of water, aliphatic primary and secondary alcohol found in cellulose, hemicellulose, lignin, extractives and carboxylic acids in extractives [44,45]. The decreasing area of this peak represented the removal of lipids. Peaks at 2859 cm−<sup>1</sup> and 2923 cm−<sup>1</sup> are recognized as asymmetric and symmetric stretching of C–H bonds in aliphatic chains that are ascribed to lipids [46]. After both treatments, the amount of lipids was reduced. The characteristic GCW peaks at 1733 cm−<sup>1</sup> and 1238 cm−<sup>1</sup> were credited to the acetylated xylan (hemicellulose), aldehyde lignin or carbonyl lipid, pectin and wax [47,48]. A peak of 1444 cm−<sup>1</sup> assigned to benzene ring stretching in lignin and =CH<sup>2</sup> vibration in polyoses (Hemicellulose) [49], while 1120 cm−<sup>1</sup> attributed to the aromatic C–H in the plane deformation mode

of the guaiacyl/syringyl units of lignin [44] disappeared after both treatments. A peak of 1367 cm−<sup>1</sup> shown as the in-plane C–H bending in hemicellulose reduced after the treatments [50]. During the 600 cm−1–1100 cm−<sup>1</sup> peak, the –OH concentration increased from UG to OG followed by ONG, revealing this was a more active site available to intermingle between filler/matrix interfaces. Table 4 demonstrates the details of band wavelength and the associated chemical group. guaiacyl/syringyl units of lignin [44] disappeared after both treatments. A peak of 1367 cm−1 shown as the in-plane C–H bending in hemicellulose reduced after the treatments [50]. During the 600 cm−1–1100 cm−1 peak, the –OH concentration increased from UG to OG followed by ONG, revealing this was a more active site available to intermingle between filler/matrix interfaces. Table 4 demonstrates the details of band wavelength and the associated chemical group.

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, % ൌ ି

The FT-IR spectroscopy in Figure 2 shows exemplary peaks of untreated GCW (UG), oil-extracted GCW (OG) and oil-extracted + NaOH GCW (ONG), respectively, at 3351 cm−1, 2923 cm−1, 2859 cm−1, 1733 cm−1, 1444 cm−1, 1367 cm−1, 1238 cm−1, 1120 cm−1 and 1031 cm−1. The absorption band of 3351 cm−1 belongs to the absorption of –OH stretching vibrations on the GCW. The –OH compound consisted of the presence of water, aliphatic primary and secondary alcohol found in cellulose, hemicellulose, lignin, extractives and carboxylic acids in extractives [44,45]. The decreasing area of this peak represented the removal of lipids. Peaks at 2859 cm−1 and 2923 cm−1 are recognized as asymmetric and symmetric stretching of C–H bonds in aliphatic chains that are ascribed to lipids [46]. After both treatments, the amount of lipids was reduced. The characteristic GCW peaks at 1733 cm−1 and 1238 cm−1 were credited to the acetylated xylan (hemicellulose), aldehyde lignin or carbonyl lipid, pectin and wax [47,48]. A peak of 1444 cm−1 assigned to benzene ring stretching in lignin and =CH2 vibration in polyoses (Hemicellulose) [49], while 1120 cm−1 attributed to the aromatic C–H in the plane deformation mode of the

.

% **(1)** 

**3. Results and Discussion**  *3.1. FT-IR Spectroscopy* 

**Figure 2.** FTIR Spectra of a) UG (green), b) OG (red) and c) ONG (black). **Figure 2.** FTIR Spectra of a) UG (green), b) OG (red) and c) ONG (black).


**Table 4.** FTIR results of UG, OG and ONG. **Table 4.** FTIR results of UG, OG and ONG.

### *3.2. TGA Analysis*

Figure 3 shows the emblematic TGA and DTG curve of HDPE, untreated GCW/HDPE (UGC), oil-extracted GCW/HDPE (OGC) and oil-extracted + NaOH/HDPE (ONGC) composites based on particle loadings of 10%. As shown in Tables 5 and 6, the decomposition began at a lower temperature in the composites than in the neat HDPE due to the presence of GCW. However, the degradation temperature of OGC and ONGC were higher than HDPE at 50% weight loss. The process of oil extractions and alkaline treatment had improved their thermal stabilities. The thermal stability increased in UGC, OGC and ONGC, respectively. The percentage of charred residue decreased from UGC to OGC and followed by ONGC. This showed the treatment effectiveness as more unwanted impurities were removed leading to less charred residue. Weight losses before 100 ◦C were due to the moisture absorption in the sample. The first step of degradation can be observed in the range of 250 ◦C–400 ◦C, marking the degradation of hemicellulose, cellulose and small amount of lignin [51]. The second step of degradation, which is also the main decomposition of GCW, can be observed at approximately 440 ◦C that contributed to the dramatic degradation of lignin, as shown in DTG curve [52]. The degradation temperature of the materials was increased in the sequence of hemicellulose < cellulose < lignin. is also the main decomposition of GCW, can be observed at approximately 440 °C that contributed to the dramatic degradation of lignin, as shown in DTG curve [52]. The degradation temperature of the materials was increased in the sequence of hemicellulose < cellulose < lignin.

1238 Hemicellulose, Lignin peak Present Reduced Reduced 1120 Lignin peak Present - - 600–1100 –OH intensity Low High High

Figure 3 shows the emblematic TGA and DTG curve of HDPE, untreated GCW/HDPE (UGC), oil-extracted GCW/HDPE (OGC) and oil-extracted + NaOH/HDPE (ONGC) composites based on particle loadings of 10%. As shown in Tables 5 and 6, the decomposition began at a lower temperature in the composites than in the neat HDPE due to the presence of GCW. However, the degradation temperature of OGC and ONGC were higher than HDPE at 50% weight loss. The process of oil extractions and alkaline treatment had improved their thermal stabilities. The thermal stability increased in UGC, OGC and ONGC, respectively. The percentage of charred residue decreased from UGC to OGC and followed by ONGC. This showed the treatment effectiveness as more unwanted impurities were removed leading to less charred residue. Weight losses before 100 °C were due to the moisture absorption in the sample. The first step of degradation can be observed in the range of 250 °C–400 °C, marking the degradation of hemicellulose, cellulose and small amount of lignin [51]. The second step of degradation, which

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(**a**) (**b**)

*3.2. TGA Analysis* 

**Figure 3.** (**a**) Thermogravimetric analysis (TGA) thermogram; and (**b**) Differential thermogravi-

metric graphics (DTG) thermogram of HDPE (smooth line), UGC, OGC and ONGC (10% particle loading). **Figure 3.** (**a**) Thermogravimetric analysis (TGA) thermogram; and (**b**) Differential thermogravimetric graphics (DTG) thermogram of HDPE (smooth line), UGC, OGC and ONGC (10% particle loading).


**Table 5** Summary of TGA data for HDPE, UGC, OGC and ONGC (10% particle loading). **Table 5.** Summary of TGA data for HDPE, UGC, OGC and ONGC (10% particle loading).

**Table 6.** Degradation temperature of HDPE, UGC, OGC and ONGC (10% particle loading).

ONGC10 459.1 482.0 492.8 17.71


## *3.3. DSC Analysis*

The sample was heated at a heating rate of 20 ◦C·min−<sup>1</sup> initially, from 50 ◦C to 180 ◦C. When the sample reached 180 ◦C, it was then left to cool off to 50 ◦C in approximately 2 min. In the second cycle, the sample was reheated again to 180 ◦C. The curves exhibited endothermic peaks, as shown in Figure 4. These peaks are attributed to the melting of crystalline domains of the HDPE matrix. Table 7 shows the important data of DSC measurement. Thermal characteristics such as crystalline melting temperature, Tm, melting enthalpy of composite according to the content of HDPE in the composite, ∆Hm, and degree of crystallinity, Xc, were identified. The content of HDPE in the composite, ∆Hm, was calculated by using the area under the curve; while X<sup>c</sup> was determined based on the ratio of ∆H<sup>m</sup> to ∆H<sup>0</sup> <sup>m</sup>, melting enthalpy of 100% crystalline HDPE (293 Jg−<sup>1</sup> ) [53]. With regard to melt temperature, the peaks were located in the same temperature range of 137 ◦C–138 ◦C. In all composite formulations, X<sup>c</sup> increased with the incorporation of GCW

content. In this case, GCW acted as nucleating agent that produced crystals to increase the percentage of X<sup>c</sup> in the composites. The nucleating effect was due to the low filler content. Furthermore, the removal of lignin from the GCW after treatment also contributed to an increase in crystallinity [54]. content. In this case, GCW acted as nucleating agent that produced crystals to increase the percentage of Xc in the composites. The nucleating effect was due to the low filler content. Furthermore, the removal of lignin from the GCW after treatment also contributed to an increase in crystallinity [54].

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**Temperature at 10% Weight Loss (°C)** 

**Composites** 

*3.3. DSC Analysis* 

**Table 6.** Degradation temperature of HDPE, UGC, OGC and ONGC (10% particle loading).

**Temperature at 15% Weight Loss (°C)** 

HDPE 456.52 462.48 469.74 479.60 UGC10 422.11 450.39 464.64 478.01 OGC10 444.50 456.08 466.70 479.88 ONGC10 448.79 458.33 467.87 480.29

The sample was heated at a heating rate of 20 °C·min−1 initially, from 50 °C to 180 °C. When the sample reached 180 °C, it was then left to cool off to 50 °C in approximately 2 min. In the second cycle, the sample was reheated again to 180 °C. The curves exhibited endothermic peaks, as shown in Figure 4. These peaks are attributed to the melting of crystalline domains of the HDPE matrix. Table 7 shows the important data of DSC measurement. Thermal characteristics such as crystalline melting temperature, Tm, melting enthalpy of composite according to the content of HDPE in the composite, ∆Hm, and degree of crystallinity, Xc, were identified. The content of HDPE in the composite, ∆Hm, was calculated by using the area under the curve; while Xc was determined based on the ratio of ∆Hm to ∆H0m, melting enthalpy of 100% crystalline HDPE (293 Jg−1) [53]. With regard to melt temperature, the peaks were located in the same temperature range of 137 °C–138 °C. In all composite formulations, Xc increased with the incorporation of GCW

**Temperature at 25% Weight Loss (°C)** 

**Temperature at 50% Weight Loss (°C)** 

**Figure 4.** DSC curves of HDPE (smooth line), UGC10, OGC10 and ONGC10.



#### ONG10 137.35 140.76 53.38 *3.4. Tensile*

The tensile properties were not as good after the inclusion of GCW. Reinforcement with a uniform circular cross-section area and with an aspect ratio more than its critical value would normally improve the tensile strength. However, the irregular shape of GCW fillers with an aspect ratio of less than two cannot support the stress transferred from the polymer matrix [55]. Therefore, the tensile strength reduced significantly. Compared to UGC, both treated GCW/HDPE composites showed an improvement in tensile strength. The tensile strength of the composites was maximized when reinforced with OGC > ONGC > UGC, respectively (Figure 5a). The tensile strength of OGC increased and reached an optimum of 16.47 MPa at a particle loading of 10%, an increment of 27% as compared to UGC, then declined. A large amount of lipids in OG was removed thus exposing a coarser surface with cellulose micro fibrils revealed, as verified by FTIR (Figure 2) and SEM (Figure 10). The exposed cellulose micro fibrils improved the effectiveness of the contact surface area thus enabling better impregnation of GCW by the HDPE matrix. On the other hand, increasing the particle loading led to the decreasing tensile strength of ONGC (Figure 5a). During a particle loading of 5%, the tensile strength was slightly better than UGC at 15.49 MPa, with 8% improvement. This may be due to less exposure of cellulose micro fibrils on its surface after the treatment. In addition, ONG became porous and weaker after alkaline treatment and this led to its having poor properties.

**Figure 5.** (**a**) Tensile Strength (**b**) Tensile Modulus (**c**) Specific Tensile Strength (**d**) Specific Tensile Modulus of HDPE, UGC, OGC and ONGC of different particle loadings. **Figure 5.** (**a**) Tensile Strength (**b**) Tensile Modulus (**c**) Specific Tensile Strength (**d**) Specific Tensile Modulus of HDPE, UGC, OGC and ONGC of different particle loadings.

*3.5. Impact Test*  A dart drop weight impact test was performed to study the impact response of UGC, OGC and ONGC with a 10% particle loading. Figure 6 and Figure 7 show the results. The ascending curve in Figure 6 signifies the stiffness of the composites. The stiffness of the composites increased from UGC, to OGC to ONGC. The maximum impact force corresponds to the onset of material damage. The maximum load was sustained by the composite before fracture increasing in the sequence of UGC < ONGC (+14%) < OGC (+28%), respectively. Meanwhile, the energy to maximum load increased from UGC < ONGC (+11%) < OGC (+51%) as shown by the dotted line in Figure 6. The decreasing curve shows the properties of the materials. The ductility of the composites experienced a similar trend due to the energy of the maximum load. Total energy absorbed increased in the sequence of UGC < ONGC (+7%) < OGC (+16%), respectively. Hydrophilic UG As observed in Figure 5b, OGC showed a slight improvement as compared to neat HDPE during 5% and 10% of particle volume fraction reaching 1263 MPa (+14%) and 1122 MPa (+1.4%), respectively. The chemical modification proved to provide a better homogeneous dispersion of GCW. The tensile properties showed a proportional increment with particle loadings until 10%. This was due to an enhancement of load transfer between the matrix and GCW interface. Nevertheless, the composites declined in tensile properties beyond optimum values (>10%) due to the agglomeration of GCW. Agglomerations create flaws and voids between GCW and the matrix thus diminishing the tensile properties [56]. Other than that, aspect ratio played a very important role. GCW had a very low aspect ratio of less than two due to insufficient stress transferred; therefore, it displayed a low tensile modulus. In contrast, the tensile modulus obtained for ONGC was lower than HDPE as ONG became weak due to the treatment.

and hydrophobic HDPE had poor interaction thus yielding a lower impact resistance. Both applied treatments had improved impact resistance. After oil extraction, the filler and matrix had a better adhesion and this gave a better uniform distribution of stress transfer. Therefore, more load and energy were required to perforate the samples. The treatments between matrix and reinforcement had significantly improved the toughness of the composite while showing the ability to absorb more energy and deform plastically before fractured. ONGC yielded lower results than OGC. The reason was supported by the results discussed earlier where ONG itself had become weaker. Overall, the impact Figure 5c displays the specific tensile strength, known as strength to weight ratio. All of the treated GCW showed lower specific tensile strength. Particle loadings of 5% and 10% for OGC showed a slight decrease to 0.0173 MPa.m3/kg (−1.1%) and 0.0167 MPa·m3/kg (−4.6%), respectively. Figure 5d illustrates the specific tensile modulus, known as stiffness to weight ratio. The specific tensile modulus of neat HDPE was 1.0335 MPa·m3/kg. The specific tensile modulus improved for every particle loading except for particle loadings of more than 15% for OGC. The specific tensile modulus had improved by 32%. The inclusion of GCW contributed to a lighter weight composite. The specific tensile properties improved

slightly as the composites exhibited a lower density. Conversely, the specific tensile properties for ONGC did not improve due to the poor condition of ONG. Another researcher reported the same trend of results using oil extraction oil palm empty fruit brunch composite [57]. Tensile, flexural and toughness properties were improved significantly more than those without extraction.

#### *3.5. Impact Test*

A dart drop weight impact test was performed to study the impact response of UGC, OGC and ONGC with a 10% particle loading. Figures 6 and 7 show the results. The ascending curve in Figure 6 signifies the stiffness of the composites. The stiffness of the composites increased from UGC, to OGC to ONGC. The maximum impact force corresponds to the onset of material damage. The maximum load was sustained by the composite before fracture increasing in the sequence of UGC < ONGC (+14%) < OGC (+28%), respectively. Meanwhile, the energy to maximum load increased from UGC < ONGC (+11%) < OGC (+51%) as shown by the dotted line in Figure 6. The decreasing curve shows the properties of the materials. The ductility of the composites experienced a similar trend due to the energy of the maximum load. Total energy absorbed increased in the sequence of UGC < ONGC (+7%) < OGC (+16%), respectively. Hydrophilic UG and hydrophobic HDPE had poor interaction thus yielding a lower impact resistance. Both applied treatments had improved impact resistance. After oil extraction, the filler and matrix had a better adhesion and this gave a better uniform distribution of stress transfer. Therefore, more load and energy were required to perforate the samples. The treatments between matrix and reinforcement had significantly improved the toughness of the composite while showing the ability to absorb more energy and deform plastically before fractured. ONGC yielded lower results than OGC. The reason was supported by the results discussed earlier where ONG itself had become weaker. Overall, the impact resistance of the composites enhanced in the sequence of UGC < ONGC < OGC. Figure 8 displays the post-impact fracture behavior of the perforated composites. The failure implicated plastic deformation of the HDPE and debonding of the GCW filler within the composites. With similar particle loading percentages, it seems that the treated GCW composites had better plastic deformation than UGC. The bottom surface of the impacted samples had a bulging effect. Delamination only occurred at the perforated area. A similar impact result trend was obtained by another researcher [58]. *Polymers* **2022**, *14*, x FOR PEER REVIEW 12 of 17 resistance of the composites enhanced in the sequence of UGC < ONGC < OGC. Figure 8 displays the post-impact fracture behavior of the perforated composites. The failure implicated plastic deformation of the HDPE and debonding of the GCW filler within the composites. With similar particle loading percentages, it seems that the treated GCW composites had better plastic deformation than UGC. The bottom surface of the impacted samples had a bulging effect. Delamination only occurred at the perforated area. A similar impact result trend was obtained by another researcher [58].

**Figure 6.** Load-Energy-Time graph of impact test for UGC, OGC and ONGC. **Figure 6.** Load-Energy-Time graph of impact test for UGC, OGC and ONGC.

**Figure 7.** Effect of filler on (**a**) load, (**b**) energy to maximum, (**c**) total energy of GCW/HDPE com-

posite at 10% particle loading.

*Polymers* **2022**, *14*, x FOR PEER REVIEW 13 of 17

**Figure 6.** Load-Energy-Time graph of impact test for UGC, OGC and ONGC.

0 5 10 15 20

Time, ms UGC10 OGC10 ONGC10

**Figure 7.** Effect of filler on (**a**) load, (**b**) energy to maximum, (**c**) total energy of GCW/HDPE composite at 10% particle loading. **Figure 7.** Effect of filler on (**a**) load, (**b**) energy to maximum, (**c**) total energy of GCW/HDPE composite at 10% particle loading. *Polymers* **2022**, *14*, x FOR PEER REVIEW 13 of 17

resistance of the composites enhanced in the sequence of UGC < ONGC < OGC. Figure 8 displays the post-impact fracture behavior of the perforated composites. The failure implicated plastic deformation of the HDPE and debonding of the GCW filler within the composites. With similar particle loading percentages, it seems that the treated GCW composites had better plastic deformation than UGC. The bottom surface of the impacted samples had a bulging effect. Delamination only occurred at the perforated area. A

similar impact result trend was obtained by another researcher [58].

(**a**) (**b**) (**c**) **Figure 8.** Photograph of fracture behavior of perforated samples (**a**) UGC, (**b**) OGC, (**c**) ONGC

0

5

10

Energy, J

15

**Figure 8.** Photograph of fracture behavior of perforated samples (**a**) UGC, (**b**) OGC, (**c**) ONGC (10% particle loading). **Figure 8.** Photograph of fracture behavior of perforated samples (**a**) UGC, (**b**) OGC, (**c**) ONGC (10% particle loading). (10% particle loading).

#### *3.6. Water Absorption 3.6. Water Absorption 3.6. Water Absorption*  Figure 9 depicts the water absorption proficiency of HDPE, UGC, OGC and ONGC.

0

400

800

1200

Load, N

1600

2000

Figure 9 depicts the water absorption proficiency of HDPE, UGC, OGC and ONGC. The water absorption results showed that the non-polar HDPE demonstrated the lowest water absorption. In comparison to similar particle loadings of 10%, the percentage of water absorption increased in HDPE, OGC, ONGC and UGC at 0.05%, 3.22%, 3.80%, 5.22%, respectively, with an improvement of 38% and 27% in UGC. The treated GCW reinforced composites exhibited a lower water absorption rate as it became more hydrophobic after the treatments. Furthermore, a better interfacial adhesion of GCW/HDPE reduced the width of the interface area and diminished water penetration to the material inner parts [59]. Figure 9 depicts the water absorption proficiency of HDPE, UGC, OGC and ONGC. The water absorption results showed that the non-polar HDPE demonstrated the lowest water absorption. In comparison to similar particle loadings of 10%, the percentage of water absorption increased in HDPE, OGC, ONGC and UGC at 0.05%, 3.22%, 3.80%, 5.22%, respectively, with an improvement of 38% and 27% in UGC. The treated GCW reinforced composites exhibited a lower water absorption rate as it became more hydrophobic after the treatments. Furthermore, a better interfacial adhesion of GCW/HDPE reduced the width of the interface area and diminished water penetration to the material inner parts [59]. The water absorption results showed that the non-polar HDPE demonstrated the lowest water absorption. In comparison to similar particle loadings of 10%, the percentage of water absorption increased in HDPE, OGC, ONGC and UGC at 0.05%, 3.22%, 3.80%, 5.22%, respectively, with an improvement of 38% and 27% in UGC. The treated GCW reinforced composites exhibited a lower water absorption rate as it became more hydrophobic after the treatments. Furthermore, a better interfacial adhesion of GCW/HDPE reduced the width of the interface area and diminished water penetration to the material inner parts [59].

*3.7. Morphological Studies* 

*3.7. Morphological Studies* 

properties.

properties.

**Figure 9.** The water absorption proficiency percent mass gain for all the composites with 10% particle loading in comparison with HDPE (smooth line). **Figure 9.** The water absorption proficiency percent mass gain for all the composites with 10% particle loading in comparison with HDPE (smooth line). **Figure 9.** The water absorption proficiency percent mass gain for all the composites with 10% particle loading in comparison with HDPE (smooth line).

Figure 10 shows the surface morphology of UG, OG and ONG. All of the GCW par-

modifications conducted on GCW had increased its surface roughness. This had been proven with the FTIR and TGA with the removal of hemicellulose, lignin, pectin, wax, lipid and impurities on the GCW surface. As seen in Figure 10b, a coarser surface was revealed, and cellulose micro fibrils were exposed on the OG surface after extracting the oil. In Figure 10c, ONG showed uneven concave and rougher surfaces with less cellulose micro fibrils exposure. Additionally, ONG was more fragile, as proven by its mechanical

modifications conducted on GCW had increased its surface roughness. This had been proven with the FTIR and TGA with the removal of hemicellulose, lignin, pectin, wax, lipid and impurities on the GCW surface. As seen in Figure 10b, a coarser surface was revealed, and cellulose micro fibrils were exposed on the OG surface after extracting the oil. In Figure 10c, ONG showed uneven concave and rougher surfaces with less cellulose micro fibrils exposure. Additionally, ONG was more fragile, as proven by its mechanical

#### *3.7. Morphological Studies Polymers* **2022**, *14*, x FOR PEER REVIEW 14 of 17

Figure 10 shows the surface morphology of UG, OG and ONG. All of the GCW particles had a dendritic appearance. The SEM pictograph in Figure 10 illustrates that the modifications conducted on GCW had increased its surface roughness. This had been proven with the FTIR and TGA with the removal of hemicellulose, lignin, pectin, wax, lipid and impurities on the GCW surface. As seen in Figure 10b, a coarser surface was revealed, and cellulose micro fibrils were exposed on the OG surface after extracting the oil. In Figure 10c, ONG showed uneven concave and rougher surfaces with less cellulose micro fibrils exposure. Additionally, ONG was more fragile, as proven by its mechanical properties. Figure 11 shows the SEM micrographs conforming to fracture surfaces from the tensile test at 10% particle loading for UGC, OGC and ONGC. For UGC, the fractured surface showed a gap in between the UG and HDPE, which resulted in poor fiber/matrix adhesion. For oil extraction application, a clearer and better interface can be seen between OG and HDPE with a layer of lipid being removed that led to a coarser surface of OG. The adhesion of ONGC was improved to a better impregnation of matrix in ONG. However, ONG became fragile and porous and thus yielded lower mechanical properties. *Polymers* **2022**, *14*, x FOR PEER REVIEW 14 of 17

**Figure 10.** SEM micrograph of GCW particles of (**a**) UG, (**b**) OG and (**c**) ONG. **Figure 10.** SEM micrograph of GCW particles of (**a**) UG, (**b**) OG and (**c**) ONG.

Figure 11 shows the SEM micrographs conforming to fracture surfaces from the tensile test at 10% particle loading for UGC, OGC and ONGC. For UGC, the fractured surface showed a gap in between the UG and HDPE, which resulted in poor fiber/matrix adhesion. For oil extraction application, a clearer and better interface can be seen between OG and HDPE with a layer of lipid being removed that led to a coarser surface of OG. The adhesion of ONGC was improved to a better impregnation of matrix in ONG. However, ONG became fragile and porous and thus yielded lower mechanical properties. (**a**) (**b**) (**c**) **Figure 10.** SEM micrograph of GCW particles of (**a**) UG, (**b**) OG and (**c**) ONG.

impurities, pectin, wax, hemicellulose, and lignin. Both treatments also produced a more hydrophobic GCW as water absorption rate had decreased. Overall, particle loadings of **Figure 11.** SEM micrograph of breaking surface of (**a**) UGC, (**b**) OGC, (**c**) ONGC (10% particle loading). **Figure 11.** SEM micrograph of breaking surface of (**a**) UGC, (**b**) OGC, (**c**) ONGC (10% particle loading).

#### 5% and 10% exhibited the optimum level for this particular type of composite since mechanical properties had improved as compared to UGC. The incorporation of GCW into **4. Conclusions**

and automotive components.

biodegradable HDPE aimed to improve the recyclability of waste in order to turn it into a value-added product, minimize ecological damage and reduce the production cost. The composites fabricated are able to degrade themselves at the end of their life cycle without leaving any negative effect on the environment. This particular type of composite has potential for application in consumables, packaging, gardening, furniture parts and automotive components. **4. Conclusion**  The tensile and impact properties, thermal analysis and analytical analysis of untreated and treated GCW/HDPE were evaluated. The characterization results show OGC to be the most impressive among all the fabricated composites. FTIR, SEM and TGA showed the effectiveness of the chemical modification performed on GCW by removing impurities, pectin, wax, hemicellulose, and lignin. Both treatments also produced a more hydrophobic GCW as water absorption rate had decreased. Overall, particle loadings of The tensile and impact properties, thermal analysis and analytical analysis of untreated and treated GCW/HDPE were evaluated. The characterization results show OGC to be the most impressive among all the fabricated composites. FTIR, SEM and TGA showed the effectiveness of the chemical modification performed on GCW by removing impurities, pectin, wax, hemicellulose, and lignin. Both treatments also produced a more hydrophobic GCW as water absorption rate had decreased. Overall, particle loadings of 5% and 10% ex-

5% and 10% exhibited the optimum level for this particular type of composite since mechanical properties had improved as compared to UGC. The incorporation of GCW into

a value-added product, minimize ecological damage and reduce the production cost. The composites fabricated are able to degrade themselves at the end of their life cycle without leaving any negative effect on the environment. This particular type of composite has potential for application in consumables, packaging, gardening, furniture parts hibited the optimum level for this particular type of composite since mechanical properties had improved as compared to UGC. The incorporation of GCW into biodegradable HDPE aimed to improve the recyclability of waste in order to turn it into a value-added product, minimize ecological damage and reduce the production cost. The composites fabricated are able to degrade themselves at the end of their life cycle without leaving any negative effect on the environment. This particular type of composite has potential for application in consumables, packaging, gardening, furniture parts and automotive components.

**Author Contributions:** Methodology, H.T.N.K. and M.Y.T.; validation, H.T.N.K., M.Z.H. and M.Y.M.Z.; formal analysis, H.T.N.K. and M.Y.T.; data curation, M.Y.T.; writing—review and editing, H.T.N.K., M.Y.T., M.Z.H. and M.Y.M.Z.; supervision, H.T.N.K. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by Universiti Malaysia Sarawak (UNIMAS) under the Cross Disciplinary Research Grant (F02/CDRG/1823/2019).

**Data Availability Statement:** Not applicable.

**Acknowledgments:** The authors gratefully acknowledge Universiti Malaysia Sarawak (UNIMAS) for providing the opportunity and support to conduct this research through Cross Disciplinary Research Grant (F02/CDRG/1823/2019).

**Conflicts of Interest:** The authors declare no conflict of interest.

## **References**


## *Article* **Investigation of the Contact Interface between Natural Fibre Metal Laminates under Tension Using Finite Element Analysis (FEA)**

**Chun Han Song, Khaled Giasin \* , Abu Saifullah and Antigoni Barouni \***

Advanced Polymers and Composites (APC) Research Group, School of Mechanical and Design Engineering, University of Portsmouth, Portsmouth PO1 3DJ, UK

**\*** Correspondence: khaled.giasin@port.ac.uk (K.G.); antigoni.barouni@port.ac.uk (A.B.)

**Abstract:** Fibre Metal Laminates (FMLs) consist of layers of metals combined with layers of fibrereinforced composites bonded together to create a laminate. The behaviour of a Fibre Metal Laminate (FML) with natural fibre composites has been investigated in this study with a specific focus on the performance of the laminate under uniaxial tension. The integration of aluminium layers with natural fibre flax/pp layers at different fibre orientations has been numerically modelled and analysed, by investigating the contact interface between natural fibre metal laminates (NFML) using finite elements (FE) implemented in ABAQUS/Explicit. The finite element model was developed by the isotropichardening behaviour of metal layers, the built-in Hashin damage model and cohesive surface-based behaviour for the interface. The results of the simulation included stress–strain response, failure sequences, delamination effect and ultimate tensile strength. It was found that those results are significantly affected by the layup sequence, giving a significant advantage to the unidirectional laminate, when the uniaxial loading is taken into consideration. This advantage is measured as a 41.9% reduction of the ultimate tensile strength when the flax fibres are oriented at [0/90] configuration between the aluminium layers and a 30% reduction when the fibres are oriented at [±45] angles.

**Keywords:** fibre metal laminates; natural fibre composites; progressive damage analysis; failure mechanism; finite element analysis; numerical analysis; tension; delamination

### **1. Introduction**

FMLs have become very common in the aerospace industry during the past decades. It is a hybrid material that is built up from metal layers and fibre-reinforced composite layers, taking advantage of the benefits of the metallic layers, such as their high bearing strength and impact resistance, as well as the good fatigue resistance and high stiffness-to-weight ratio of the composites [1,2]. The main benefit of the FMLs is sourced from the bridging effect between the composite layer and metal layer, which can significantly slow down the fatigue crack growth rate. The bridging effect is taking place between the composite layer and the delamination plane. When the metal layer cracks, the composite layer will continue to hold and reduce the stress intensity [3–6]. Hence, the stress concentration will reduce, which can lead to a slow crack growth rate. This improved fatigue behaviour of FMLs makes them ideal for aircraft structures. Asundi et al. [7] found that using glass fibrereinforced aluminium laminates as a material of the fuselage skin can reduce the weight by 15–25% without compromising the performance. The most successful application is used as the upper fuselage skin material of the Airbus A380. FMLs in aerospace applications might experience constant tensile and compressive loading [8], making the mechanical and failure behaviour of FMLs under tension very critical, which can include matrix cracking, fibre breakage and debonding of layers [9].

Over the years, the fibre-reinforced composite materials have been used in a variety of engineering applications and domains, where the tensile fatigue behaviour is investigated

**Citation:** Song, C.H.; Giasin, K.; Saifullah, A.; Barouni, A. Investigation of the Contact Interface between Natural Fibre Metal Laminates under Tension Using Finite Element Analysis (FEA). *Polymers* **2022**, *14*, 4650. https:// doi.org/10.3390/polym14214650

Academic Editors: R. A. Ilyas, S. M. Sapuan, Emin Bayraktar, Shukur Abu Hassan, Nabil Hayeemasae and Khubab Shaker

Received: 30 September 2022 Accepted: 26 October 2022 Published: 1 November 2022

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2022 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

for glass fibre laminates [10], and the compressive strength is improved on glass fibre composites by winding additional layers around the longitudinal glass fibres [11], and the development of data-driven models and neural network models has been attempted for the prediction of the flexural performance of fibre-reinforced concrete columns [12,13]. Various attempts have taken place in order to investigate the performance of fibre metal laminates under different loading scenarios such as bending, tensile, compression, impact and shear [14]. Khalili et al. [15] investigated the mechanical properties of fibre metal laminates with different lay-ups of laminates and metal layers by three different testing scenarios: bending, impact and tensile. The results indicated that the damage tolerance limit and energy absorption of FMLs are superior to fibre-reinforced composites. Detailed work has been done by Sharma et al. [16] on quasi-static testing of titanium-based glass fibre-reinforced laminates. The tensile behaviour of four different layups was evaluated. The overall thickness of fibre metal laminates was fixed while changing the position and thickness of the metal layers in the stack. From the results, the layup sequence rarely affected the initial modulus of the fibre metal laminates. Moussavi et al. [17] combined classical laminate theory with the elastic-plastic behaviour of the aluminium layer to predict the stress–strain behaviour of fibre metal laminates. The fibre metal laminates model was loaded by uniform tensile force along the fibre direction while no cohesive behaviour was considered between any two layers. They concluded that FMLs with zero orientation fibre layer show improvement in maximum tensile strength. In another experiment, Hashemi et al. [18] investigated the performance of 3D glass fibre-based fibre metal laminates under tensile loading, when the laminates were manufactured under a different set of parameters (i.e., temperature, pressure and time), with the pressure contributing more to the mechanical properties of the laminate.

Numerical simulations are increasingly being used to study the performance of fibre metal laminates. Soltani et al. [19] simulated the tensile behaviour of FMLs under in-plane loading using Finite Element (FE) modelling, proposing a model able to analyse GLARE laminates in structural applications. Sharma et al. [20] investigated the tensile response of FMLs using three different specimens with constant volume but different configurations of metal and composite layers. Stress–strain curves of the FMLs were presented and indicated the failure stage of each specimen. It was stated that the configuration of metal and laminate layers would affect the tensile behaviour of the FML, where the existence of a metal layer between the composite layers demonstrates weaker performance than without a metal layer between the composite layers. A work that discusses the progressive damage and failure mechanisms of open-hole fibre metal laminates was done by Du et al. [21]. The performance of three different carbon fibre-reinforced PEEK prepregs layups, namely unidirectional, cross-ply and quasi-isotropic layups, was investigated. It was found that the failure of the FMLs began from fibre breakage. After the laminates met the yield point, the fibres would break first, followed by the failure of the matrix. The facture of metal layers took place once the delamination entirely occurred between the metal and laminate layer. Furthermore, higher-order shear deformation theory has been implemented to predict the shear strains through the thickness of composite laminated plates, with the use of FE modelling [22] as well as 3D deformations and progressive failures have been predicted for conventional fibre metal laminates using a user subroutine [23].

In the last few years, the production of sustainable products has gained more attention from manufacturers and researchers, leading them to integrate natural fibre composites into different applications [24,25]. Natural fibre composites are lightweight, easier to produce from natural resources and exhibit reduced environmental impact compared to synthetic fibres, making them a very promising replacement to synthetic fibres in demanding applications, such as automotive [26]. Fidelis et al. [27] investigated the tensile strength of several natural fibre composites, including jute, sisal, curaua, coir and piassava, concluding that the curaua fibres have the highest young modulus and tensile strength among other natural fibres. The use of natural fibre-reinforced composites has been increasingly used in semi-structural applications, where impact toughness is important [28], as well as in more

demanding applications with the cycle loading performance of the material is critical [29]. However, the use of natural fibres is still limited to applications where the demand for structural integrity is high and critical, such as aerospace. The combination of natural fibre composites with metallic layers, to provide a sustainable alternative to FMLs, could replace the conventional FMLs in aerospace and improve the carbon footprint of the structures significantly, bringing a novelty to aerospace materials. Several implications can appear during this attempt, with the highest would be the compatibility of the two materials, with a greater focus on the delamination between the metallic and composite layers. Therefore, a preliminary numerical study is essential to understand the behaviour of these novel materials.

This paper aims to perform a finite element analysis for simulating a natural fibre metal laminate (NFML) composite structure and its failure behaviour under tension. Hence, the main focus is to study the damage initiation and damage evolution of the composite structure under tensile loading. In the simulation, cohesive-surface-based behaviour was applied to observe the delamination and the Hashin 2D criterion was involved to capture the damage initiation and evolution in the composite. Three types of layup configurations were considered, such as unidirectional, cross-ply and quasi-isotropic, to study the effect of fibre orientation on the failure modes, which is an important factor affecting the strength of the composite. The results obtained from the simulation such as the stress distribution and damage evolution of each layer will be discussed in the results section.

#### **2. Materials and Methods**

#### *2.1. Metallic Layers*

The metallic layer used in this study is Aluminium 2024-T3. The elastic and plastic behaviour of the metallic layer was modelled by using isotropic hardening. The isotropic hardening data for Al 2024-T3 are listed in Table 1 [20].


**Table 1.** Isotropic hardening data for Aluminium 2024-T3.

#### *2.2. Composite Layers*

Two different composite materials were involved in this study; E-glass/Epoxy composite laminates and flax/polypropylene composite. E-glass/Epoxy composite laminates were modelled in the first part of the simulation to validate the developed FE model using the work of Sharma [20] and the properties are shown in Table 2. The second part of this study investigated the performance of natural fibre flax/polypropylene composite. The material properties used for natural fibre flax/polypropylene composite are listed in Table 3 [30].


**Table 2.** Material properties of the E-glass/Epoxy composite.

#### **Table 3.** Material properties of the flax/polypropylene.


For the prediction of the damage initiation and evolution, a cohesive-surface-based behaviour method was implemented, using the material properties summarised in Table 4 [30].


**Table 4.** Material properties of the cohesive interface.

#### *2.3. Configuration of NFMLs 2.3. Configuration of NFMLs*

To validate the developed FE model in this work, a validation study was initially conducted with already existing literature. The fibre metal laminate with glass fibre composite layers, named GLARE 3/2-0.4, was used for the initial validation study [19]. To validate the developed FE model in this work, a validation study was initially conducted with already existing literature. The fibre metal laminate with glass fibre composite layers, named GLARE 3/2-0.4, was used for the initial validation study [19].

#### 2.3.1. Glare 3/2-0.4 2.3.1. Glare 3/2-0.4

The Glare3/2-0.4 laminate was made of three sheets of aluminium 2024-T3 alloy alternating with two layers of E-glass/Epoxy composite laminates. The schematic diagram of the specimen is shown in Figure 1. Due to the symmetry of the geometry and loading of the specimen, only a quarter of the model was simulated. The total thickness of the specimen is 3.7 mm, while the aluminium thickness is 0.4 mm each and the composite layer is 0.625 mm each. The Glare3/2-0.4 laminate was made of three sheets of aluminium 2024-T3 alloy alternating with two layers of E-glass/Epoxy composite laminates. The schematic diagram of the specimen is shown in Figure 1. Due to the symmetry of the geometry and loading of the specimen, only a quarter of the model was simulated. The total thickness of the specimen is 3.7 mm, while the aluminium thickness is 0.4 mm each and the composite layer is 0.625 mm each.

**Figure 1.** Schematic diagram of GLARE 3/2-0.4. **Figure 1.** Schematic diagram of GLARE 3/2-0.4.

#### 2.3.2. Aluminium/Flax Fibre Metal Laminates 2.3.2. Aluminium/Flax Fibre Metal Laminates

The main work for this study was conducted on a natural fibre metal laminate (NFML), where flax-fibre-reinforced laminates were used combined with Aluminium layers. The material is subjected to tensile loadings at one end of the specimen. The effect of lay-up configuration on fibre metal laminates with flax fibre reinforcement is subsequently investigated in this study. Three different laminate configurations were used for this investigation, where three aluminium layers and two embedded flax-fibre/polypropylene composite layers were used for each configuration. The laminate configurations used include unidirectional, cross-ply and biaxial configurations. The thickness of the aluminium and flax/pp layers is listed in Table 5. The schematics of the fibre metal laminates with different stacking sequences are presented in Figure 2. The size and geometry of the model remain the same as described in ASTM D3039. The main work for this study was conducted on a natural fibre metal laminate (NFML), where flax-fibre-reinforced laminates were used combined with Aluminium layers. The material is subjected to tensile loadings at one end of the specimen. The effect of layup configuration on fibre metal laminates with flax fibre reinforcement is subsequently investigated in this study. Three different laminate configurations were used for this investigation, where three aluminium layers and two embedded flax-fibre/polypropylene composite layers were used for each configuration. The laminate configurations used include unidirectional, cross-ply and biaxial configurations. The thickness of the aluminium and flax/pp layers is listed in Table 5. The schematics of the fibre metal laminates with different stacking sequences are presented in Figure 2. The size and geometry of the model remain the same as described in ASTM D3039.

**Type** 

A [AL/0/0/AL/0/0/AL] FLAX 0.4 0.625 B [AL/0/90/AL/90/0/AL] FLAX 0.4 0.625 C [AL/+45/−45/AL/+45/−45/AL] FLAX 0.4 0.625

**Al Thickness (mm)** 

**Flax/PP Thickness (mm)** 

**Table 5.** Details of fibre metal laminates considered.

**Laminate** 


**Table 5.** Details of fibre metal laminates considered.

**Figure 2.** The various layup of Aluminium/Flax fibre metal laminates. **Figure 2.** The various layup of Aluminium/Flax fibre metal laminates. **Figure 2.** The various layup of Aluminium/Flax fibre metal laminates.

#### **3. FE Modelling Approach 3. FE Modelling Approach 3. FE Modelling Approach**

elements in the geometry.

elements in the geometry.

For the investigation of the tensile behaviour of NFMLs, an FE model was developed using ABAQUS/Explicit, using the geometry as described in the experimental testing standard ASTM D3039. The dog-bone-shaped geometry was used where one-quarter of the geometry was being modelled, due to symmetry reasons, as shown in Figure 3. For the investigation of the tensile behaviour of NFMLs, an FE model was developed using ABAQUS/Explicit, using the geometry as described in the experimental testing standard ASTM D3039. The dog-bone-shaped geometry was used where one-quarter of the geometry was being modelled, due to symmetry reasons, as shown in Figure 3. For the investigation of the tensile behaviour of NFMLs, an FE model was developed using ABAQUS/Explicit, using the geometry as described in the experimental testing standard ASTM D3039. The dog-bone-shaped geometry was used where one-quarter of the geometry was being modelled, due to symmetry reasons, as shown in Figure 3.

FMLs will undergo different types of failure modes during the deformation process, such as delamination between two layers, fibre breakage, fracture of the metal layer, and

FMLs will undergo different types of failure modes during the deformation process, such as delamination between two layers, fibre breakage, fracture of the metal layer, and the debonding process in the laminates layer and matrix cracking [31]. The cohesive sur-

tween two surfaces. The cohesive-surface-based behaviour method uses traction-separation theory to define cohesive interaction between two surfaces without having cohesive

tween two surfaces. The cohesive-surface-based behaviour method uses traction-separation theory to define cohesive interaction between two surfaces without having cohesive

Traction-separation behaviour involves three stages, namely, the linear elastic behaviour, the damage initiation and the damage evolution. Linear elastic behaviour with

Traction-separation behaviour involves three stages, namely, the linear elastic behaviour, the damage initiation and the damage evolution. Linear elastic behaviour with

FMLs will undergo different types of failure modes during the deformation process, such as delamination between two layers, fibre breakage, fracture of the metal layer, and the debonding process in the laminates layer and matrix cracking [31]. The cohesive surfacebased behaviour was used for this numerical modelling to define the interaction between two surfaces. The cohesive-surface-based behaviour method uses traction-separation theory to define cohesive interaction between two surfaces without having cohesive elements in the geometry.

Traction-separation behaviour involves three stages, namely, the linear elastic behaviour, the damage initiation and the damage evolution. Linear elastic behaviour with uncoupled traction separation was used, meaning shear direction cohesive forces are not affected by pure normal separation and normal direction cohesive forces are not affected by pure shear slip. Hence, the uncoupled behaviour relation is written as follows.

$$t = \begin{Bmatrix} t\_n \\ t\_s \\ t\_t \end{Bmatrix} = \begin{bmatrix} \mathbb{K}\_{nn} & 0 & 0 \\ 0 & \mathbb{K}\_{ss} & 0 \\ 0 & 0 & \mathbb{K}\_{tt} \end{bmatrix} \begin{Bmatrix} \delta\_n \\ \delta\_s \\ \delta\_t \end{Bmatrix} \tag{1}$$

where *tn*, *t<sup>s</sup>* , *t<sup>t</sup>* indicate the nominal traction vectors in the orthogonal direction. The value of penalty stiffness for the traction separation behaviour is denoted as *K*, and relative displacement is denoted as *δ*. Both are involved in the orthogonal direction.

Damage initiation represents the beginning of the degradation of the contact surface under loading conditions. The degradation of the material starts when the contact stresses meet the damage initiation criterion defined by the user. A value of 1 in the damage initiation criterion represents the failure of the contact surface. The quadratic stress criterion was chosen for this simulation. It involves a peak value of the contact stress to determine the degradation.

$$
\left(\frac{t\_n}{t\_n^0}\right)^2 + \left(\frac{t\_s}{t\_s^0}\right)^2 + \left(\frac{t\_t}{t\_t^0}\right)^2 = 1\tag{2}
$$

where *t* 0 *n* , *t* 0 *<sup>s</sup>* and *t* 0 *t* represent the interface strength parameter in the orthogonal direction.

Damage evolution indicates the degradation of the material, which is cohesive stiffness after fulfilling the initiation criteria. The damage evolution is according to softening laws. Either displacement or fracture energy can determine the softening laws for simulating the delamination. The Benzeggagh–Kenane fracture criterion is applied to the simulation by assuming the same value of fracture energy in the first and second shear direction (*G C <sup>s</sup>* = *G C t* ). All parameters of the cohesive-surface-based behaviour method used in this study were summarized in Table 2.

To predict the behaviour of composite layers, a constitutive model is used to describe the response of linear elastic behaviour of fibre-reinforced material. The relationship between stress and strain before damage initiates is defined as *σ* = *C<sup>d</sup> ε*. The damaged elasticity matrix replaces the elasticity matrix once damage initiation happens. The damaged elasticity matrix *C<sup>d</sup>* is given below.

$$\mathbf{E}[\mathbf{C}\_{d}] = \begin{bmatrix} (1 - d\_{f})\mathbf{E}\_{1} & \begin{pmatrix} 1 - d\_{f} \end{pmatrix} (1 - d\_{m}) v\_{21} \mathbf{E}\_{1} & \mathbf{0} \\ (1 - d\_{f})(1 - d\_{m}) v\_{12} \mathbf{E}\_{2} & (1 - d\_{m}) \mathbf{E}\_{2} & \mathbf{0} \\ \mathbf{0} & \mathbf{0} & (1 - d\_{s}) \mathbf{G}\_{12} \mathbf{D} \end{bmatrix} \tag{3}$$

$$D = 1 - \left(1 - d\_f\right) (1 - d\_m) v\_{12} v\_{21} \tag{4}$$

$$d\_f = d\_f^t \text{ if } \sigma\_{11} \ge 0, \text{ or } d\_f^c \text{ if } \sigma\_{11} < 0 \tag{5}$$

$$d\_m = d\_m^t \text{ if } \sigma\_{22} \ge 0 \text{, or } d\_m^c \text{ if } \sigma\_{22} < 0 \tag{6}$$

$$d\_s = 1 - \left(1 - d\_f^t\right) \left(1 - d\_f^c\right) \left(1 - d\_m^t\right) \left(1 - d\_m^c\right) \tag{7}$$

where the subscripts *f* , *m*, *s* represent fibre, matrix and shear, respectively, and the subscript d is damage. *E*<sup>1</sup> and *E*<sup>2</sup> are Young's modulus in the fibre and matrix direction, respectively, *G*<sup>12</sup> and *v*<sup>12</sup> are the shear modulus and Poisson's ratio in the X-Y direction.

Equation (3) above defines how the damage of fibre and matrix changes the material's stiffness according to four different failure modes, as described in the 2D Hashin failure criterion. ABAQUS offers the Hashin 2D criterion to predict the four different failure modes of the shell element. Failure of the composite is defined as the element reaching the yield stress and leading to damage occurrences. Maximum stress or strain theory will be needed to determine the initial damage to the matrix or fibre [32]. The purpose of using the Hashin-based damage criterion is to evaluate different failure modes in a different direction [33]. The general form of four different damage initiation modes is defined below.

Fibre tension *σ*<sup>11</sup> ≥ 0

$$F\_f^t = \left(\frac{\sigma\_{11}}{X^t}\right)^2 + \mathfrak{a}\left(\frac{\sigma\_{12}}{S^L}\right)^2\tag{8}$$

Fibre compression *σ*<sup>11</sup> < 0

$$F\_f^\mathbb{C} = \left(\frac{\sigma\_{11}}{X^\mathbb{C}}\right)^2 \tag{9}$$

Matrix tension *σ*<sup>22</sup> ≥ 0

$$F\_m^t = \left(\frac{\sigma\_{22}}{Y^t}\right)^2 + \left(\frac{\sigma\_{12}}{S^L}\right)^2\tag{10}$$

Matrix compression *σ*<sup>22</sup> < 0

$$F\_m^\mathbb{C} = \left(\frac{\sigma\_{22}}{2S^t}\right)^2 + \left[\left(\frac{Y^\mathbb{C}}{2S^T}\right)^2 - 1\right] \left(\frac{\sigma\_{22}}{Y^\mathbb{C}}\right) + \left(\frac{\sigma\_{12}}{S^L}\right)^2\tag{11}$$

where *F t f* , *F C f* indicate failure mode of fibre in tension and compression; *F t <sup>m</sup>*, *F C <sup>m</sup>* indicate the failure mode of the matrix in tension and compression. *σ*11, *σ*<sup>22</sup> and *σ*<sup>12</sup> represent the effective stress tensor. *X <sup>t</sup>* and *X <sup>C</sup>* are the tensile strength and compressive strength of laminate in the longitudinal direction while *Y <sup>t</sup>* and *Y <sup>C</sup>* are the tensile strength and compressive strength of laminate in the transverse direction. *S <sup>L</sup>* and *S T* , respectively, represent the shear strength of the laminate in the longitudinal and transverse directions.

Due to the symmetry of the geometry and loading of the specimen, only a quarter of the specimen was modelled. The FE model is subjected to tensile loadings at one end of the specimen, whereas encastre boundary conditions were applied on the other end. The built-in fibre-reinforced material damage model will be used in this simulation. A mesh convergence study was performed to determine the minimum element size. ABAQUS offers general contact in ABAQUS/Explicit solver, where each interface of two layers was defined. For the interface between metal and composite layers, the friction of 0.3 was used, while a friction value of 0.5 was used as input for the interface between the composite layers [34]. The element type used for the metal layer was eight-node solid brick elements (C3D8R) and for the composite layer was the eight-node continuum shell element (SC8R). ABAQUS explicit option was selected for this simulation as it is more suitable for highly nonlinear analysis and can reduce the required time. In the GLARE 3/2-0.4 FE model, no damage evolution was applied. However, the model for aluminium/flax fibre metal laminates does involve a ductile damage model. Once the damage initiation has been satisfied, Hillerborg's fracture energy damage evolution criterion will be introduced. This ductile damage model includes the fracture strain and fracture energy to modify the failure of metal.

#### **4. Results and Discussion**

To initially validate the developed FE model, the strain–stress curves were compared to the results presented in [20] for three different material types: (a) a 0.4 mm thick aluminium plate, as shown in Figure 4a, (b) a cross-ply GFRP laminate, as shown in Figure 4b, (c) a GLARE 3/2-0.4 laminate, as shown in Figure 4c. All three curves exhibit very good agreement with the current FE model. Also, by comparing the ultimate tensile strength, the percentage error is 0.958%. As can be seen, all curves agree very well, especially around the initiation of delamination, a fact which enhances the use of the current FE model for further simulations. *Polymers* **2022**, *14*, x FOR PEER REVIEW 9 of 18

**Figure 4.** Stress–strain curves of Present FE model and Sharma et al. FE model [20] for (**a**) AL-2024 T3 alloy; (**b**) of [0°/90°/90°/0°] GFRP; (**c**) GLARE 3/2-0.4. **Figure 4.** Stress–strain curves of Present FE model and Sharma et al. FE model [20] for (**a**) AL-2024 T3 alloy; (**b**) of [0◦/90◦/90◦/0◦ ] GFRP; (**c**) GLARE 3/2-0.4.

The stress–strain curves, failure sequences, delamination effect, ultimate tensile strength and effect of composite layup configurations will be discussed in this section. The stress–strain curves, failure sequences, delamination effect, ultimate tensile strength and effect of composite layup configurations will be discussed in this section.

The stress–strain response for each configuration A, B and C, under tensile loading is described in Figures 5 and 6. According to Figure 5, laminate A with unidirectional fibre orientation has the highest tensile strength, followed by laminate C, and laminate B exhibits the lowest UTS among the chosen laminates. Specimen A has the highest tensile strength mainly due to the fibre orientation being aligned to the tensile load in the longitudinal direction. Figure 5 shows a significant difference in ultimate tensile strength when changing the layup of the lamina. For example, laminate B with a cross-ply configuration demonstrates a reduction of 41% of ultimate tensile strength compared to laminate A. The stress–strain response for each configuration A, B and C, under tensile loading is described in Figures 5 and 6. According to Figure 5, laminate A with unidirectional fibre orientation has the highest tensile strength, followed by laminate C, and laminate B exhibits the lowest UTS among the chosen laminates. Specimen A has the highest tensile strength mainly due to the fibre orientation being aligned to the tensile load in the longitudinal direction. Figure 5 shows a significant difference in ultimate tensile strength when changing the layup of the lamina. For example, laminate B with a cross-ply configuration demonstrates a reduction of 41% of ultimate tensile strength compared to laminate A.

Laminate A Laminate B Laminate C

**Type of Laminate Al/Flax laminate**

**Ultimate tensile strength (MPa)**

0 0.05 0.1 0.15 0.2 0.25

 Present FE model FE model (Sharma et.al, 2017)

Strain

Stress [MPa]

T3 alloy; (**b**) of [0°/90°/90°/0°] GFRP; (**c**) GLARE 3/2-0.4.

Stress [MPa]

**Figure 5.** Comparison of ultimate tensile strength of the three laminate configurations. **Figure 5.** Comparison of ultimate tensile strength of the three laminate configurations.

(**a**) (**b**)

Present model

FE model (Sharma et al, 2017)

Stress [MPa]

0 0.005 0.01 0.015 0.02 0.025

 Present FE model FE model (Sharma et.al, 2017)

Strain

(**c**) **Figure 4.** Stress–strain curves of Present FE model and Sharma et al. FE model [20] for (**a**) AL-2024

0 0.02 0.04 0.06 0.08 0.1 0.12 0.14

Strain

strength and effect of composite layup configurations will be discussed in this section.

The stress–strain curves, failure sequences, delamination effect, ultimate tensile

The stress–strain response for each configuration A, B and C, under tensile loading is described in Figures 5 and 6. According to Figure 5, laminate A with unidirectional fibre orientation has the highest tensile strength, followed by laminate C, and laminate B exhibits the lowest UTS among the chosen laminates. Specimen A has the highest tensile strength mainly due to the fibre orientation being aligned to the tensile load in the longitudinal direction. Figure 5 shows a significant difference in ultimate tensile strength when changing the layup of the lamina. For example, laminate B with a cross-ply configuration demonstrates a reduction of 41% of ultimate tensile strength compared to laminate A.

**Figure 6.** Stress–strain curves of (**a**) laminate A; (**b**) laminate B; (**c**) laminate C. **Figure 6.** Stress–strain curves of (**a**) laminate A; (**b**) laminate B; (**c**) laminate C.

Due to the same metal distribution and metal volume, the curves of the three specimens in Figure 6 look similar but fail at different strain values. For laminate A, at the beginning of loading, the maximum stresses are mainly concentrated on the aluminium layers and only appear on the edges of the composite layers. The stress slowly grows into the whole structure and is distributed evenly. This is because of the characteristics of 0° degree fibre orientation which increases the tensile loading capability of the laminate. At this point, the longitudinal stresses on the laminate will continue to rise and reach the ultimate tensile strength. Once the ultimate tensile strength is reached, cracks will appear in the metal layer and the whole laminate will fail. Due to the same metal distribution and metal volume, the curves of the three specimens in Figure 6 look similar but fail at different strain values. For laminate A, at the beginning of loading, the maximum stresses are mainly concentrated on the aluminium layers and only appear on the edges of the composite layers. The stress slowly grows into the whole structure and is distributed evenly. This is because of the characteristics of 0◦ -degree fibre orientation which increases the tensile loading capability of the laminate. At this point, the longitudinal stresses on the laminate will continue to rise and reach the ultimate tensile strength. Once the ultimate tensile strength is reached, cracks will appear in the metal layer and the whole laminate will fail.

Figure 7 presents the results for the damage evolution in both matrix and fibres in tension for each of the composite layers of laminate A at 2.6 mm displacement, with the Figure 7 presents the results for the damage evolution in both matrix and fibres in tension for each of the composite layers of laminate A at 2.6 mm displacement, with the

naming of the flax layers following the convention of Figure 2. The damage evolution started with the matrix cracking that was first initiated at both 0° flax composite layers, as

capable of holding stress due to their characteristics. After the matrix failure, the fibre starts to break. This led to the failure of the specimen. As shown from the stress distribution in Figure 8, at 2.8 mm displacement, the middle aluminium layer has already failed, and cracking starts to appear at the 0° Flax-2 layer too. For the interfaces between the layers, which are numbered in Figure 9, the delamination at the interface between each layer after the failure of the specimen is shown in Figure 10. The red-coloured areas in these contours represent the highly delaminated areas of each interface (with a value of 1.0 of the damage indicator) whereas the blue-coloured areas indicate no delamination (value 0.0 of the damage indicator). The edges of the layer between the top aluminium and bottom aluminium layers have less delamination than those in between the two com-

posite layers and mid-aluminium layers.

naming of the flax layers following the convention of Figure 2. The damage evolution started with the matrix cracking that was first initiated at both 0◦ flax composite layers, as shown in Figure 7. At the same time, the fibres of both 0◦ flax composite layers are still capable of holding stress due to their characteristics. After the matrix failure, the fibre starts to break. This led to the failure of the specimen. As shown from the stress distribution in Figure 8, at 2.8 mm displacement, the middle aluminium layer has already failed, and cracking starts to appear at the 0◦ Flax-2 layer too. For the interfaces between the layers, which are numbered in Figure 9, the delamination at the interface between each layer after the failure of the specimen is shown in Figure 10. The red-coloured areas in these contours represent the highly delaminated areas of each interface (with a value of 1.0 of the damage indicator) whereas the blue-coloured areas indicate no delamination (value 0.0 of the damage indicator). The edges of the layer between the top aluminium and bottom aluminium layers have less delamination than those in between the two composite layers and mid-aluminium layers. *Polymers* **2022**, *14*, x FOR PEER REVIEW 11 of 18 *Polymers* **2022**, *14*, x FOR PEER REVIEW 11 of 18

0° FLAX-2

0° FLAX-2

0° FLAX-1

0° FLAX-1

**Figure 7.** Damage evolution for fibres and matrix of top composite layers of Laminate A at 2.6 mm displacement. **Figure 7.** Damage evolution for fibres and matrix of top composite layers of Laminate A at 2.6 mm displacement. displacement.

**Figure 8.** Stress distribution of the four top layers of laminate A at 2.8 mm displacement. **Figure 8.** Stress distribution of the four top layers of laminate A at 2.8 mm displacement. **Figure 8.** Stress distribution of the four top layers of laminate A at 2.8 mm displacement.

0° FLAX-1 DAMAGEFT

displacement.

0° FLAX-1 DAMAGEMT

**Figure 7.** Damage evolution for fibres and matrix of top composite layers of Laminate A at 2.6 mm

AL-top 0° FLAX-1 0° FLAX-2 AL-Middle

**Figure 8.** Stress distribution of the four top layers of laminate A at 2.8 mm displacement.

0° FLAX-2 DAMAGEFT

0° FLAX-2 DAMAGEMT

**Figure 10.** Delamination of each layer of laminate A after failure; (**a**) 1st interface; (**b**) 2nd interface; (**c**) 3rd interface; (**d**) 4th interface; (**e**) 5th interface; (**f**) 6th interface. **Figure 10.** Delamination of each layer of laminate A after failure; (**a**) 1st interface; (**b**) 2nd interface; (**c**) 3rd interface; (**d**) 4th interface; (**e**) 5th interface; (**f**) 6th interface.

The stress distribution of laminate B is mainly focused on the 0° ply and aluminium layers as shown in Figure 11 at a displacement of 0.5 mm. The longitudinal stress carried by the aluminium layer and 0° ply is around 371 MPa and 216 MPa, respectively, while the 90° ply only carried 1.6 MPa of stress, due to the matrix cracking occurring at the 90° ply during the early stage of tension loading. With further loading, the propagation of matrix cracking will be distributed across the 90° ply. The fibres belonging to the 0° ply will start to break, and the direction of propagation is perpendicular to the fibre direction, as seen in Figure 12. The stress distribution of laminate B is mainly focused on the 0◦ ply and aluminium layers as shown in Figure 11 at a displacement of 0.5 mm. The longitudinal stress carried by the aluminium layer and 0◦ ply is around 371 MPa and 216 MPa, respectively, while the 90◦ ply only carried 1.6 MPa of stress, due to the matrix cracking occurring at the 90◦ ply during the early stage of tension loading. With further loading, the propagation of matrix cracking will be distributed across the 90◦ ply. The fibres belonging to the 0◦ ply will start to break, and the direction of propagation is perpendicular to the fibre direction, as seen in Figure 12.

The edge delamination formation occurred first between the 90° flax composite layer and the mid-aluminium layer of laminate B (Figure 13c). Unlike laminate A, the delamination happens randomly of the edges during the early stage. After the delamination occurred on the edges of the 90° flax composite and mid-aluminium layers, it started to distribute into adjacent layers. The 1st and 6th interfaces have delamination on the end edge The edge delamination formation occurred first between the 90◦ flax composite layer and the mid-aluminium layer of laminate B (Figure 13c). Unlike laminate A, the delamination happens randomly of the edges during the early stage. After the delamination occurred on the edges of the 90◦ flax composite and mid-aluminium layers, it started to distribute into adjacent layers. The 1st and 6th interfaces have delamination on the end

of the laminate while other interfaces do not, as shown in Figure 13. Again here, the redcoloured areas in these contours represent the highly delaminated areas of each interface

delamination (value 0.0 of the damage indicator). The area of delamination across interfaces 1 and 6, that is, between Al and 0° composite layer layers, is larger than the delamination observed across interfaces 3 and 4, that is, between Al and 90° composite layers. It can be concluded that the primary failure starts from matrix cracking of 90° ply, and fibre tension damage in the 0° layer occurred during the early stage of loading. Once the matrix cracking of 90° ply started, the 90° ply lost its ability to withstand the stress. Then the fibre of the 0° ply and aluminium layer will be responsible for holding the stress. Once it

reached the ultimate strain, the specimen failed.

edge of the laminate while other interfaces do not, as shown in Figure 13. Again here, the red-coloured areas in these contours represent the highly delaminated areas of each interface (with a value of 1.0 of the damage indicator) whereas the blue-coloured areas indicate no delamination (value 0.0 of the damage indicator). The area of delamination across interfaces 1 and 6, that is, between Al and 0◦ composite layer layers, is larger than the delamination observed across interfaces 3 and 4, that is, between Al and 90◦ composite layers. It can be concluded that the primary failure starts from matrix cracking of 90◦ ply, and fibre tension damage in the 0◦ layer occurred during the early stage of loading. Once the matrix cracking of 90◦ ply started, the 90◦ ply lost its ability to withstand the stress. Then the fibre of the 0◦ ply and aluminium layer will be responsible for holding the stress. Once it reached the ultimate strain, the specimen failed. *Polymers* **2022**, *14*, x FOR PEER REVIEW 13 of 18 *Polymers* **2022**, *14*, x FOR PEER REVIEW 13 of 18

**Figure 11.** Stress distribution of each top layer of laminate B at 0.5 mm displacement and. **Figure 11.** Stress distribution of each top layer of laminate B at 0.5 mm displacement and. **Figure 11.** Stress distribution of each top layer of laminate B at 0.5 mm displacement and.

**Figure 12.** Damage evolution of each composite layer of the laminate B at 2 mm displacement. **Figure 12.** Damage evolution of each composite layer of the laminate B at 2 mm displacement. **Figure 12.** Damage evolution of each composite layer of the laminate B at 2 mm displacement.

(**a**) (**b**)

(**a**) (**b**)

**Figure 12.** Damage evolution of each composite layer of the laminate B at 2 mm displacement.

AL-top 0° FLAX-1 90° FLAX-2 AL-Middle

**Figure 11.** Stress distribution of each top layer of laminate B at 0.5 mm displacement and.

90° FLAX-2 DAMAGEFT

90° FLAX-2 DAMAGEMT

0° FLAX-1 DAMAGEMT

0° FLAX-1 DAMAGEFT

**Figure 13.** Delamination of each layer of laminate B after failure; (**a**) 1st interface; (**b**) 2nd interface; (**c**) 3rd interface; (**d**) 4th interface; (**e**) 5th interface; (**f**) 6th interface. **Figure 13.** Delamination of each layer of laminate B after failure; (**a**) 1st interface; (**b**) 2nd interface; (**c**) 3rd interface; (**d**) 4th interface; (**e**) 5th interface; (**f**) 6th interface.

The stress distribution of laminate C follows the fibre orientation of the composite layer. Figure 14 shows the displacement at 3 mm, where the stress concentration of the 45° flax layer and the −45° flax layer is aligned to the fibre direction. Figure 14 shows that the stresses at the aluminium layers are concentrated around the edges. It is worth observing that, when the angle of the principal loading direction is more than 45 degrees compared to the fibre orientation, the metal layers are dominating the load carrying for the whole FML, which can be also verified from Figure 14, with the stresses developed on the Al layers being around 460 MPa, as opposed to circa 205 MPa for the composite layers. The results presented here show that the aluminium layers carry almost double the magnitude of stress under tensile loading compared with the composite layers. Figure 15 shows the different damage modes developed on the flax layers of the NFML. It is evident that the matrix cracking is the first model to be observed and is extended along the ±45° direction, as expected. The stress distribution of laminate C follows the fibre orientation of the composite layer. Figure 14 shows the displacement at 3 mm, where the stress concentration of the 45◦ flax layer and the −45◦ flax layer is aligned to the fibre direction. Figure 14 shows that the stresses at the aluminium layers are concentrated around the edges. It is worth observing that, when the angle of the principal loading direction is more than 45 degrees compared to the fibre orientation, the metal layers are dominating the load carrying for the whole FML, which can be also verified from Figure 14, with the stresses developed on the Al layers being around 460 MPa, as opposed to circa 205 MPa for the composite layers. The results presented here show that the aluminium layers carry almost double the magnitude of stress under tensile loading compared with the composite layers. Figure 15 shows the different damage modes developed on the flax layers of the NFML. It is evident that the matrix cracking is the first model to be observed and is extended along the ±45◦ direction, as expected.

Unlike laminate B, the delamination of laminate C does not have sequences. It rapidly grows on the edge of the specimen. Figure 16 shows that the delamination occurs at the failure area and follows the 45° fibre orientation pattern. The red-coloured areas in these contours represent the highly delaminated areas of each interface (with a value of 1.0 of the damage indicator) whereas the blue-coloured areas indicate no delamination (value 0.0 of the damage indicator). The damage evolution of laminate C, starting with matrix damage, occurred at both 45° and −45° flax layers. Fibre breakages happened when the specimen failed. The failure of the specimen also follows the fibre orientation. Unlike laminate B, the delamination of laminate C does not have sequences. It rapidly grows on the edge of the specimen. Figure 16 shows that the delamination occurs at the failure area and follows the 45◦ fibre orientation pattern. The red-coloured areas in these contours represent the highly delaminated areas of each interface (with a value of 1.0 of the damage indicator) whereas the blue-coloured areas indicate no delamination (value 0.0 of the damage indicator). The damage evolution of laminate C, starting with matrix damage, occurred at both 45◦ and −45◦ flax layers. Fibre breakages happened when the specimen failed. The failure of the specimen also follows the fibre orientation.

AL-top +45° FLAX-1 −45° FLAX-2 AL-Mid

**Figure 14.** Stress distribution of each top layer of laminate C at 3 mm displacement.

specimen failed. The failure of the specimen also follows the fibre orientation.

(**c**) (**d**)

 (**e**) (**f**) **Figure 13.** Delamination of each layer of laminate B after failure; (**a**) 1st interface; (**b**) 2nd interface;

The stress distribution of laminate C follows the fibre orientation of the composite layer. Figure 14 shows the displacement at 3 mm, where the stress concentration of the 45° flax layer and the −45° flax layer is aligned to the fibre direction. Figure 14 shows that the stresses at the aluminium layers are concentrated around the edges. It is worth observing that, when the angle of the principal loading direction is more than 45 degrees compared to the fibre orientation, the metal layers are dominating the load carrying for the whole FML, which can be also verified from Figure 14, with the stresses developed on the Al layers being around 460 MPa, as opposed to circa 205 MPa for the composite layers. The results presented here show that the aluminium layers carry almost double the magnitude of stress under tensile loading compared with the composite layers. Figure 15 shows the different damage modes developed on the flax layers of the NFML. It is evident that the matrix cracking is the first model to be observed and is extended along the ±45°

Unlike laminate B, the delamination of laminate C does not have sequences. It rapidly grows on the edge of the specimen. Figure 16 shows that the delamination occurs at the failure area and follows the 45° fibre orientation pattern. The red-coloured areas in these contours represent the highly delaminated areas of each interface (with a value of 1.0 of the damage indicator) whereas the blue-coloured areas indicate no delamination (value 0.0 of the damage indicator). The damage evolution of laminate C, starting with matrix damage, occurred at both 45° and −45° flax layers. Fibre breakages happened when the

(**c**) 3rd interface; (**d**) 4th interface; (**e**) 5th interface; (**f**) 6th interface.

direction, as expected.

**Figure 14.** Stress distribution of each top layer of laminate C at 3 mm displacement. **Figure 14.** Stress distribution of each top layer of laminate C at 3 mm displacement.

**Figure 15.** Damage evolution of each top composite layer of specimen C at 1 mm displacement. **Figure 15.** Damage evolution of each top composite layer of specimen C at 1 mm displacement.

(**a**) (**b**)

(**e**) (**f**) **Figure 16.** Delamination of each layer of laminate C after failure; (**a**) 1st interface; (**b**) 2nd interface;

In summary, the three laminate configurations with different layup orientations are being investigated in this section when loaded under tension, where it was found that the stress distributions of each laminate are greatly affected by fibre orientation. For example, all layers of laminate A are capable of carrying load due to the longitudinal direction, while the 90° flax layer of laminate B does not contribute to the longitudinal loading. In addition, the stress distribution of laminate C is along either 45° or −45° fibre orientation. The maximum stresses developed in each laminate are shown in Figure 17, where the first 4 layers are presented, that is, the top and middle aluminium layers as well as the flax

(**c**) 3rd interface; (**d**) 4th interface; (**e**) 5th interface; (**f**) 6th interface.

45° FLAX-1 DAMAGEFT

45° FLAX-1 DAMAGEMT

−45° FLAX-2 DAMAGEFT

−45° FLAX-2 DAMAGEMT

**Figure 15.** Damage evolution of each top composite layer of specimen C at 1 mm displacement.

**Figure 16.** Delamination of each layer of laminate C after failure; (**a**) 1st interface; (**b**) 2nd interface; (**c**) 3rd interface; (**d**) 4th interface; (**e**) 5th interface; (**f**) 6th interface. **Figure 16.** Delamination of each layer of laminate C after failure; (**a**) 1st interface; (**b**) 2nd interface; (**c**) 3rd interface; (**d**) 4th interface; (**e**) 5th interface; (**f**) 6th interface.

In summary, the three laminate configurations with different layup orientations are being investigated in this section when loaded under tension, where it was found that the stress distributions of each laminate are greatly affected by fibre orientation. For example, all layers of laminate A are capable of carrying load due to the longitudinal direction, while the 90° flax layer of laminate B does not contribute to the longitudinal loading. In addition, the stress distribution of laminate C is along either 45° or −45° fibre orientation. The maximum stresses developed in each laminate are shown in Figure 17, where the first 4 layers are presented, that is, the top and middle aluminium layers as well as the flax In summary, the three laminate configurations with different layup orientations are being investigated in this section when loaded under tension, where it was found that the stress distributions of each laminate are greatly affected by fibre orientation. For example, all layers of laminate A are capable of carrying load due to the longitudinal direction, while the 90◦ flax layer of laminate B does not contribute to the longitudinal loading. In addition, the stress distribution of laminate C is along either 45◦ or −45◦ fibre orientation. The maximum stresses developed in each laminate are shown in Figure 17, where the first 4 layers are presented, that is, the top and middle aluminium layers as well as the flax layers in between (Figure 9), with different orientations for each laminate. It is evident from this bar chart that both the aluminium layers are carrying most of the load in all three laminates, with the 0◦ flax layers being mostly loaded, as compared to the other orientations. It is also interesting to observe that the 90◦ fibre orientation carries the lowest load, due to being perpendicular to the loading direction. Delamination often occurs on the edges of the structure and the breakage area but only the interface between the 0◦ fibre and aluminium layers shows delamination on the fixed end side. *Polymers* **2022**, *14*, x FOR PEER REVIEW 16 of 18 layers in between (Figure 9), with different orientations for each laminate. It is evident from this bar chart that both the aluminium layers are carrying most of the load in all three laminates, with the 0° flax layers being mostly loaded, as compared to the other orientations. It is also interesting to observe that the 90° fibre orientation carries the lowest load, due to being perpendicular to the loading direction. Delamination often occurs on the edges of the structure and the breakage area but only the interface between the 0° fibre and aluminium layers shows delamination on the fixed end side.

**Figure 17.** Maximum stress comparison between all three laminate types for the first four layers of each laminate**.**  formance of the composite laminate, with the unidirectional laminate performing better **Figure 17.** Maximum stress comparison between all three laminate types for the first four layers of each laminate.

inate configurations using an FE model developed in Abaqus. The importance of understanding fully the behaviour of these novel materials is essential, leading to their wider adoption into secondary and primary engineering applications. The combined benefit of the proposed NFMLs in terms of their lower carbon footprint and their overall reduced CO2 emissions has motivated this study to provide further insight into their performance. The approach used in this paper comprises a selected set of laminate configurations combining Aluminium layers with flax fibre-reinforced composite layers at different fibre orientations, in order to investigate the effect of the layup on the tensile behaviour of the material. The study is particularly focused on the contact interface between the natural fibre composite and the metal layers under tension loading. The finite element modelling approach is meso-level, which means each layer is modelled individually. The meso-level modelling approach is sufficient to analyse the behaviour of layers and the interaction between them. The developed FE model used the Hashin-2D built-in damage of composite layer and cohesive surface-based criterion in Abaqus. The damage in the interface and failure mechanisms of the fibre metal laminates was discussed in the results part of aluminium/flax fibre metal laminates. The damage to the interface, such as delamination, was also discussed and related to the fibre orientation. Three different types of layup sequences were investigated in this study and showed a different behaviour under tension loading, with the Laminate A exhibiting a higher toughness in tension, due to the fibre orientation. As expected, the fibre orientation plays a critical role in the strength and per-

**5. Conclusions** 

#### **5. Conclusions**

In this paper, the performance of fibre metal laminates reinforced with natural fibre composite layers (NFMLs) under tension is investigated and studied under different laminate configurations using an FE model developed in Abaqus. The importance of understanding fully the behaviour of these novel materials is essential, leading to their wider adoption into secondary and primary engineering applications. The combined benefit of the proposed NFMLs in terms of their lower carbon footprint and their overall reduced CO<sup>2</sup> emissions has motivated this study to provide further insight into their performance. The approach used in this paper comprises a selected set of laminate configurations combining Aluminium layers with flax fibre-reinforced composite layers at different fibre orientations, in order to investigate the effect of the layup on the tensile behaviour of the material. The study is particularly focused on the contact interface between the natural fibre composite and the metal layers under tension loading. The finite element modelling approach is meso-level, which means each layer is modelled individually. The meso-level modelling approach is sufficient to analyse the behaviour of layers and the interaction between them. The developed FE model used the Hashin-2D built-in damage of composite layer and cohesive surface-based criterion in Abaqus. The damage in the interface and failure mechanisms of the fibre metal laminates was discussed in the results part of aluminium/flax fibre metal laminates. The damage to the interface, such as delamination, was also discussed and related to the fibre orientation. Three different types of layup sequences were investigated in this study and showed a different behaviour under tension loading, with the Laminate A exhibiting a higher toughness in tension, due to the fibre orientation. As expected, the fibre orientation plays a critical role in the strength and performance of the composite laminate, with the unidirectional laminate performing better under tension. However, the delamination in the interfaces of the unidirectional laminate (Figure 10) exhibits a greater extend compared to the cross-ply laminate (Figure 13). The debonding between the metallic and the composite layers is of great importance and it's evident from the stress distribution analysis of all three laminates examined in the present study that the aluminium layers are mainly responsible for carrying a greater load compared to the composite layers.

The limitations observed in this study are related to the prediction of the full stress state of the laminate, due to the use of the 2D Hashin model. However, the results produced are of great importance, as they offer an initial understanding of the performance of the novel materials. The feasibility of this study could be further improved in a future paper, to accommodate a 3D constitutive model for the full prediction of the behaviour of such NFMLs. Only one type of natural fibre composite was tested in this study. A comparative study is required in configuration with two or more natural fibre composites within the fibre metal laminates. Additionally, multiple alternative metal materials can also be applied, which can bring in more room for exploring a new combination of fibremetal laminates. Lastly, the feasibility of applying natural fibre composite to the aerospace structural component is also worth investigating.

**Author Contributions:** Conceptualization, A.B.; methodology, A.B. and C.H.S.; software, C.H.S.; validation, C.H.S.; formal analysis, C.H.S. and A.B.; investigation, C.H.S.; data curation, C.H.S.; writing—original draft preparation, C.H.S.; writing—review and editing, A.B., K.G. and A.S.; visualization, C.H.S.; supervision, A.B. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research received no external funding.

**Institutional Review Board Statement:** Not applicable.

**Data Availability Statement:** The data used in the manuscript can be found in the work of Aransáez Ortega, M. Numerical Simulation of the Impact Response of Natural Fibre Composites. 2019. Available online: https://trepo.tuni.fi/handle/123456789/27698. Any other data regarding the FE model can be requested by contacting the corresponding author Antigoni Barouni.

**Conflicts of Interest:** The authors declare no conflict of interest.

## **References**


## *Article* **High-Efficiency Carbon Fiber Recovery Method and Characterization of Carbon FIBER-Reinforced Epoxy/4,4**0 **-Diaminodiphenyl Sulfone Composites**

**Yong-Min Lee 1,2, Kwan-Woo Kim 1,\* and Byung-Joo Kim 3,\***

	- Jeonju 54896, Republic of Korea
	- Tel.: +82-63-219-3713 (K.-W.K.); +82-63-220-3293 (B.-J.K.)

**Abstract:** Globally, the demand for carbon fiber-reinforced thermosetting plastics for various applications is increasing. As a result, the amount of waste from CFRPs is increasing every year, and the EU Council recommends recycling and reuse of CFRPs. Epoxy resin (EP) is used as a matrix for CFRPs, and amine hardeners are mainly used. However, no research has been conducted on recycling EP/4,4'-diaminodiphenyl sulfone (DDS)-based CFRP. In this study, the effect of steam and air pyrolysis conditions on the mechanical properties of re-cycled carbon fiber (r-CF) recovered from carbon fiber-reinforced thermosetting (epoxy/4,40 -diaminodiphenyl sulfone) plastics (CFRPs) was investigated. Steam pyrolysis enhanced resin degradation relative to N<sup>2</sup> . The tensile strength of the recovered r-CF was reduced by up to 35.12% due to oxidation by steam or air. However, the interfacial shear strength (IFSS) tended to increase by 9.18%, which is considered to be due to the increase in functional groups containing oxygen atoms and the roughness of the surface due to oxidation. The recycling of CFRP in both a steam and an air atmosphere caused a decrease in the tensile strength of r-CF. However, they were effective methods to recover r-CF that had a clean surface and increased IFSS.

**Keywords:** carbon fiber-reinforced plastic (CFRP); recycling; epoxy; DDS; mechanical properties

### **1. Introduction**

Carbon fiber-reinforced plastics (CFRPs) are high-performance composite materials composed of carbon fibers and thermosetting or thermoplastic polymer matrices [1–3]. A global goal is to achieve net-zero carbon dioxide emissions by 2050 [4]. Therefore, studies on improving fuel efficiency have been conducted in the fields of aerospace [5,6], ships [7,8], and automobiles [9,10] by building vehicle bodies and parts with CFRP, which is light, high intensity, and high elasticity for light weight [11–13]. As CFRP can replace metals due to the properties mentioned above, as well as the advantages of excellent corrosion resistance [14,15] and durability [16], demand has also increased in structural applications such as renewable energy [17,18], sports [19], and civil engineering and architecture [20–22]. As the production of and demand for CFRP increases, the amount of waste also increases yearly [23–25]. In the European Union, automobiles generate 8-9 million tons of waste each year at the end of their lifespan. Accordingly, the European Parliament and the Council adopted a directive on end-of-life vehicles in 1997 to address this issue. This directive requires that waste be reused and recovered. Thus, giving priority to reuse and recycling is the basic principle of the directive [26]. Meanwhile, delivery of approximately 40,000 aircraft is expected by 2038, with an annual air traffic growth rate of at least 4.3%. Of these aircraft, 56% are expected to accommodate market growth, and the remaining

**Citation:** Lee, Y.-M.; Kim, K.-W.; Kim, B.-J. High-Efficiency Carbon Fiber Recovery Method and Characteri zation of Carbon FIBER-Reinforced Epoxy/4,40 -Diaminodiphenyl Sulfone Composites. *Polymers* **2022**, *14*, 5304. https://doi.org/10.3390/ polym14235304

Academic Editors: R.A. Ilyas, S.M. Sapuan, Emin Bayraktar, Shukur Abu Hassan, Nabil Hayeemasae and Khubab Shaker

Received: 14 November 2022 Accepted: 2 December 2022 Published: 4 December 2022

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2022 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

deliveries are expected to replace approximately 75% of current aircraft over the next 20 years. This will result in the retirement of approximately 19,000 commercial aircraft over the next 18 years. Due to regulations, OEMs consider the end of life of an aircraft at the design stage, and research on aircraft recycling is underway in industry and academia [27]. CFRP is used in 53% of the Airbus a350 XWB, and in 50% of the Boeing 787 [28]. The recent launch of clean aviation to achieve climate neutrality has driven the goal of replacing 75% of existing civil aircraft with hydrogen- and electric-powered aircraft [29]. By 2050, the aviation industry is expected to collect approximately 500,000 tons of end-of-life (EOL) waste from CFRP [30]. CFRP has been disposed of by landfill in the past. However, disposal by landfill is restricted or prohibited because it causes environmental problems due to the lifespan of polymers [31–33]. Therefore, various recycling methods for CFRP have recently been developed. There are three typical recycling methods for CFRP: mechanical, chemical, and thermal. Mechanical recycling is the simple process of physically crushing CFRP. However, CF and resin are not separated during the crushing, which is a drawback, and long fibers are broken during the grinding process, which results in poor physical properties [34]. The chemical method is a method that uses a solvent to decompose important chemical bonds of the used thermosetting resin. Acids [35–37], hydrogen peroxide [34,38], and supercritical fluids [39–41] are used as solvents. However, acids take a long time to decompose CFRP, and hydrogen peroxide is difficult to manage due to its high reactivity. In addition, although supercritical fluids have a high CF recovery rate and can recover CF with a clean surface and excellent physical properties, high temperatures and high-pressure conditions are required, and the cost of r-CF recovery is high. In addition, large-capacity and continuous processes are difficult, and environmental problems may occur due to processing with solvents. Conventional thermal recycling is a simple and continuous process that decomposes resin and recovers CF at a high temperature of approximately 500 ◦C or higher in N<sup>2</sup> or air atmosphere. However, a large amount of energy is consumed, and undecomposed impurities generated during the process are adsorbed on the CF surface, leaving a residue. In addition, pyrolysis produces CO and CO2, which can lead to environmental problems [42]. To compensate for these shortcomings, Kim et al. suggested a decomposition method using superheated steam in 2017 [23]. Superheated steam can uniformly apply heat to CFRP and penetrate the resin to decompose rapidly. In addition, unlike other pyrolysis methods, it is possible to recover in the form of tar and thus generate less CO2. Epoxy is the most commonly used thermosetting resin in CFRP production, and amine curing agents are used primarily for epoxy curing. Among the commonly used 4,40 -diaminodiphenyl methane (DDM) and 4,40 -diaminodiphenyl sulfone (DDS), many studies have been conducted on recycling methods of EP/DDM-based CFRP, but EP/DDSbased CFRP recycling methods have not been studied. In this study, EP/DDS-based CFRPs were pyrolysis in steam and air atmosphere. The temperature conditions in each atmosphere were essential to remove the residual epoxy on the CF surface. Optimal process conditions such as atmosphere gas type, temperature, and holding time were derived, and the mechanical properties of the recovered r-CF were investigated.

#### **2. Materials and Methods**

For the CFRPs used in this study, polyacrylonitrile-based unidirectional carbon fiber (CF, T700, Toray, Tokyo, Japan) was used as a reinforcing material, and an amine curing agent (4,40 -diaminodiphenyl sulfone, DDS, Sigma-Aldrich, Saint Louis, MO, USA, AHEW of 62.08 g/eq, T<sup>m</sup> of 172–175 ◦C) and a bifunctional epoxy resin (EP, diglycidyl ether of bisphenol A, Kukdo Chem., Seoul, Republic of Korea, YD-128, EEW of 184–190 g/eq, a viscosity of 11,500–13,500 cps at 25 ◦C) were mixed in an equivalent ratio and then impregnated using the hand lay-up method. At this time, the impregnated EP was 35 ± 2 wt.%. Curing was performed at 190 ◦C (10 ◦C/min) for 1 h based on the differential scanning calorimeter (DSC) results shown in Figure 1. The structures of EP and DDS are shown in Figure 2.

°C/min) for 1 h based on the differential scanning calorimeter (DSC) results shown in Fig-

°C/min) for 1 h based on the differential scanning calorimeter (DSC) results shown in Fig-

*Polymers* **2022**, *14*, x FOR PEER REVIEW 3 of 12

ure 1. The structures of EP and DDS are shown in Figure 2.

ure 1. The structures of EP and DDS are shown in Figure 2.

**Figure 1.** Curing temperature of epoxy/4,4′-diaminodiphenyl sulfone mixture using differential scanning calorimetry. **Figure 1.** Curing temperature of epoxy/4,40 -diaminodiphenyl sulfone mixture using differential scanning calorimetry. **Figure 1.** Curing temperature of epoxy/4,4′-diaminodiphenyl sulfone mixture using differential scanning calorimetry.

**Figure 2.** Chemical structures of epoxy and 4,4′-diaminodiphenyl sulfone. **Figure 2.** Chemical structures of epoxy and 4,40 -diaminodiphenyl sulfone.

**Figure 2.** Chemical structures of epoxy and 4,4′-diaminodiphenyl sulfone.

#### *2.1. CFRP Pyrolysis*

*2.1. CFRP Pyrolysis*  A quartz tube furnace (length of 950 mm, inner diameter of 8 mm) was used for the CFRP pyrolysis. The fabricated CFRP was placed on an alumina crucible and loaded into *2.1. CFRP Pyrolysis*  A quartz tube furnace (length of 950 mm, inner diameter of 8 mm) was used for the CFRP pyrolysis. The fabricated CFRP was placed on an alumina crucible and loaded into A quartz tube furnace (length of 950 mm, inner diameter of 8 mm) was used for the CFRP pyrolysis. The fabricated CFRP was placed on an alumina crucible and loaded into a quartz tube. The heating steam was generated at 170 ◦C using triple-distilled water (5 cc/min) in a pre-heater. CFRP was heated to 500–600 ◦C at a rate of 10 ◦C/min and pyrolyzed for 180 min, followed by post-treatment at 500–600 ◦C for 10 min with air (O2- 21% under N2; 200 cc/min). Finally, recycled carbon fiber (r-CF) was recovered after cooling to room temperature. The nomenclature for r-CF according to the pyrolysis conditions is shown in Table 1.


**Table 1.** Nomenclature for recycled carbon fibers and carbon fiber-reinforced plastics pyrolysis conditions.

#### *2.2. Fixed Pyrolysis Conditions*

The curing temperature of the EP/DDS mixture used as a matrix was measured using a differential scanning calorimeter (DSC, DSC 3, Mettler Toledo, Columbus, OH, USA). Approximately 7 mg of the mixture was placed in a 100 µL aluminum pan and heated from 30 ◦C to 300 ◦C at a rate of 10 ◦C/min under an N<sup>2</sup> atmosphere (50 cc/min).

The thermal decomposition temperatures of the CF and EP were measured using a simultaneous thermal analyzer (STA, TGA/DSC 3+, Mettler Toledo). Approximately 4 mg of the samples were placed in a 70 µL aluminum pan and heated from room temperature to 500, 550, and 600 ◦C at a rate of 10 ◦C/min under an N<sup>2</sup> and air atmosphere (50 cc/min).

#### *2.3. Physical Properties*

The tensile properties of a single fiber of Ar-CF and r-CF were measured using a universal testing machine (UTM, Seron, Seoul, Republic of Korea) according to ASTM D3379. At this time, the gauge length of the CF was set to 10 mm, and the draw-off clamp speed was set to 1 mm/min. After placing the CF between the grips of the UTM, both sides of the sample window were cut at the midpoint of the gauge. For each sample, at least 50 CFs were measured; a schematic diagram of the CF tensile strength sampling is shown in Figure 3a. In the tensile test, an error may occur if the carbon fiber is placed diagonally rather than in the center of the window. *Polymers* **2022**, *14*, x FOR PEER REVIEW 5 of 12 diagram for the IFSS sampling of CFs is shown in Figure 3b. The IFSS test can be considered a failure if the carbon fiber is located diagonally to the window or if the epoxy bond is not fully cured and the carbon fiber rather than the EP droplet is pulled out.

**Figure 3.** Schematics of physical properties test; (**a**) tensile strength and (**b**) interfacial shear strength. **Figure 3.** Schematics of physical properties test; (**a**) tensile strength and (**b**) interfacial shear strength.

Functional groups on the Ar-CF and r-CF surfaces were detected in the 400–4000 cm−<sup>1</sup> range using Fourier-transform infrared spectroscopy (FT-IR, Nicolet iS10, Thermo Scien-

The surface properties of Ar-CF and r-CF were observed using field emission scanning electron microscopy (FE-SEM, S-4800, Hitachi, Japan). Ar-CF and r-CF were fixed in the SEM sample holder using carbon tape, and all images were obtained at 104 magnifica-

CF and cured EP/DDS thermal decomposition behaviors were measured by STA under a gas (N2 or air) atmosphere, and the thermal decomposition behavior under a steam atmosphere was measured by a self-made TGA furnace. The results are shown in Figure 4, and a schematic diagram is shown in Figure 5. The mass of CF was decreased by 2.81% (N2), 4.58% (steam), and 7.62% (air) at 600 °C. In addition, EP showed a mass reduction of 71.09% (N2), 85.40% (steam), and 96.64% (air) at 600 °C. The mass loss was the least under N2 atmosphere and the degree of damage to the CF was therefore also small. On the other hand, the degree of damage to CF was large under the air atmosphere because the mass loss was the highest. Under the steam atmosphere, the decrease in mass was between that of N2 and air. Even though the amount of EP char remaining after treatment showed a slight difference (approximately 14.31%) between N2 and steam, it exerted a significant influence during the post-treatment (air) time. The post-treatment can have a significant impact on CF damage. In conclusion, it is considered that the CF damage can be minimized by shortening the duration of air post-treatment because EP char can be quickly

tassium bromide (KBr, Sigma-Aldrich, Saint Louis, MO, USA) together and applying a clamp force of 7 tons for 2 min using a hydraulic press (CrushIR, PIKE Technologies, Mad-

tions at 1.0 × 10−5 torr and an acceleration voltage of 15 kV.

*3.1. Pyrolysis Conditions of EP/DDS and Carbon Fiber* 

*2.4. Component Analysis* 

ison, Wisconsin, WI, USA).

**3. Results and Discussion** 

generated under steam.

The interfacial shear strength (IFSS) between the Ar-CF and r-CF and the EP resin was also measured using UTM. A single fiber was placed at the center of a paper window and adhered with an EP bond, and EP droplets with a diameter of approximately 100–200 µm were placed using a sharp needle. The EP droplet is a difunctional epoxy (EP, diglycidyl ether of bisphenol A, Kukdo Chem., YD-128, EEW of 184–190 g/eq, a viscosity of 11,500–13,500 cps at 25 ◦C) and an amine curing agent (4,40 -diaminodiphenyl methane, Sigma-Aldrich, Saint Louis, MO, USA, AHEW of 49.565 g/eq, Tm of 89 ◦C) was used. It was mixed for 20 min on a hot plate at 80 ◦C, and the temperature was maintained during the droplet operation. EP droplets were cured at 150 ◦C (10 ◦C/min) for 1 h. At least 25 CFs were measured for each sample, and the schematic diagram for the IFSS sampling of CFs is shown in Figure 3b. The IFSS test can be considered a failure if the carbon fiber is located diagonally to the window or if the epoxy bond is not fully cured and the carbon fiber rather than the EP droplet is pulled out.

### *2.4. Component Analysis*

Functional groups on the Ar-CF and r-CF surfaces were detected in the 400–4000 cm−<sup>1</sup> range using Fourier-transform infrared spectroscopy (FT-IR, Nicolet iS10, Thermo Scientific, Waltham, MA, USA). FT-IR samples were prepared as discs by grinding CF and potassium bromide (KBr, Sigma-Aldrich, Saint Louis, MO, USA) together and applying a clamp force of 7 tons for 2 min using a hydraulic press (CrushIR, PIKE Technologies, Madison, Wisconsin, WI, USA).

The surface properties of Ar-CF and r-CF were observed using field emission scanning electron microscopy (FE-SEM, S-4800, Hitachi, Tokyo, Japan). Ar-CF and r-CF were fixed in the SEM sample holder using carbon tape, and all images were obtained at 10<sup>4</sup> magnifications at 1.0 <sup>×</sup> <sup>10</sup>−<sup>5</sup> torr and an acceleration voltage of 15 kV.

#### **3. Results and Discussion**

#### *3.1. Pyrolysis Conditions of EP/DDS and Carbon Fiber*

CF and cured EP/DDS thermal decomposition behaviors were measured by STA under a gas (N<sup>2</sup> or air) atmosphere, and the thermal decomposition behavior under a steam atmosphere was measured by a self-made TGA furnace. The results are shown in Figure 4, and a schematic diagram is shown in Figure 5. The mass of CF was decreased by 2.81% (N2), 4.58% (steam), and 7.62% (air) at 600 ◦C. In addition, EP showed a mass reduction of 71.09% (N2), 85.40% (steam), and 96.64% (air) at 600 ◦C. The mass loss was the least under N<sup>2</sup> atmosphere and the degree of damage to the CF was therefore also small. On the other hand, the degree of damage to CF was large under the air atmosphere because the mass loss was the highest. Under the steam atmosphere, the decrease in mass was between that of N<sup>2</sup> and air. Even though the amount of EP char remaining after treatment showed a slight difference (approximately 14.31%) between N<sup>2</sup> and steam, it exerted a significant influence during the post-treatment (air) time. The post-treatment can have a significant impact on CF damage. In conclusion, it is considered that the CF damage can be minimized by shortening the duration of air post-treatment because EP char can be quickly generated under steam.

#### *3.2. Morphology of the r-CF*

The scanning electron microscopy (SEM) images of the surface morphology of Ar-CF and r-CF are shown in Figure 6. In steam pyrolysis at 500–600 ◦C, the EP of CFRP did not completely decompose. Therefore, some EP remained on the CF surface. However, an increase in the pyrolysis temperature affected the EP decomposition. Then, air pyrolysis induced the complete removal of EP on the CF surface within a short time. However, air pyrolysis at 500 ◦C did not completely remove the EP char on the r-CF and remained in the form of crumbs. On the other hand, air pyrolysis at 550 ◦C completely removed the EP char on the r-CF surface, and the degree of etching on the surface was small. Air pyrolysis at 600 ◦C also completely removed the EP char on the CF surface. However, the

r-CF surface was etched, and a curve was formed due to excessive air treatment. Based on the results above, the conditions of the 50-55-10 sample were effective for EP/DDS-based CFRP recycling.

*Polymers* **2022**, *14*, x FOR PEER REVIEW 6 of 12

**Figure 4.** TGA and DTG graphs of as-received carbon fiber and epoxy resin pyrolysis under gas atmosphere: (**a**) cured epoxy/4,4′-diaminodiphenyl sulfone and (**b**) carbon fiber. **Figure 4.** TGA and DTG graphs of as-received carbon fiber and epoxy resin pyrolysis under gas atmosphere: (**a**) cured epoxy/4,40 -diaminodiphenyl sulfone and (**b**) carbon fiber.

**Figure 5.** Weight loss of as-received carbon fiber and epoxy resin pyrolysis under gas atmosphere: (**a**) N2, (**b**) steam, and (**c**) air. **Figure 5.** Weight loss of as-received carbon fiber and epoxy resin pyrolysis under gas atmosphere: (**a**) N<sup>2</sup> , (**b**) steam, and (**c**) air.

CFRP recycling.

*3.2. Morphology of the r-CF* 

The scanning electron microscopy (SEM) images of the surface morphology of Ar-CF and r-CF are shown in Figure 6. In steam pyrolysis at 500–600 °C, the EP of CFRP did not completely decompose. Therefore, some EP remained on the CF surface. However, an increase in the pyrolysis temperature affected the EP decomposition. Then, air pyrolysis induced the complete removal of EP on the CF surface within a short time. However, air pyrolysis at 500 °C did not completely remove the EP char on the r-CF and remained in the form of crumbs. On the other hand, air pyrolysis at 550 °C completely removed the EP char on the r-CF surface, and the degree of etching on the surface was small. Air pyrolysis at 600 °C also completely removed the EP char on the CF surface. However, the r-CF surface was etched, and a curve was formed due to excessive air treatment. Based on the

**Figure 6.** Scanning electron microscopy images of the as-received and recycled carbon fibers under varying pyrolysis conditions; (**a**) Ar-CF, (**b**) 50, (**c**) 55, (**d**) 60, (**e**) 50-50-10, (**f**) 50-55-10, (**g**) 50-60-10, (**h**) 55-60-10, and (**i**) 60-60-10. **Figure 6.** Scanning electron microscopy images of the as-received and recycled carbon fibers under varying pyrolysis conditions; (**a**) Ar-CF, (**b**) 50, (**c**) 55, (**d**) 60, (**e**) 50-50-10, (**f**) 50-55-10, (**g**) 50-60-10, (**h**) 55-60-10, and (**i**) 60-60-10. ever, the IFSS of r-CF after air treatment at 500–550 °C was calculated as 46.42 MPa (82.76%) and 49.95 MPa (89.05%), respectively, compared to Ar-CF. As steam-air pyrolysis proceeded, EP char debris was formed on the surface of r-CF, and EP droplets did not

#### *3.3. Mechanical Properties of the r-CF* completely adhere to CF but instead formed on the EP char. Meanwhile, the IFSS of the r-

used to calculate the tensile strength.

respectively.

CF and EP.

*3.3. Mechanical Properties of the r-CF*  The tensile strength of the recovered r-CF is shown in Figure 7a. Equation (1) was The tensile strength of the recovered r-CF is shown in Figure 7a. Equation (1) was used to calculate the tensile strength. CF after air treatment at 600° C for 10 min was calculated as 59.14 MPa (105.44%), 59.45 MPa (105.99%), and 61.24 MPa (109.18%), respectively, compared with Ar-CF. As shown in the SEM images in Figure 5, the specific surface area increased due to the curvature of

$$Strength = F/A\tag{1}$$

ℎ = ⁄ (1) where *F* is the force at the break of the fiber, and *A* is the average cross-sectional area of where *F* is the force at the break of the fiber, and *A* is the average cross-sectional area of the fiber. by the increase in functional groups containing oxygen atoms on r-CF by air.

crumbs, which reduced the tensile strength and seemed to be the cause of the large error rate. The tensile strength of the sample subjected to secondary pyrolysis at 550–600 °C was **Figure 7.** Physical properties of the as-received and recycled carbon fibers under varying pyrolysis conditions: (**a**) tensile strength and (**b**) interfacial shear strength. **Figure 7.** Physical properties of the as-received and recycled carbon fibers under varying pyrolysis conditions: (**a**) tensile strength and (**b**) interfacial shear strength.

*3.4. Functional Group of r-CF*  Figure 8 shows the FT-IR analysis that was done to identify the change in functional groups containing oxygen atoms in Ar-CF and r-CF. The hydroxyl (-OH) group on the CF The tensile strength of r-CF was decreased as the pyrolysis temperature increased. In the case of Ar-CF, the tensile strength decreased from 4.67 GPa to 3.03 GPa (60-60-10). As the temperature increased to 500–600 ◦C, the tensile strength of r-CF tended toward 22.48% (4.49 to 3.62 GPa) in the case of the sample that only underwent steam pyrolysis.

was observed at 3431 cm−1 [43,44]. The peaks corresponding to the sp2 C-H bond (=CH)

for the ketone (C=O) and C-O bond were observed at 1638 cm−1 and 1100 cm−1 **[43–45]**,

The FT-IR spectra of r-CF were similar to those of Ar-CF. As the steam-air heat treatment proceeded, the peaks corresponding to the -OH, C=O, and C-O groups increased, and the peaks corresponding to the -CH, =CH, and C=C bonds decreased, suggesting that steam and air heat treatment converted -CH, =CH, and C=C bonds to functional groups containing oxygen atoms. Therefore, it is considered that the introduction of functional groups containing oxygen atoms improved the interfacial bonding interaction between

It is considered that the EP on the CF surface was decomposed by steam to form EP char, and thus the tensile strength was reduced as the CF was oxidized. The tensile strength of the 50-50-10 sample, which was subjected to air post-treatment after steam pyrolysis at 500 ◦C, was reduced by 11.99% (4.11 GPa), and the error rate was calculated to be large compared with that of the other samples. The EP char formed by steam was not completely decomposed in the air atmosphere and remained on the surface in the form of crumbs, which reduced the tensile strength and seemed to be the cause of the large error rate. The tensile strength of the sample subjected to secondary pyrolysis at 550–600 ◦C was reduced by up to 35.12% (3.85–3.03 GPa). Short air pyrolysis at temperatures above 550 ◦C induced a decrease in tensile strength due to oxidation and etching of the r-CF surface. However, the EP char on the CF surface was completely removed.

The IFSS of the recovered r-CF is shown in Figure 7b. Equation (2) was used to calculate the IFSS.

$$\text{IFSS} = \text{F} / \pi \text{DL} \tag{2}$$

where *F* is the equipment force when the fiber breaks, *D* is the droplet size, and *L* is the diameter of the carbon fiber.

The IFSS of Ar-CF was 56.09 MPa. The IFSS of r-CF recovered under steam pyrolysis conditions was lower than the IFSS of Ar-CF, regardless of temperature conditions. However, the IFSS of r-CF after air treatment at 500–550 ◦C was calculated as 46.42 MPa (82.76%) and 49.95 MPa (89.05%), respectively, compared to Ar-CF. As steam-air pyrolysis proceeded, EP char debris was formed on the surface of r-CF, and EP droplets did not completely adhere to CF but instead formed on the EP char. Meanwhile, the IFSS of the r-CF after air treatment at 600◦ C for 10 min was calculated as 59.14 MPa (105.44%), 59.45 MPa (105.99%), and 61.24 MPa (109.18%), respectively, compared with Ar-CF. As shown in the SEM images in Figure 5, the specific surface area increased due to the curvature of the r-CF surface. In addition, it is thought that the improvement in the IFSS was caused by the increase in functional groups containing oxygen atoms on r-CF by air.

#### *3.4. Functional Group of r-CF*

Figure 8 shows the FT-IR analysis that was done to identify the change in functional groups containing oxygen atoms in Ar-CF and r-CF. The hydroxyl (-OH) group on the CF was observed at 3431 cm−<sup>1</sup> [43,44]. The peaks corresponding to the sp<sup>2</sup> C-H bond (=CH) and sp<sup>3</sup> C-H bond (-CH) were observed at 2920 cm−<sup>1</sup> and 2851 cm−<sup>1</sup> [43–45], respectively. In addition, the peak for the C=C bond was observed at 1575 cm−<sup>1</sup> [43–45], and the peaks for the ketone (C=O) and C-O bond were observed at 1638 cm−<sup>1</sup> and 1100 cm−<sup>1</sup> [43–45], respectively.

The FT-IR spectra of r-CF were similar to those of Ar-CF. As the steam-air heat treatment proceeded, the peaks corresponding to the -OH, C=O, and C-O groups increased, and the peaks corresponding to the -CH, =CH, and C=C bonds decreased, suggesting that steam and air heat treatment converted -CH, =CH, and C=C bonds to functional groups containing oxygen atoms. Therefore, it is considered that the introduction of functional groups containing oxygen atoms improved the interfacial bonding interaction between CF and EP.

**Figure 8.** Fourier-transform infrared spectra of the as-received and recycled carbon fibers with varying pyrolysis conditions. **Figure 8.** Fourier-transform infrared spectra of the as-received and recycled carbon fibers with varying pyrolysis conditions.

#### **4. Conclusions 4. Conclusions**

In this study, CFRPs were prepared using epoxy (EP) and 4,4′-diaminodiphenyl sulfone (DDS) as a curing agent. The prepared CFRPs were pyrolyzed under steam and air atmospheres to recover r-CF, and the structural and mechanical properties of the r-CF were evaluated. As the pyrolysis temperature increased, r-CF was oxidized, and the tensile strength was reduced by up to 64.88% relative to that of Ar-CF. As the thermal decomposition temperature increased, r-CF was oxidized and the tensile strength decreased by up to 64.88% compared to Ar-CF, but the IFSS tended to increase by 9.18% due to the increase in oxygen functional groups. In conclusion, a short-duration air post-treatment after steam pyrolysis had the effect of the complete removal of the residual EP char on the CF surface. Based on these results, clean r-CF was recovered from EP/DDS-based CFRP after air post-treatment. Although the tensile strength decreased, the IFSS improved. Therefore, the r-CF with improved IFSS has advantageous for manufacturing high-quality CFRP because of its excellent interfacial bonding with EP. In this study, CFRPs were prepared using epoxy (EP) and 4,40 -diaminodiphenyl sulfone (DDS) as a curing agent. The prepared CFRPs were pyrolyzed under steam and air atmospheres to recover r-CF, and the structural and mechanical properties of the r-CF were evaluated. As the pyrolysis temperature increased, r-CF was oxidized, and the tensile strength was reduced by up to 64.88% relative to that of Ar-CF. As the thermal decomposition temperature increased, r-CF was oxidized and the tensile strength decreased by up to 64.88% compared to Ar-CF, but the IFSS tended to increase by 9.18% due to the increase in oxygen functional groups. In conclusion, a short-duration air post-treatment after steam pyrolysis had the effect of the complete removal of the residual EP char on the CF surface. Based on these results, clean r-CF was recovered from EP/DDS-based CFRP after air post-treatment. Although the tensile strength decreased, the IFSS improved. Therefore, the r-CF with improved IFSS has advantageous for manufacturing high-quality CFRP because of its excellent interfacial bonding with EP.

**Author Contributions:** Conceptualization, B.-J.K.; methodology, B.-J.K.; validation, Y.-M.L.; formal analysis, Y.-M.L.; investigation, Y.-M.L.; writing—original draft preparation, Y.-M.L.; writing—review and editing, K.-W.K.; visualization, Y.-M.L.; supervision, K.-W.K., B.-J.K.; project administration, K.-W.K. All authors have read and agreed to the published version of the manuscript. **Author Contributions:** Conceptualization, B.-J.K.; methodology, B.-J.K.; validation, Y.-M.L.; formal analysis, Y.-M.L.; investigation, Y.-M.L.; writing—original draft preparation, Y.-M.L.; writing—review and editing, K.-W.K.; visualization, Y.-M.L.; supervision, K.-W.K., B.-J.K.; project administration, K.-W.K. All authors have read and agreed to the published version of the manuscript.

**Funding:** This study was supported by the "Industrial Core Technology Development Program" (no. 20012817). The development of functional parts manufacturing technology, which reduced the cost by more than 15% by upcycling carbon fiber and intermediate materials, was funded by the Ministry of Trade, Industry, and Energy, Republic of Korea. This study was supported by the Material Parts Technology Development Project (Project No. 20015733) funded by the Ministry of Trade, Industry Energy (MOTIE), Republic of Korea. **Funding:** This study was supported by the "Industrial Core Technology Development Program" (no. 20012817). The development of functional parts manufacturing technology, which reduced the cost by more than 15% by upcycling carbon fiber and intermediate materials, was funded by the Ministry of Trade, Industry, and Energy, Republic of Korea. This study was supported by the Material Parts Technology Development Project (Project No. 20015733) funded by the Ministry of Trade, Industry Energy (MOTIE), Republic of Korea.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** The data presented in this study are available on request from the corresponding author.

**Conflicts of Interest:** The authors declare no conflict of interest.

## **References**


## *Article* **Fiber Orientation and Strain Rate-Dependent Tensile and Compressive Behavior of Injection Molded Polyamide-6 Reinforced with 20% Short Carbon Fiber**

**Joonhee Lee , Hyungyil Lee and Naksoo Kim \***

Department of Mechanical Engineering, Sogang University, Seoul 04107, Republic of Korea **\*** Correspondence: nskim@sogang.ac.kr; Tel.: +82-010-2703-8635

**Abstract:** As the interest in short-fiber reinforced polymer (SFRP) composites manufactured by injection molding increases, predicting the failure of SFRP structures becomes important. This study aims to systemize the prediction of failure of SFRP through mechanical property evaluation considering the anisotropy and strain rate dependency. To characterize the mechanical properties of polyamide-6 reinforced with carbon fiber of a weight fraction of 20% (PA6-20CF), tensile and compressive experiments were conducted with different load-applying directions and strain rates. Additionally, the results were discussed in detail by SEM image analysis of the fracture faces of the specimen. FE simulations based on the experimental condition were constructed, and the numerical model coefficients were derived through comparison with experimental results. The coefficients obtained were verified by bending tests of the specimens manufactured from composite cross members fabricated by injection molding. Predicting under static and high strain rate conditions, small errors of about 9.6% and 9.3% were shown, respectively. As a result, it proves that explained procedures allow for better failure prediction and for contribution to the systematization of structural design.

**Keywords:** short fiber-reinforced composites; injection molding; thermoplastic material characterization; numerical simulation; complex geometry

### **1. Introduction**

Short fiber reinforced thermoplastic polymers are largely utilized in the industry due to their versatility, recyclability, and low cost and cycle time related to the injection molding process [1]. For these reasons, they have been largely employed in the automotive industry, whose fuel efficiency and emission reduction regulations have been strengthening in recent years. Among the recent regulations, the European one concerning CO2 emission poses a serious challenge to emission reduction (< 95 g/km within 2020). Thus, several automotive manufacturers are researching the direction of replacing aluminum and steel components with reinforced thermoplastic ones because of the considerable weight reduction [2]. In order to properly apply fiber-reinforced thermo-plastic materials to high-performance structural components, the local material properties resulting from the material flow during the thermoplastic injection molding (TIM) process should be properly accounted for. Hence, in the literature, several authors focused their attention on estimating the variation of the mechanical properties of short fiber reinforced (SFR) material, as summarized hereafter.

The effect of the anisotropy developing in injection-molded short glass fiber reinforced polybutylene terephthalate (PBT) polymer was investigated by means of cutting several specimens along 0◦ , 45◦ , and 90◦ , in comparison to the injection direction [3]. One of the most interesting aspects is related to the shell-core structure developing in the material during the injection molding process. This fact results in the fiber in the shell region, close to the surfaces of the specimen, being strongly aligned with the polymer flow direction, whereas, in the core of the specimen, they are transversally aligned with the polymer flow. This aspect was also investigated by means of numerical simulations [4–8]. In addition,

**Citation:** Lee, J.; Lee, H.; Kim, N. Fiber Orientation and Strain Rate-Dependent Tensile and Compressive Behavior of Injection Molded Polyamide-6 Reinforced with 20% Short Carbon Fiber. *Polymers* **2023**, *15*, 738. https://doi.org/ 10.3390/polym15030738

Academic Editors: R.A. Ilyas, S.M. Sapuan, Emin Bayraktar, Shukur Abu Hassan, Nabil Hayeemasae and Khubab Shaker

Received: 1 December 2022 Revised: 25 January 2023 Accepted: 29 January 2023 Published: 31 January 2023

**Copyright:** © 2023 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

much effort has also been focused on investigating the influence of the fiber length on the mechanical properties of SFR thermoplastics. Fu and Lauke presented an analytical model for the estimation of the probabilistic fiber length and orientation distribution in injection molded components and validated it by comparing its prediction to experimental results [9]. In their conclusions, they highlighted how the strength of the composite increases with the fiber length as well as with the reduction of the critical fiber length or, in other words, with the increase of the interfacial adhesion between fibers and matrix polymer. Similar conclusions have been also drawn by other researchers [10,11].

Many studies related to the experimental characterization of the mechanical behavior of SFRP under various loading conditions have also been performed [12–27]. Experiments were conducted by cutting the specimens with various angles from the plates made of glass fiber-reinforced polyamide-6 (PA6) and polyphenylene sulfide (PPS). Changes in mechanical properties according to glass fiber content and fiber orientation angle can be confirmed. Uniaxial tensile and compressive tests for θ = 0◦ ,90◦ specimens were conducted, and the experimental results were verified with the simulation results and application for complex geometry. Similarly, the stress triaxiality-dependent failure behavior of the polyamide-6 with 20% carbon fiber (PA6-20CF) was investigated [22]. In addition, many studies have been conducted to investigate the mechanical behavior and failure mechanism of polymer composites reinforced with glass fiber and carbon fiber under compressive load [28–32]. The progress of the kink-band and the fracture mode were especially described in detail in terms of fractography [28,29]. Additionally, there were experimental results that confirmed that the mechanical behavior of fiber-reinforced composites can be different depending on the strain rate [6,24,28,29,33,34]. The deviation in stress–strain response was influenced by the interfacial properties because of the strain rate sensitivity of the fiber and matrix. In addition, multi-scale modeling and simulation for characterizing the material mechanical behavior have been utilized [8,35–37]. As mentioned earlier, in the cases of composites, since consideration of anisotropy greatly affects the accurate prediction of mechanical properties, attempts to numerically model the fiber orientation through experiments and simulation have been made steadily [38–43].

For the mechanical properties of polyamide-6 reinforced with carbon fiber studied in this paper, Karsli and Aytac investigated the thermal and morphological properties [44]. They focused on the influences of fiber length and fiber content. Additionally, the failures of fibers and matrix were analyzed through SEM image analysis of fracture surfaces of the specimens. In addition, Zhang investigated the strain rate and temperature-dependent mechanical properties of carbon fiber-reinforced thermoplastic composites fabricated by hot press molding [45]. This research presented that the modulus and strength of continuous carbon fiber reinforced thermoplastic (CCFRT) increase with a higher strain rate. Additionally, the relationship is non-linear, and it can be confirmed that it is due to the increase in the strength of the thermoplastic matrix and the fiber interface. In this study, changes in mechanical properties of PA6-20CF manufactured through injection molding according to fiber orientation distribution and strain rate are identified. Additionally, by appropriately considering local tensile-compressive mechanical properties caused by injection molding, contributed to the systematization of the structural design of complex structures. The research is articulated according to the following goals. First, experimental characterization of PA6-20CF through specimens of different thicknesses and angles fabricated by injection molding. Second, verification by injection molding simulation. Last, consideration of local mechanical properties and verification of tensile-compressive mechanical properties through the mapping process. Therefore, this paper presents descriptions of the material and the geometry of the specimen used in the experiment to obtain tensile-compressive mechanical properties according to the fiber orientation. Then, the experimental results are discussed and compared in detail with the analysis of the failure mechanism observed in SEM images. Based on the anisotropic elastoplastic mechanical behavior of the material, the numerical model is explained, and simulations were conducted. Considering the strain ratedependent tensile-compressive mechanical properties, the numerical model coefficients

according to the fiber orientation were derived by calibration of the load–displacement results of the experiment and simulation. As a final step, mechanical properties are input based on the obtained coefficients, and a bending test under static and high strain rate conditions was conducted which was machined from a cross member manufactured by injection molding of PA6-20CF (each jig speed: 1 mm/min, 20 mm/min). It allowed evaluation of how much the previously described mechanism can improve the failure prediction accuracy in the analysis of complex geometries. Under the static and high strain rate conditions, failure was predicted with small errors of about 9.6% and 9.3%, respectively. Therefore, the validity of the derived sets of properties was confirmed, suggesting that these research procedures can improve the reliability of the structural analysis of composite materials with complex shapes under complex stress states and contribute to design systematization.

#### **2. Materials and Methods**

#### *2.1. PA6-20CF Plane Specimen*

In this study, PA6-20CF material reinforced with KN111 polyamide-6 (Kolon Plastics Inc., Gimcheon, Republic of Korea) and Torayca T700S (Toray Carbon Fibers America Inc., Tacoma, USA.) short carbon fiber with a weight fraction of 20% was investigated. PA6-20CF plate specimens with a thickness of *t* = 1.8 mm and 4.0 mm were manufactured through the injection molding of the PA6-20CF pellet, and the specifications are shown in Figure 1a, b. Before the injection process of the plate specimen, the pellet was dried at 120 ◦C for 6 h to prevent deterioration of properties due to hygroscopicity and was injected under the following process conditions: the mold surface temperature of 85 ◦C, melt temperature of 285 ◦C, injection time with 2 sec, and cooling time with 20 sec. After that, specimens were manufactured within a small processing error range of 0.25 mm through wire cutting of the plates. Tensile and compression specimens were prepared for three orientation angles *θ* = 0◦ , 45◦ , and 90◦ , which are defined as the angle between flow direction with the direction of applied load during tests (see Figure 1b,c).

**Figure 1.** (**a**) PA6-20CF plate manufactured by injection molding, and injection direction, (**b**) Schematic illustration of different fiber orientation (*θ*) specimens obtained from the injectionmolded plate (X = loading axis), (**c**) fiber alignment in specimen with fiber orientation (*θ*) = 0◦ , 45◦ , 90◦ .

In addition to that, the machining locations of the specimen from the plate were decided through finite volume method (FVM) analysis, and a detailed description will be presented in Section 2.2. After wire cutting, all the machined specimens were stored in a sealed plastic bag with silica gel to prevent moisture absorption. To prevent deviations in physical properties due to the hygroscopicity of all specimens, prior to the test, the HKOD-50 electric furnace (heat flow: 12 m3/h) was used for drying at 120 ◦C for 6 h, and then the temperature was lowered for 1 h at room temperature. After that, the tests were conducted. The tensile tests were carried out according to ASTM-D638-02a-Type IV standard and the machined positions on the plate are the same at the thickness t = 1.8 and 4.0 mm of the plate as shown in Figure 2a [46]. An Epsilon 3542 extensometer with a gauge

length of 25 mm was used. Additionally, the tests were performed with different jig speeds of 1.5, 15, and 150 mm/min for considering strain rate dependency (strain rates of 0.001, 0.01, and 0.1 s−<sup>1</sup> , respectively).

**Figure 2.** (**a**) ASTM-D 638-02a-Type IV specimen dimensions and cutting location of the tensile specimens from the plate, (**b**) ASTM-D3410M-16 specimen dimensions and cutting location of the compressive specimens from the plate.

The compressive tests were carried out according to ASTM D3341M-16 standard and the machined positions on the plate are the same at the thickness *t* = 1.8 and 4.0 mm of the plate as shown in Figure 2b [47]. The tests were performed with different jig speeds of 0.6, 6, and 60 mm/min for considering strain rate dependency (strain rates of 0.001, 0.01, and 0.1 s−<sup>1</sup> , respectively). To prevent buckling that occurs in the specimen when a compressive load is applied, the specimen has a relatively short gauge length compared to the grip part. Additionally, it was confirmed that only acceptable failure characteristic results within the ASTM D3341M-16 standard were derived. For repeatability verification, *θ* = 0◦ , 90◦ : 5 times, and *θ* = 45◦ : 3 times or more were performed according to each test condition (fiber orientation and strain rate) both tensile and compressive.

#### *2.2. Numerical Implementation*

In the research, three different commercial software were utilized. Autodesk Moldflow Insight (AMI) 2019 (San Rafael, CA, USA) was employed for the estimation of the fiber orientation distributions in the injection-molded plates and specimens. The Autodesk Helius 2019 (San Rafael, CA, USA) allowed for the mapping of the local mechanical properties, as a result of the fiber length and orientation distributions. The mapping was carried out from the results thermoplastic injection molding (TIM) process simulation into the input file of the ABAQUS/Standard structural simulation. Finally, ABAQUS 2017 (Dassault Systemes, Velizy-Villacoublay, France) was employed for the estimation of local stress and strain distribution on specimens.

#### 2.2.1. Thermoplastic Injection Molding Simulation

The thermoplastic injection molding simulation was developed by means of the commercial software Autodesk Moldflow Insight 2019. To properly locate the specimens on the plate used for machining the tensile and compressive specimens, a finite volume method simulation of injection molding of the plate was conducted as shown in Figure 3. By this action, cross-sections at the center of the specimens could have a consistent fiber orientation distribution.

**Figure 3.** Results of the fiber orientation distribution from cut faces of the specimens from the plate of thickness *t* = (**a**) 1.8 mm, (**b**) 4.0 mm.

In the simulation, the same process parameters utilized in the real injection molding process have been utilized, as summarized previously. The fiber orientation distribution was derived through the ARD-RSC model presented in Equation (1) [48,49]. The ARD-RSC model considered the deformation rate tensor ( . *γ*) as a polynomial rather than a constant in the original Folgar–Tucker formulation [50]. Thus, the overestimation of the rate of change of the fourth-order fiber orientation tensor was adjusted. As a result, it was improved to better describe the interactions between fibers with a length of 1 mm or longer.

*∂*aij *<sup>∂</sup>*<sup>t</sup> <sup>=</sup> <sup>−</sup><sup>1</sup> 2 <sup>w</sup>ikakj <sup>−</sup> <sup>a</sup>ikwkj + <sup>1</sup> <sup>2</sup>*λ* . *γ*ikakj + aik . *γ*kj − 2 . *<sup>γ</sup>*kl<sup>h</sup> aijkl + (1 − k) <sup>L</sup>ijkl <sup>−</sup> <sup>M</sup>ijkl · <sup>a</sup>mnkli + . *γ* h 2 <sup>c</sup>ij <sup>−</sup> (<sup>1</sup> <sup>−</sup> <sup>k</sup>)cklMijkl − 2k · ckkaij − 5 <sup>c</sup>ikakj <sup>+</sup> <sup>a</sup>ikckj <sup>+</sup> 10ckj aijkl + (1 − k) <sup>L</sup>ijkl <sup>−</sup> <sup>M</sup>ijkl · <sup>a</sup>mnkli (1)

In Equation (1), the constant *k* is a dynamic scalar factor describing the ability to realign the fibers along the mold after injection. A low *k* means that the orientation ability of the fiber itself is low and, therefore, the degree of agreement between the fiber orientation distribution and the polymer resin flow is low. Conversely, when the value of *k* is large, the fibers are realigned to a high degree of agreement with the direction of resin flow in the mold between injections. In this research for the PA6-20CF, the orientation dynamic scalar factor *k* = 0.5 of ARD-RSC was utilized, which was utilized and verified through the comparison of fiber orientation distribution confirmed by SEM analysis and simulation in previous studies [22,23]. In the TIM simulation, 486,195 and 1,285,994 tetrahedral elements were utilized with *t* = 1.8 mm plates and for *t* = 4.0 mm plates, respectively. The mechanical properties of the PA6-20CF material were inputted in the Moldflow numerical simulation and the constants for the Ramberg–Osgood flow stress model, Equation (2), were calculated and are reported in the next section.

$$
\sigma = E^{1/n}(\mathsf{K})^{(n-1)/n} \left( \varepsilon\_{p, \varepsilon f f} \right)^{1/n} \tag{2}
$$

This procedure allowed mapping the mechanical properties on the injection molding simulation mesh during the filling phase, sparing time during the mapping operation. Along with the flow stress model of Equation (2), a modified Hill '48 yield function in Equation (3) was utilized. It allowed considering the variation of the mechanical properties, for each element of the mesh, caused by the local fiber length and orientation distributions.

$$\sigma\_{eff} = \left[\frac{1}{2}\left\{\left(a\sigma\_{11} - \beta\sigma\_{22}\right)^2 + \left(\beta\sigma\_{22} - \beta\sigma\_{33}\right)^2 + \left(\beta\sigma\_{33} - a\sigma\_{11}\right)^2\right.\tag{3}$$

$$+ 6\left[\left(\sigma\_{12}\right)^2 + \left(\sigma\_{23}\right)^2 + \left(\sigma\_{31}\right)^2\right]\right\}\right]^{1/2} \tag{3}$$

The *α* and *β* parameters in Equation (3) account for the improvement of the polymer matrix mechanical properties due to the presence of the carbon fibers, along the direction parallel and transversal to the local flow direction, respectively, and were calculated by Equation (4).

$$\begin{aligned} \alpha(\lambda\_I) &= \theta + \left[ \frac{(a\_m - \theta)}{(\lambda\_{m,l} - 1/2)} \right] (\lambda\_I - 1/2) \\ \beta(\lambda\_I) &= \theta + \left[ \frac{(\beta\_m - \theta)}{(\lambda\_{m,l} - 1/2)} \right] (\lambda\_I - 1/2) \end{aligned} \tag{4}$$

For each element of the mesh, *α* and *β* are calculated based on *λ*<sup>I</sup> , *α*m, *β*<sup>m</sup> and *λ*m, I constants. The *λ*<sup>I</sup> defines the first eigenvalue of the fiber orientation tensor of the considered element, and the *λ*m, I define the first eigenvalue of the fiber orientation tensor in strong fiber alignment along the flow direction region.

#### 2.2.2. Mapping Procedure

The mapping procedure was carried out in the Autodesk Helius software and allowed us to map the local mechanical properties, as a result of the TIM process, into the ABAQUS numerical simulation. This procedure is an important task in determining how the original fiber orientation distribution is reflected in the specimen when the specimen was machined at each location on the plate. After the mapping operation, the ABAQUS input file containing a specific set of mechanical properties, for each element of the mesh, was created and ought to be run from the Autodesk Helius command shell.

The Ramberg–Osgood model constants are based on isotropic behavior, and the respective σ-ε curves for tension and compression were considered. Once the stress state is determined, the program inputs the mechanical properties suitable for the stress state of each element. At each iteration step, both tensile and compressive solutions are calculated using a set of tensile and compressive mechanical properties. Then, the stress state is determined according to two verification procedures as shown in Figure 4, and a corresponding appropriate solution is used for each step.

**Figure 4.** Schematic diagram of the procedure how to define the stress state.

#### 2.2.3. Structural Simulation

The structural numerical simulations were constructed into ABAQUS/Standard utilizing the C3D10—quadratic hexahedral elements. Both tension and compression specimen geometries were modeled identically to the specifications of the specimen used in the test as previously shown in Figure 2. The region connected to the jig was given a rigid body condition, and in the case of the tensile specimen, displacement at the gauge distance node was used as output data. For *t* = 1.8 mm specimens, 9216, 22,400 elements were used for tensile and compressive specimens. Additionally, for *t* = 4.0 mm specimens, 18,432, 40,800 elements were used, respectively. The results of the ABAQUS numerical simulations are reported in detail in Sections 3.3 and 3.4 of the paper.

#### **3. Results and Discussions**

#### *3.1. Experimental Results*

In this section, the results of tensile and compressive tests performed at room temperature for specimen thicknesses *t* = 1.8, 4.0 mm, fiber orientation angles *θ* = 0◦ , 45◦ , 90◦ , and

strain rates 0.001, 0.01, and 0.1 s−<sup>1</sup> are presented. Furthermore, the mechanical properties and the behavior of the anisotropic elastoplastic material derived for each test condition are discussed. Afterward, it was analyzed based on the fractographic study of the SEM image, which is discussed in detail in Section 3.2.

#### 3.1.1. Tensile Tests

Engineering stress-engineering strain curves obtained as a result of the tensile tests are shown in Figures 5 and 6. Additionally, the anisotropic elastoplastic material behavior (with hardening) of short fiber reinforced polymer (SFRP) according to the strain rate can be observed. As the fiber orientation angle *θ* decreased and strain rate increased at *t* = 1.8 mm and 4.0 mm, ultimate tensile strength *σ*UTS, elastic modulus, and yield strength tended to increase. On the other hand, ultimate tensile strain *ε*UTS tended to decrease as the fiber orientation angle *θ* decreased and the strain rate increased. Based on this, it shows that the material became brittle as the strain rate increased. When the thickness *t* of the plate is different (*t* = 1.8 and 4.0 mm), different material behavior is shown because of the deviation in the degree of anisotropy. For *θ* = 0◦ , the strength is greater at *t* = 1.8 mm, but for *θ* = 45◦ and 90◦ , the strength is greater at *t* = 4.0 mm. This is because the thinner the plate, the higher the shell ratio in the shell-core structure (higher average fiber orientation index), and the better the fibers are aligned in the main flow direction (Figure 3). Additionally, for the same reason, when *t* = 1.8 mm, it shows big differences in mechanical properties between *θ* = 0◦ and *θ* = 45◦ , 90◦ , but small at *t* = 4.0 mm.

**Figure 5.** Tensile engineering stress–engineering strain curves of *t* = 1.8 mm specimens with different strain rates. (**a**) *θ* = 0◦ , (**b**) *θ* = 45◦ , (**c**) *θ* = 90◦ .

**Figure 6.** Tensile engineering stress–engineering strain curves of *t* = 4.0 mm specimens with different strain rates. (**a**) *θ* = 0◦ , (**b**) *θ* = 45◦ , (**c**) *θ* = 90◦ .

#### 3.1.2. Compressive Tests

Engineering stress-engineering strain curves obtained as a result of the compressive tests are shown below in Figures 7 and 8. Similar to the tensile test results, the behavior of the anisotropic elastoplastic material of the short fiber reinforced polymer according to the strain rate could be observed, and different results were obtained when the thickness *t* was changed. At *t* = 1.8 mm, the strain rate sensitivity of *θ* = 0◦ is greater than that of *θ* = 45◦ , 90◦ . At both *t* = 1.8 mm and 4.0 mm, the ultimate compressive strength *σ*UCS, elastic modulus, and yield strength tended to increase as the strain rate increased. On the other hand, the ultimate compressive strain *ε*UCS tended to decrease as the strain rate increased.

**Figure 7.** Compressive engineering stress–engineering strain curves of *t* = 1.8 mm specimens with different strain rates. (**a**) *θ* = 0◦ , (**b**) *θ* = 45◦ , (**c**) *θ* = 90◦ .

**Figure 8.** Compressive engineering stress–engineering strain curves of *t* = 4.0 mm specimens with different strain rates. (**a**) *θ* = 0◦ , (**b**) *θ* = 45◦ , (**c**) *θ* = 90◦ .

Comparing the results of the compressive test with the tensile test, it can be observed that the short fiber-reinforced polymer exhibited completely different mechanical behavior under tensile and compressive loading conditions. First, under the compressive load condition at the same strain rate, it showed that the *σ*UCS at *θ* = 45◦ was the lowest and the *ε*UCS was the biggest. In addition, it can be observed that the difference in *σ*UTS between *t* = 1.8 mm and *t* = 4.0 mm was relatively smaller than the difference in σUCS. Furthermore, in compression, when *t* = 4.0 mm, it showed that the strength was higher than that of *t* = 1.8 mm, at all *θ*. From this, it can be concluded that the mechanical properties of the short-fiber reinforced polymer are relatively strongly influenced by the matrix in the compressive state [31,32,45].

#### *3.2. SEM Image Analysis*

In this section, the experimental results explained in Section 3.1 are analyzed and verified through SEM (Scanning Electron Microscope) image analysis of the fracture surface of the specimen. For SEM analysis, a JSM-7100F FE-SEM machine (JEOL Ltd., Tokyo, Japan) with a resolution of 3.0 mm (at 15.0 kV) was utilized. The specimen fracture faces were coated with Pt, and SEM imaging was taken perpendicular to the fracture surface. The fiber orientation at the fracture surface and failure characteristics of the matrix and fiber were analyzed.

#### 3.2.1. Tensile Specimens

Through SEM analysis of the tensile specimen, it can be confirmed that the alignments of the fibers observed at the fracture surface were different depending on the fiber orientation *θ* (see Figure 9). It can be observed that the fibers of the *θ* = 0◦ specimen are cut perpendicular to the fracture plane, and in the case of the *θ* = 45◦ and 90◦ specimens, the fibers are aligned parallel to the fracture plane. In addition to that, the length of the fibers observed at the fracture surface of the *θ* = 90◦ specimen is the longest, and the length of *θ* = 0◦ is the shortest [3,21].

**Figure 9.** Fracture faces observed from tensile specimens which were tested with a strain rate of 0.001 s−<sup>1</sup> (magnification: ×35, scale: 100 µm).

In the case of *θ* = 0◦ , since the fracture plane is perpendicular to the fiber direction, it can be expected that the fiber fracture occurs predominantly. Various microscopic damage mechanisms according to the fiber orientation *θ* under tensile load can be observed. It is possible to observe fiber failure of shortened fibers after breakage and fiber pull-outs in which fibers were pulled out from the matrix in Figure 10a. Additionally, as shown in Figure 10b, debonding occurred as traces in which fiber separated from the matrix [51]. The fracture face of the fiber in Figure 10c is consistent with the fracture characteristics of the carbon fiber under the tensile load studied previously [52–54].

**Figure 10.** Fracture mechanisms observed for *θ* = 0◦ tensile specimen: (**a**) Fiber failure and fiber pullouts (magnification: ×200, scale: 100 µm), (**b**) fiber-matrix debonding (magnification: ×1000, scale: 10 µm), (**c**) fracture face of the fiber after the tensile failure (magnification: ×10,000, scale: 1 µm).

At the failure surface of the *θ* = 90◦ specimen, completely different failure characteristics were observed, as shown in Figure 11. The average length of the fibers was long. Additionally, the fiber tracks left as traces after the fibers were separated from the matrix and the hackles distributed along them can be observed (Figure 11a,b). Moreover, the matrix that was bonded to the fibers was weakened and a smeared matrix appeared. Through the deformation of the matrix, the ductile behavior of *θ* = 90◦ can be explained. In the case of *θ* = 45◦ , the behavior is intermediate between *θ* = 0◦ and *θ* = 90◦ (Figure 12). Although the average fiber length observed was long, a large number of fibers in which breakage has occurred can also be found. Therefore, it can be confirmed that the failure characteristics

that can be seen at *θ* = 0◦ and 90◦ coexist. The fracture face of the fiber was the roughest at *θ* = 0◦ , and conversely, the smoothest at *θ* = 90◦ .

**Figure 11.** Fracture mechanisms observed for *θ* = 90◦ tensile specimen: (**a**) Fiber tracks, hackles, and smeared matrix (magnification: ×200, scale: 100 µm), (**b**) fiber tracks and broken long fiber (magnification: ×500, scale: 10 µm), (**c**) ductile failure of the matrix (magnification: ×1000, scale: 10 µm), (**d**) smooth fracture face of the fiber (magnification: ×10,000, scale: 1 µm).

**Figure 12.** Fracture mechanism observed for *θ* = 45◦ tensile specimen, (**a**) fiber pull-outs and fiber tracks (magnification: ×200, scale: 100 µm), (**b**) smeared matrix and fiber pull-outs (magnification: ×1000, scale: 10 µm), (**c**) fracture surface of the fiber (magnification: ×10,000, scale: 1 µm).

#### 3.2.2. Compressive Specimens

In compressive specimen SEM images in Figure 13, the direction of fibers aligned at each fiber orientation *θ* show similar results with tensile (see Figure 9). However, the average length of fibers observed at the fracture surface did not show a significant difference between *θ* = 45◦ and *θ* = 90◦ . Moreover, the deformed matrix appears brighter at *θ* = 45◦ than at *θ* = 90◦ . Through this, more deformation of the matrix can be grasped at *θ* = 45◦ .

**Figure 13.** Fracture surfaces observed for compressive specimens with a strain rate of 0.001 s−<sup>1</sup> (magnification: ×40, scale: 100 µm).

For *θ* = 0◦ compressive specimen, the failure characteristics were similar to those of the tensile specimen, but in compression, fiber breakage was promoted by the buckling of the fiber, and fibrillation induced by strong shear ductile damage occurred, as shown in Figure 14a,b. Additionally, fiber-matrix debonding can be observed in Figure 14b,c [51]. The fracture surface of the fiber in Figure 15c is comparable to the fracture characteristics of the carbon fiber under flexural load studied previously [52,53].

**Figure 14.** Fracture mechanism observed for *θ* = 0◦ compressive specimen, (**a**) fiber failure and fibrillation (magnification: ×300, scale: 10 µm), (**b**) fiber-matrix debonding (magnification: ×500, scale: 10 µm), (**c**) fracture face of the fiber after flexural failure (magnification: ×5000, scale: 1 µm).

**Figure 15.** Fracture mechanisms observed for *θ* = 90◦ compressive specimen: (**a**) fiber failure, smeared matrix, and fiber tracks (magnification: ×300, scale: 10 µm), (**b**) fiber tracks and fibrillation (magnification: ×400, scale: 10 µm), (**c**) fracture face of the fiber after flexural failure (magnification: ×10,000, scale: 1 µm).

Both in *θ* = 90◦ and 45◦ , fiber tracks, the smeared matrix near the fiber, and fibrillation occurred frequently as shown in Figures 15 and 16. In compression, fibers commonly showed flexural failure due to buckling, and ductile shear damage was dominant [28–32,52–54]. Fiber failure was observed more frequently at *θ* = 90◦ than at *θ* = 45◦ , conversely, there was more deformation on the matrix in the case of *θ* = 45◦ . These facts support that the σUCS, 90◦ , is bigger than σUCS, 45◦ .

**Figure 16.** Fracture mechanism observed for *θ* = 45◦ compressive specimen: (**a**) smeared matrix and fiber tracks (magnification: ×200, scale: 100 µm), (**b**) smeared matrix and fiber tracks (magnification: ×350, scale: 10 µm), (**c**) fracture surface of fiber due to flexural failure (magnification: ×10,000, scale: 1 µm).

#### *3.3. Numerical Model Verification*

To prove the reliability of the mapping procedure and the anisotropic elastoplastic behavior of PA6-20CF, EXP-FE load–displacement curves for tensile and compression specimens were compared. The same Ramberg–Osgood coefficients were applied to both *t* = 1.8 and 4.0 mm at the same load condition and strain rate. Additionally, the comparison result of the load–displacement curves of the EXP-FE under each test condition (load condition and strain rate) is shown in Figures 17 and 18. The input coefficient sets were as shown in Table 1. In both tension and compression, as the strain rate increased, *K*, *n* decreased, *α*, *β*, and the elastic modulus of matrix and fiber tended to increase. Moreover, in the case of compression, the change in the elastic modulus was more significant. As mentioned earlier, different fiber orientation distribution causes a difference in mechanical properties, and it is important to consider this properly. By the mapping procedure, local mechanical properties as a result of the injection molding process can be input for each element in FE simulation, so the reliability of the results can be improved. The possibility to consider the thickness-dependent (fiber orientation-dependent) mechanical properties with one set of model coefficients is a huge advantage for the simulation of complex geometries. The coefficients derived from the EXP-FE comparison were finally verified by being used for structural analysis of structures with complex shapes under a complex stress state in the next section—Section 3.4.

**Figure 17.** Comparison between experimental and numerical results for ASTM-D638 tensile specimens.

**Figure 18.** Comparison between experimental and numerical results for ASTM-D3410 compressive specimens.

**Table 1.** Model constants for the Ramberg–Osgood flow stress model and for the modified Hill'48 yield function.


#### *3.4. Application: Composite Cross Member Rear Part Bending Specimen*

In order to verify the effect of considering the tensile-compressive mechanical behavior of PA6-20CF according to the fiber orientation and load conditions, a bending test of a specimen machined from PA6-20CF composite cross member was conducted (see Figures 19 and 20). The cross member was manufactured by injection molding under the following process conditions (Table 2), and TIM analysis was performed by inputting the same process parameters in the simulation. In addition to that, the model coefficients set shown in Table 1 was input to evaluate the reliability of the derived PA6-20CF mechanical properties. By the mapping procedure, the local mechanical properties due to the injection molding were considered in the specimen machined by cutting the cross member. In order to secure the reliability of the mechanical properties at the static and high strain rates derived previously, the jig speed was varied, and EXP-FE comparison was performed for two cases (jig speed: 1 mm/min, 20 mm/min). The experiment setup is the same as shown in Figure 20, and for the bending experiments, a universal test-one machine was used.

**Figure 19.** (**a**) PA6-20CF composite cross-member with aluminum insert fabricated by injection molding (bottom view), (**b**) fiber orientation distribution result of the cross member and detailed view at the machined location.

**Figure 20.** (**a**) Dimensions and configuration of the specimen after cutting, (**b**) FE model and boundary conditions.

**Table 2.** Injection molding process parameters of PA6-20CF composite cross member.


In Figure 21, the bending test load–displacement curves for two different jig speeds were compared. The simulation results in which only tensile properties were input and results in which both tensile and compressive properties were entered were compared with the experimental results. In the case of the jig speed with 1 mm/min, as seen in Figure 21a, when only the tensile properties were considered, the result was 3.1 kN higher than the experimental result (relative error: 27.2%), whereas when tensile and compressive properties were simultaneously considered, the result was 1.1 kN lower (relative error: 9.6%). In addition to that, the result of the jig speed with 20 mm/min is shown in Figure 21b. When only the tensile properties were input, there was an error of 5.5 kN (relative error: 42.6%), and when both were considered, an error of 1.2 kN was obtained (relative error: 9.3%). By comparing the load at the failure displacement, it can be seen that the error was improved by about 26% at the jig speed of 1 mm/min and by about 13% at 20 mm/min. In the case of PA6-20CF, the tensile and compressive mechanical behaviors are different due to carbon fibers, and this reinforces the argument that these should be properly considered. Additionally, comparing Figure 21a,b, it can be confirmed that the prediction error was significantly reduced when the strain rate was properly considered. As a result, these show that investigated procedures can contribute to predicting the failure of the PA6-20CF structures.

**Figure 21.** Bending test results comparison between EXP and FE with jig speed: (**a**) 1 mm/min, (**b**) 20 mm/min.

### **4. Conclusions**

The main objectives of this paper were as follows: first, failure mechanism investigation through experimental characterization and SEM verification of injection molded PA6- 20CF; second, defining the advantage of simultaneous consideration of tensile-compression properties of anisotropic elastoplastic material behavior in the complex stress state; third, design systemization of SFRP which is fabricated by injection molding considering mechanical properties according to fiber orientation distribution and strain rate.

By experiments and SEM image analysis, mechanical behavior of PA6-20CF with various load conditions and fiber orientation was identified. Additionally, it was found that the failure mechanism was very different depending on the fiber orientation and the loading conditions. Afterwards, utilizing the mapping procedure, inputting local mechanical properties induced by the injection molding process into the structural numerical simulation was carried out. By this method, it was possible to accurately consider the anisotropic mechanical behavior of the uniaxial tensile and compressive specimen, and by the comparison of EXP-FE results, the Ramberg–Osgood flow stress model coefficients were derived. A structural analysis model was constructed in which the mechanical properties at each load condition and strain rate were inputted and, to verify it, a bending test was conducted with a specimen manufactured by machining a composite crossmember. It demonstrated high EXP-FE correlation under static conditions and high strain rate conditions to prove the validity and applicability of modeling.

**Author Contributions:** Conceptualization, J.L.; methodology, J.L.; software, J.L.; validation, J.L.; formal analysis, J.L.; investigation, J.L.; resources, N.K.; writing—original draft preparation, J.L.; writing—review and editing, N.K.; visualization, J.L.; supervision, N.K.; project administration, H.L., N.K.; All authors have read and agreed to the published version of the manuscript.

**Funding:** This research received no external funding.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** The data presented in this study are available on request from the corresponding author.

**Acknowledgments:** This research was supported by the Carbon Industrial Cluster Development Program (1008 3609) funded by the Ministry of Trade, Industry & Energy (MOTIE, Republic of Korea) and by the Material Component Technology Development Program (2000 4983) funded by the Ministry of Trade, Industry & Energy (MOTIE, Republic of Korea).

**Conflicts of Interest:** The authors declare no conflict of interest.

### **References**


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## *Article* **Thermal Effects on Mechanical Strength of Additive Manufactured CFRP Composites at Stable and Cyclic Temperature**

**Isyna Izzal Muna 1,\* , Magdalena Mieloszyk <sup>1</sup> , Ruta Rimasauskiene <sup>2</sup> , Nabeel Maqsood <sup>2</sup> and Marius Rimasauskas <sup>2</sup>**


**Abstract:** Additive manufacturing (AM) techniques can be applied to produce carbon-fiber-reinforced polymer (CFRP) elements. Such elements can be exposed to different environmental factors, e.g., temperature, moisture, and UV radiation, related to their operational conditions. From a variety of environmental factors, the temperature is one of the most typical. Temperature strongly influences matrix material joining together CFRP components, resulting in material strength reduction. Therefore, it is important to understand processes in the composite material caused by temperature. This experimental work investigated the thermal effects on the performances of AM CFRP composites. Specimens with unidirectional (UD) alignments of the fiber reinforcement were printed using the fused deposition modeling (FDM) technique. The printed specimens were subjected to two different thermal conditions: stable continuous at 65 ◦C and cyclic temperature between 50 and 70 ◦C. Tensile testing was performed to study the mechanical strength and Young's modulus of AM UD-CFRPs. In order to investigate the morphological structure on the surface of AM specimens, an optical microscope, scanning electron microscope (SEM), and digital microscope were utilized. Untreated (intact) samples attained the highest average tensile strength value of 226.14 MPa and Young's modulus of 28.65 GPa. The ultimate tensile strength of the sample group subjected to stable heat treatment decreased to 217.99 MPa, while the thermal cycling group reduced to 204.41 MPa. The Young's modulus of the sample group subjected to stable thermal exposure was decreased to 25.39 GPa, while for the thermal cycling group, it was reduced to 20.75 GPa. The visual investigations revealed that the intact or untreated specimen group exhibited lateral damage in top failure mode (LAT), the thermally stable group underwent edge delamination in the middle (DGM) as the nominated failure mode, and the explosive breakage at gauge in the middle (XGM) failure mode occurred in the sample from the thermal cycling group. Based on morphological observations at the microscale, the delamination, fiber pull-out, and matrix cracking were the dominant damages in the 3D-printed tensile-tested specimens. The molecular chains of the polymer changed their structure into an amorphous one, and only local motions of stretching occurred when the specimens were exposed to stable heating (prolonged). In the case of thermal cycling, the strain gradients were accumulated in the matrix material, and the local stresses increased as a result of the reheating and re-cooling exposure of the polymeric composites; the molecular motion of the long-range polymer structure was reactivated several times. Micro-cracking occurred as a result of internal stresses, which led to material failure and a reduction of the mechanical properties.

**Keywords:** thermal cycle; thermally stable; 3D-printing technology; fused deposition modeling; CFRP composites; mechanical degradation; morphological structure

**Citation:** Muna, I.I.; Mieloszyk, M.; Rimasauskiene, R.; Maqsood, N.; Rimasauskas, M. Thermal Effects on Mechanical Strength of Additive Manufactured CFRP Composites at Stable and Cyclic Temperature. *Polymers* **2022**, *14*, 4680. https:// doi.org/10.3390/polym14214680

Academic Editor: R.A. Ilyas

Received: 29 September 2022 Accepted: 31 October 2022 Published: 2 November 2022

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2022 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

#### **1. Introduction**

The technological advancement in additive manufacturing (AM) over the last decade has been a major driver for the development of composite materials. It is widely known that AM technologies, specifically the fused deposition modeling (FDM) method, have numerous benefits owing to the low price, flexibility, and simplicity with small amounts of waste material, which indeed open up endless possibilities in their utilization in engineering and industrial fields [1–3]. Carbon-fiber-reinforced polymer (CFRP) has potential applications in the industrial field such as automotive bumper brackets [4] and the development of more durable wind turbines [5]. The printed composite materials must have a long service life in real applications. However, these materials are more likely exposed to harsh environmental conditions specifically related to temperature. Depending on its different magnitude and period time, the exposure to temperature of 3D-printed composites could affect the mechanical behavior of the materials differently. Researchers usually use different thermal cycling profiles to study the behavior of aerospace vehicles such as satellite components, which depend on their case study or thermal cycling test apparatus [6].

There have been a large number of papers that experimentally studied the effect of thermal treatment of the mechanical behavior of polymeric composites manufactured with the FDM method at a continuous temperature [7–12]. Several results reported that the heat treatment leads to the better mechanical properties of 3D-polymer-based composites [10–12]. Handweker et al. investigated the influence of a heat treatment (annealing) process on the mechanical performance of continuous- and chopped-fiberreinforced polyamide 6 in the build-up direction. It was found that Young's modulus increased by a factor of three, while the ultimate tensile strength (UTS) increased by 50% for the chopped-carbon-fiber-reinforced material and 186% for the continuous-glass-fiberreinforced material [11]. In Wang et al. [12], heat treatment was performed on three printed continuous-carbon-fiber-reinforced polymer (CCFRP) composite specimens with different layer distributions and heating conditions. The dimensional change of CCFRCs during heat treatment was shown to be much more closely related to the microstructure change than the material crystallization. Furthermore, heat treatment could improve the mechanical properties of CCFRCs by decreasing the porosity while also strengthening the interface. Heat treatment, specifically, delayed the initiation of cracks in CCFRCs during bending tests by changing the failure mode of the matrix layers and improving the bonding between fiber and bundles/matrix. Another experimental work by Nassar et al. revealed that heat treatment was able to enhance the bonding of filament layers and reduce the porosity content of CFRP composite printed using the FDM method, which resulted in an improvement of its tensile modulus. However, there was a limited improvement in the tensile strength and modulus of elasticity values for the samples treated at low heat treatment temperature [10]. Despite the benefits of using FDM to generate durable materials with high mechanical properties, this method is also known to have some limitations, such as nozzle clogging [13,14] and the appearance of satellite droplets [15].

However, some other experiments on CFRP composites showed a decreased mechanical behavior after thermal treatment [7,16]. Zhang et al. investigated the effect of high temperature on pure epoxy resin and 3D carbon/epoxy braided composites [7]. It was revealed that a constant decline in compressive properties occurred as a result of the combination of matrix degradation and fiber/resin interface cracking. At temperatures lower than the glass transition temperature (Tg) of the pure epoxy resin, the matrix aging was the main aging process, although fiber/matrix interface debonding can be detected at temperatures greater than the Tg, such as 180 ◦C. After thermal aging, their micro-morphologies and compressive behaviors were reported. Jia et al. experimentally investigated the influence of heat treatment performed at a stable temperature on CFRP composites generated conventionally, ranging from −100 ◦C to 100 ◦C [16]. The results showed that, under static and dynamic three-point bending tests, CFRP composites had improved flexural strength, maximum deflection, and energy absorption at lower temperatures (−60 ◦C, −100 ◦C), but performed poorly at higher temperatures (100 ◦C). At various temperatures, experimental

photos from post-mortem photographs, scanning electron microscopy, and high-speed films were used to investigate various failure behaviors such as micro-buckling, kinking, and fiber breakage. In more recent experimental research by [6], 3D-printed CF/PEEK was exposed to thermal cycling and then evaluated using tensile and arc heating tests. It was reported that the thermal cycle resulted in decreased tensile strength and the length of the samples increased after the heating test.

Until recently, the number of experiments of thermal treatment conducted in another thermal mode, e.g. cyclic temperature, remains very limited since most heat treatment experiments on 3D-printed composites were mostly performed at a stable continuous temperature. However, some papers investigated the effect of thermal cycling at low temperatures on composites produced conventionally. Ghasemi et al. studied the residual stress during thermal cycling on a graphite/epoxy composite [17]. The cycle was performed one time and used as a change in temperature from 70 ◦C to 100 ◦C and back to 70 ◦C. This profile has a heating/cooling (transition) rate of 3–5 ◦C per minute, while the duration of the dwelling time at the temperature extremes was about 15 min. It was reported that the mechanical properties of this composite were degraded when the material was exposed to thermal cycling conditions. Up to a certain period of exposure, the composite may retain its strength and stiffness above its allowable limits. However, as time passes, the strength and stiffness may become very low.

Another thermal cycling treatment of a carbon fiber/epoxy composite generated conventionally was carried out by Gonzalez et al. [18]. On CFRP samples, 500 thermal cycles at 15 ◦C/min and a 5 min dwell time were applied during the experimental investigation. Two different atmospheres were used in the thermal cycling test. The first represented an inert atmosphere (nitrogen gas) for isolating any oxidative agents' effects. For the second atmosphere, dry air was used to simulate material oxidation effects. At high temperatures, matrix oxidation was expected, but matrix cracking occurred due to a mismatch in the thermal expansion coefficient. The combination of both effects resulted in CFRP aging acceleration. Lafarie et al. studied the characterization of the damage processes of carbon/epoxy laminates under isothermal aging and thermal cycling conditions in a vacuum and oxidative atmosphere [19]. The thermal cycling experiments were 500 thermal cycles with maximum and minimum temperatures of 180 and −50 ◦C, respectively, and constant cooling and heating rates of 4 ◦C/min. It was found that, during a thermal cycling test in an oxidative atmosphere such as air, there was a coupling effect between matrix oxidation, which occurs at the highest temperatures of the cycle, and matrix cracking caused by thermo-mechanical ply stresses caused by prohibited differential expansions of the plies. Previously, heating treatment of unidirectional (UD) continuous CFRP composite at an elevated temperature was experimentally and numerically investigated by Muna et al. [20]. The coupled thermal–displacement analysis was used during the numerical calculations. The strain in the sample due to its exposure to elevated temperature was measured using fiber Bragg grating (FBG) sensors.

Thermal degradation of additive manufactured polymer composite has not been explored by researchers comprehensively, yet it is an important phenomenon that will be useful in the structural health prediction of lightweight materials. Up to the present, the majority of research studies on the thermal degradation of 3D-printed composites have only been investigated on specimens at a stable temperature for a certain duration, while for the thermal cycling treatment, most of the research work has been performed on the specimens produced not with 3D-printing technology, but rather, conventionally. Furthermore, few explorations are available on the work performed on 3D-printed polymeric composites under both thermal conditions at continuous and cyclic temperatures and the effect on the mechanical degradation. This paper aims to investigate the effect of thermal exposure on the mechanical and morphological behavior of 3D continuous-carbon-fiber-reinforced polymers at stable continuous and cyclic temperatures.

#### **2. Experimental Work**

#### *2.1. Additive Manufacturing of CFRP Composite Samples*

Additive manufacturing (AM) is a well-known processing technique for adding materials to manufacture parts based on a 3D computer model in a layer-by-layer manner, which can also be referred to as 3D printing or rapid prototyping (RP). In this experimental work, the FDM method was chosen due to its simplicity and ability to print complex geometries with low cost and flexibility to use different materials. Unidirectional (UD) continuous CFRP composites were printed as experimental specimens. The benefit of employing continuous carbon fibers (CCFs) as a reinforcement agent in polymer-based composites is that their weight remains light with the increased strength. Thermoplastic polylactic acid (PLA) was chosen as the matrix agent.

The CFRP composite specimens were additively manufactured with the fused deposition modeling (FDM) method at Kaunas University of Technology. A modified 3D printer MeCreator2 with dual inputs at the extrusion printing head, which was designed and developed by Rimasauskas et al., was employed [14]. The pre-impregnation process was applied for the CCF reinforcement for better printing quality and adhesion. For the impregnation process, PLA pellets were dissolved in dichloromethane solution from Eurochemicals with the ratio concentration of 90 g/10 g using the magnetic LBX H01 mini-stirrer at 600 rpm. Later, the standard non-impregnated CCF was passed through this resin and simultaneously dried with a hot air gun temperature of 220 ◦C. After the impregnation process, the CCF was further printed with the PLA matrix filament to form a composite specimen. The continuous carbon fiber used was T300D-3000, consisting of 3000 fibers in an untwisted tow (Toray, France). The materials' properties are shown in Table 1. The information of all materials concerning the mechanical properties was provided by the material suppliers. The 3 K CCF was used as the reinforcement material and was kept the same throughout the experimental work. An extrusion multiplier was used to control the content, but it was also kept the same at 0.5 for all the manufactured composite specimens. The prepared materials were then ready to be used for additive manufacturing process with the FDM method. The 3D model of the specimens was created in a CAD package, which was then exported as a stereo lithography (STL) file and subsequently loaded into 3D printer slicing software (Simplify3D, Cincinnati, OH, USA) to generate the G-code, which the 3D printer used to print out each test specimen. The CAD files of the specimens were used to extrude and deposit molten thermoplastic, which was built up in layers from a horizontal base. Two types of FDM printers were used during the experiments, a modified one with two inputs at the extruder and another 3D printer with one input from Prusa company (for printing the gripping tabs). The printing parameters are shown in Table 2.

**Table 1.** Material properties of continuous carbon fiber and polylactic acid.


**Table 2.** Printing parameters.


The extrusion printing was updated in order to include two inputs (one for the polymer filament and the other for the carbon fiber tow) and one output that allowed the polymer to be fused with the CF for printing. As was discussed previously in the experimental research by [14], the working principle of PLA and CF mixing consisted of several steps. Firstly, PLA filament was fed through a PTFE tube by using a standard filament feeding system. When the filament reached the printing head, which was heated up to 200 ◦C, the polymer liquefied. Subsequently, the impregnated carbon fiber tow was inserted through the PTFE tube to the printing head, directly to the printing nozzle. The molten polymer came into contact with the carbon fiber in the mixing zone, and secondary carbon fiber impregnation was performed. The molten polymer was constantly pushed together with the carbon fiber through the nozzle. It should be emphasized that the carbon fiber was fed solely by the molten polymer and did not require any additional feeding equipment. Later, the carbon fiber coated with polymer was extruded through a printing nozzle onto borosilicate glass mounted on an aluminum plate.

Generally, the manufacturing process of CFRP composites with the FDM method requires a long period of time (e.g., in this experimental case, about 1 h for printing 1 specimen), and the 3D printer is not designed to handle errors. Printing mistakes are frequently expected, and continuing the printing process when an error appears may endanger the 3D printer, and the final part of the printed sample typically cannot be used. One common printing error is nozzle clogging [13,14]. In order to minimize this detrimental issue, numerous efforts have been made by researchers to overcome nozzle clogging during the printing process [13,21,22]. Gutierrez et al. proposed a method to decrease the clogging deposition rate of alumina inclusions in continuous casting nozzles through three simultaneous measures: flow modification, the use of raw materials with a low impurity content, and smoothed internal surfaces [13]. Internal flow regulation entails avoiding dead zones and generating symmetric patterns. The feasibility of these methods was tested using a mathematical model. Using this approach, the adherence of inclusions to the nozzle wall applied a boundary condition based on the thickness of the sub-laminar boundary, rather than the typical "trap" boundary condition. The generic boundary condition produces deposition rates that are unaffected by the inclusion size. The proposed boundary condition uses particle sizes to distinguish against the clogging deposition rate. Plant studies with these nozzles, together with water modeling showed that the current strategy may significantly reduce clogging occurrence. More recently, as a response to the problem of nozzle clogging, Sampedro et al. created a real-time process monitoring system capable of accurately forecasting abnormal behaviors in the FDM printing method [21]. A network of collaborative sensors was utilized to collect time series data, which were then processed by the suggested machine learning method. The multi-head encoder–decoder temporal convolutional network (MH-ED-TCN) gathers features from the data, assesses their impact on the many processes that occur throughout an operational printing cycle, and distinguishes between normal and problematic manufacturing operations. The studies conducted yielded a 97.2% accuracy in predicting the future behavior of a 3D printer. More recent experimental work has been conducted to improve the quality of the 3D-printed composite specimen using a controlled cooling effect [23]. In regard to achieving this printing quality, the effect of printing parameters (extrusion temperature, extrusion width, and extrusion multiplier) on the tensile and flexural properties were examined. The highest Young's and flexural modulus of 31.50 GPa and 15.68 GPa, respectively, were reached when the composite structure was manufactured using printing parameters E.W. −1.3 mm and E.M. −0.5.

Next, thermoplastic PLA reinforced with CF composite was extruded through the printing nozzle, which was printed on borosilicate glass mounted on the aluminum plate bed surface at a certain temperature according to the optimized printing parameters in the literature (i.e., 90 ◦C). Before starting the printing process, adhesive spray should be spread out on top of the borosilicate glass to avoid the dragging of printed filaments. Once the printing of the sample completed, a cutter was used to separate the filament

spool and printed part. Subsequently, the borosilicate glass should be removed from the aluminum plate bed and to allowed cool down, then the 3D-printed sample can be removed from the glass by the aid of blades. The dimension of unidirectional (UD) CFRP composites was 150 × 13 × 2 mm with 6 layers of the same fiber alignment in all layers. The volume fraction of the composite was kept constant throughout the experimental work due to having the same printing parameters. The carbon fiber reinforcement content was measuring using the tool length path, and the estimated carbon fiber content was 18.2% (wt.). The total number of specimens printed was 15 samples in which 5 samples for each group of thermal treatment were required in accordance with the D3039 ASTM standard used for tensile testing. The specimens are depicted in Figure 1. After the manufacturing stage completed, the weight and size dimensions of the composite parts were measured using a digital scale and a caliper, respectively. The weight and dimension measurement was necessary in order to compare the changes in size and mass before and after thermal loading. Furthermore, it was important for the calculation of the mechanical properties after the mechanical testing.

**Figure 1.** The 15 CFRP composite specimens printed with the FDM method for (**a**) intact/no treatment, (**b**) continuous temperature, and (**c**) cyclic temperature.

#### *2.2. Thermal Treatment on Printed Specimens*

In an attempt to investigate the effect of temperature exposures on the printed CFRP composites via FDM, the specimens were thermally treated under two different thermal conditions: stable continuous and cyclic at low temperature around the glass transition (Tg) of the PLA polymers used in this work. An air-circulated environmental oven "Memmert" Model UFB-400 was utilized for stable and cyclic heating, which supplied a continuous renewal of oxygen in a certain percentage of ambient air, as presented in Figure 2. For the stable continuous heat treatment, the specimens were placed inside the pre-heated oven at 65 ◦C and kept for 6 h, then naturally cooled at room temperature. As for the thermal cycling, the oven temperature was manually adjusted from 50 to 70 ◦C according to the cyclic plan (see Figure 3) for 6 cycles. One cycle consisted of two extreme temperatures with a dwell time of 10 min, a heating rate of 1 ◦C/min, and a cooling rate of 2.5 ◦C/min. These temperatures were chosen according to the glass transition temperature of the PLA used in this work, which lies around 65 ◦C. It was expected that the specimens would undergo cross-linking and a crystallization change of the polymer up to the state where its structural shape would particularly deform.

Dwell time means how long of an exposure period is appropriate at each temperature limit. If a prolonged dwell time is employed, the duration of the test will lengthen until the number of cycles decreases. Then, again, in some instances, a long dwell period can potentially speed up the testing process if it generates structural changes that increase the stresses during temperature fluctuations. At high temperatures, polymer materials, for instance, relax or creep—the polymer chains within the material move to alleviate the stresses imposed by the high temperature. When the temperature is decreased, these alterations may result in a large increase in the stresses imposed on buildings. However, sufficient dwell time is necessary for these modifications to take place.

**Figure 2.** An air-circulated oven for hot temperature treatment.

**Figure 3.** Thermal cycling profile for one cycle. A = dwell time; B = heating rate; C = cooling rate.

#### *2.3. Static Tensile Testing*

The tensile testing was performed after thermal treatment to examine the mechanical strength and Young's modulus of the treated specimens, as well as the untreated samples (intact). A Tilnius Olsen H25KT (capacity 25 kN) universal testing machine was utilized to perform tensile testing with the standard head displacement speed of 2 mm/min and at room temperature. The tensile testing setup is shown in Figure 4. The ASTM standard D3039 was used for the in-plane tensile testing, which requires each tested specimen to have four tabs in gripping position at the top and bottom. Prior to the tensile testing, PLA tabs having dimensions 50 × 12.5 × 2 mm were printed separately, and four points were marked 15 mm from the center of the specimens to measure the elastic strain. The printed tabs then were adhered to the specimens with a universal structural bonder (adhesive glue) by applying pressure via clamps to hold and secure the parts tightly. In this study, the ASTM D3039 standard was used to perform the tensile testing of the specimens, in which five samples were required for each group of thermal treatment to determine the average material properties.

#### *2.4. Morphological Investigation*

The micro-morphological analyses on the surface were performed before and after performing the cyclic and stable thermal loadings by utilizing an optical microscope device (Nikon Eclipse LV100ND) equipped with a high-definition color camera (Nikon DS-Ri2). The imaging software (NIS Elements 4.5.1.00, Nikon Europe B.V., Amstelveen, The Netherlands) was used to prepare and process the data at 5× magnification. The

maximum sample size observed with the optical microscope was 150 × 150 mm. A scanning electron microscope device (FE-SEM SU5000, Hitachi Co., Tokyo, Japan) was employed to investigate the micro-structure damage of different specimen groups after tensile testing. The maximum specimen size observed with the SEM was 200 mm in diameter and 80 mm in height. A digital microscope (Levenhuk, DTX 500 LCD, Warsaw, Poland) was utilized to capture the macroscopic images of the tensile-fractured samples. The maximum specimen size observed with the digital microscope was about 150 × 100 mm. The microscopes are presented in Figure 5.

**Figure 4.** Experimental setup of tensile testing.

**Figure 5.** (**a**) Optical microscope; (**b**) scanning electron microscope; (**c**) digital microscope.

#### **3. Results**

#### *3.1. Dimensional Change*

The samples were measured before and after the heat treatment in order to determine the amount of geometrical changes, as shown in Table 3. Any thermal treatment is expected to have an effect on the dimensions and shape of a plastic part. For the test samples, the length, as well as the height and width were measured using a digital vernier caliper dimensional measurement tool. Table 1 shows the change of the specimens' length, width, and thickness before heat treatment (after 3D printing) and after heat treatment. The difference between the nominal design from the CAD software and the printed dimension in practice is referred as variations of the relative values of the length, width, and thickness. These values were then acquired and averaged. In comparison to the nominal model, the

dimensions of printed CFRP samples in the xy-plane (length and width) were undersized around (0.62 ± 0.93)% and (0.07 ± 0.27)%, respectively. However, in comparison with the 3D model, the dimension in the z direction (thickness) was oversized, and a mean error of 13.44 ± 0.05% was obtained. Overall, these findings are consistent with previous research on the dimensional correctness of 3D-printed composite parts [8,24]. The noticeable difference in the dimension of the specimens can be hypothetically attributed to the difficulty of depositing impregnated continuous carbon fiber filament. During the printing process, the clogging of the printer nozzle could possibly occur due to the improper proportion of polymer matrix filament mounted into the impregnated CCF material. In addition to that, the issue with nozzle clogging might lead to the breakage of the CCF filament throughout the deposition process. The material deposition issue was encountered by the authors in this experimental work. In accordance with previous experimental research, a similar problem of nozzle clogging occurred when the amount of carbon fiber filament was higher than the polymer matrix, which caused a greater tension at the nozzle tip and resulted in damaged carbon fiber filament [14]. It is worth noting that carbon fiber filament was printed according to the pre-defined path where the carbon content was calculated by evaluating the path traveled by the printing head and carbon weight per unit length. Therefore, in an attempt to tackle this issue during the materials' deposition, the cooling fan shall be adjusted to prevent either a lack or excessive polymer extruded onto the printing bed.

**Table 3.** Dimensions of 3D-printed specimens before and after heating treatment.


#### *3.2. Static Tensile Results*

The tensile properties of the 3D-printed specimens of each thermal group (untreated, continuous, and cyclic) are shown in Figure 6. From the bar graph plot, it can be seen that the untreated (intact) samples attained the highest average tensile strength value of 226.14 MPa and Young's modulus of 28.65 GPa. The ultimate tensile strength of the sample group subjected to stable heat treatment decreased to 217.99 MPa, while for the thermal cycling group, it reduced to 204.41 MPa. The Young's modulus of the sample group subjected to stable thermal exposure decreased to be 25.39 GPa, while for the thermal cycling group, it reduced to 20.75 GPa. The mean values in the bar plots were used to depict the trends of each printed group specimen and the range of their tensile characteristics effects. The degradation in the mechanical strength and elastic modulus after thermally stable and cyclic loading was attributed to the difference in the coefficient of thermal expansion (CTE) between the matrix and fiber, which was caused by the reduced crosslinking of the polymers [25]. This CTE discrepancy resulted in thermal stress, and it can cause the fibers to pull out due to fiber–matrix debonding, which then leads to the mechanical deterioration of composite specimens [18,26]. ANOVA analysis was performed to identify statistical differences and significance between the unheated group and the stable continuous and cyclic temperature on the tensile strength with a significance level of 5%. A one-way ANOVA revealed that there was a statistically significant difference in tensile strength between the three tested groups with a *p*-value of 0.0108 and test statistic F of 6.764 (F(2, 27) = [6.764], *p* = 0.0108).

**Figure 6.** Experimental results of the tensile properties of 3D-printed samples. (**a**) Tensile strength; (**b**) Young's modulus.

Figure 7 depicts the representative stress–strain curves of unidirectional CFRP composites during the static tensile tests at various thermal conditions. All of the stress–strain curves showed a linear elastic regime, followed by a stress drop. From the tensile stress–strain curves, it can be observed that the untreated/intact composite specimen reached the maximum stress level, followed by the stable continuous and cyclic composite specimens. The stable continuous 3D-fabricated composite specimen was slightly less compared to the untreated specimen with almost similar behavior in strain, while the cyclic composite specimen attained the lowest stress level among the composite group and reached beyond the strain level of an untreated and stable continuous composite specimen, which indicated its greater elasticity.

**Figure 7.** Average stress–strain curve for the tensile samples.

The visual investigation of the failure area of the 15 specimens was studied in accordance with ASTM standard D3039. The failure mode of each thermal group is presented in Figure 8. It was indicated that the intact or untreated specimen group exhibited the lateral-at tab-top (LAT) failure mode and well agreed with the result reported by researchers [27]. The intact sample after longitudinal tension loading exhibited lateral fiber breakage and fiber splitting completely at the top region near the tab, which is gripped by the gauge. Edge delamination in the middle (DGM) was the nominated failure mode in the stable continuous heated group's specimen. It can be seen that there was debonding between layers in the middle due to the lower strength of the interfacial adhesion formed after the continuous thermal loading for 6 h at 65 ◦C. The explosive breakage at gauge in the middle (XGM) failure mode occurred in the sample from the treated group after thermal exposure with cycling mode between 50 ◦C and 70 ◦C for six cycles, and this result was similar to what was studied by Ghasemi et al. [28]. In this damage mode, a large number of fibers pull out near the gauge, and this was presumed due to reduce the chain

scission and cross-linking of the polymer matrix caused by thermal cycling. Moreover, it can be observed from Figure 8c that layer debonding also occurred after thermal cycling.

**Figure 8.** Fractured samples after performing tensile testing. (**a**) Intact; (**b**) continuous; (**c**) cyclic.

#### *3.3. Morphological Investigation*

Morphological analyses were performed before and after performing the thermal loading exposures by utilizing an optical microscope device, while a scanning electron microscope was employed to investigate the micro-structure of the different specimen groups after tensile testing. The microscopes are shown in Figure 5.

In order to represent the difference on the structural surface, one specimen of each heating group before and after thermal loading was compared. It can be seen from Figure 9 that there was a slight change in the morphological surface before and after the thermal performance in this experiment. The polymer PLA as a matrix material had a slightly smoother and finer appearance on the continuous and cyclic treatment samples. However, these specimens did not show some deformed shapes (some wrinkle shapes along the longitudinal direction) since it did not undergo sufficiently high-temperature exposure.

**Figure 9.** Optical micrographs of specimen before and after stable temperature at 65 ◦C and cyclic temperature between 50 and 70 ◦C.

From the SEM micrographs shown in Figure 10, it can be observed that the effect of thermal cycling and continuous heating on the 3D-printed CFRP specimens caused different microstructural fracture damages after tensile testing in both the polymers and carbon fibers compared to the untreated (intact) specimens. The PLA polymers from the thermally treated specimens were split up with a relatively micro and massive cracked hole located in the middle for continuous and cyclic heating, respectively. As for the intact sample, its polymer structure remained morphologically unseparated, and fracture did not exist. The presented micro-crack in the matrix material was one of the main damage mechanisms caused by the thermal cycles.

**Figure 10.** Scanning electron microscope images of the specimen after destructive tensile testing: (**a**) intact (untreated); (**b**) stable continuous at 65 ◦C; (**c**) cyclic between 50 and 70 ◦C in a 60◦ , 0◦ , and 90◦ view angle, respectively.

#### **4. Discussion**

Based on the experimental results of the tensile test, the untreated group displayed superior strength results when compared to the heated group subjected to stable continuous and cyclic temperatures. The mechanical response after stable temperature exposure showed a moderate loss in tensile strength and Young's modulus. The behavior trends of Young's modulus and the tensile strength were almost similar with a higher detrimental effect under thermal cycling. The mechanical strength of 3D CFRP specimens subjected to the continuous temperature and the cyclic temperature was reduced by 3.6% and 9.6%, respectively. The Young's modulus of the specimens subjected to the continuous temperature and the cyclic temperature was decreased by 11.4% and 26.5%, respectively. These values obtained from the tensile testing indicated the degradation of the mechanical properties after the thermal exposures.

It was decided to investigate the fracture interface of the 3D-printed samples in order to monitor the deformation behavior and how the fracture happened prior to and after mechanical testing. To study such occurrences, an optical microscope was employed to investigate the interfacial adhesion between the continuous carbon fiber (CCF) and PLA thermoplastic matrix prior to destructive tensile testing. The SEM microscope was used to examine the fracture interface of the 3D CFRP specimens and changes in the matrix microstructure from each group. To investigate such outcomes, one specimen from each group was chosen that best reflected the failure mode.

The thermal cycling performed in this work was with an oxidative condition, where environmental air was used during the cooling process by opening the ventilator to the maximum scale to allow the air to enter the oven chamber. This oxidative atmosphere led to accelerated aging of the CFRP composites due to a mismatch in the thermal expansion coefficient (CTE), which caused the polymer matrix's oxidation and cracking [18,19]. While microscopic-scale damage such as micro-cracking occurred as a result of the impact and internal stresses, micro-cracking was the leading cause of material failure due to its undetectable nature, as well as the induced structure fragmentation, which led to a reduction in the mechanical properties such as strength, stiffness, and dimensional stability [29].

The visual investigation of the failure area of the 15 specimens was studied in accordance with ASTM standard D3039. The failure mode of each thermal group is presented in Figure 8. It was indicated that the intact or untreated specimen group exhibited the lateral-at tab-top (LAT) failure mode and well-agreed with the result reported by researchers [27]. The intact sample after longitudinal tension loading exhibited lateral fiber breakage and fiber splitting completely at the top region near the tab, which was gripped by the gauge. Edge delamination in the middle (DGM) was the nominated failure mode in the stable continuous heated group's specimen. It can be seen that there was debonding between layers in the middle due to the lower strength of interfacial adhesion formed after the continuous thermal loading for 6 h at 65 ◦C. The explosive breakage at gauge in the middle (XGM) failure mode occurred in the sample from the treated group after thermal exposure with cycling mode between 50 ◦C and 70 ◦C for six cycles, and this result was similar to what was studied by Ghasemi et al. [28]. In this damage mode, a large amount of fibers pulled out near the gauge, and it was presumed to be due to the reduced chain scission and cross-linking of the polymer matrix caused by the thermal cycling. Moreover, it can be observed from Figure 8c that layer debonding also occurred after the thermal cycling.

The appearance of breakage phenomena can be also triggered by the interfacial forces exerted between the fiber and the matrix, as presented in Figure 8. It has been recently reported in research work where functionalized surfaces with pure cellulose nanocrystals [30] are faced with composites made by matrices with a polar nature and cellulose-rich hemp/flax fibers [31]. These studies concluded that the breakage of bundles and the single hemp/flax fibers was due to these high interfacial forces. In addition, the early stages of cycling were also linked to the production of microvoids, followed by interfacial sliding, which caused composite quality degradation [32].

The mechanical properties of polymeric composites are more matrix-dependent, making them more susceptible to thermal aging. However, it is worth noting that the type of thermal loading can also influence composite behavior. As mentioned previously, the detrimental effect on the mechanical behavior of the 3D CFRP composite was attributed to the difference of the coefficient of thermal expansion (CTE) between the matrix and fiber, and it was caused by the reduced cross-linking of the polymers [25]. This CTE disparity caused local thermal stress, which might cause fiber pull-out due to fiber–matrix debonding. The debonding of fiber and matrix occurs in an oxidative environment due to the significant strain gradients present in matrix locations near highly stiff fibers, causing high local stresses, which lead to crack initiation and mechanical deterioration of composite specimens [18,19,26].

The visual inspection of the morphological surface shown in Figure 9 showed that there was a slight change to the morphological surface before and after the thermal treatment in this experiment. On stable (prolonged) and cyclic treatment samples, the polymer PLA as the matrix material had a slightly smoother and finer appearance. However, because they were not exposed to sufficiently high temperatures, these specimens did not exhibit visible warpage or deformed shapes (some wrinkle shapes along the longitudinal direction). Moreover, the warpage of the 3D CFRP specimen was governed by the stress level released in the macroscopic dimension [33]. It can be observed from Figure 10 that matrix cracks occurred after the thermally stable treatment, and in the cycled group, the crack was larger. When the specimens were exposed to stable heating (prolonged) for 6 h at 65 ◦C, which

is a few degrees above the glass transition temperature of the PLA polymer used, the molecular chains of the polymer changed their structure into an amorphous one, and only local motions of stretching occurred [20]. In addition to that, the prolonged heating allowed the formation of adsorptive inter-layer bonds and volatile groups. As a result, the polymer structure changed, and inner stress relaxation occurred as a result of this thermoprocess, which is known as thermo-relaxation [34]. In the case of thermal cycling, the strain gradients were accumulated in the matrix material, and the local stresses increased as a result of the reheating and re-cooling exposure of the polymeric composites; the molecular motion of the long-range polymer structure was reactivated several times.

#### **5. Conclusions**

In summary, we investigated the effect of stable and cyclic temperature on the mechanical properties of unidirectional 3D CFRP composites. Static tensile testing was performed to obtain the mechanical strength, Young's modulus, and failure damage of these composites. The results indicated that unidirectional CFRP composites undergo thermal exposures, revealing a degraded mechanical strength and Young's modulus under stable temperature at 65 ◦C and cyclic temperature between 50 and 70 ◦C. The specimens under thermal cyclic possessed lower mechanical performance compared to the untreated group by having a 9.6% reduction in tensile strength and 26.5% in the elastic modulus. Meanwhile, the 3D CFRP composite group under continuous thermal exposure exhibited a lower decrease of the mechanical properties by having a 2.4% reduction in tensile strength and 11.4% in the elastic modulus. The degradation in the mechanical strength and elastic modulus after thermally stable and cyclic loading was attributed to the difference in the coefficient of thermal expansion (CTE) between the matrix and fiber, which was caused by the reduced cross-linking of the polymers.

Based on morphological investigations using an optical microscope, the failure mode of 15 specimens was studied in accordance with ASTM standard D3039. It was indicated that the intact or untreated specimen group exhibited lateral damage at top failure mode (LAT), the thermally stable group underwent edge delamination in the middle (DGM) as the nominated failure mode, and the explosive breakage at gauge in the middle (XGM) failure mode occurred in the sample from the thermal cycling group.

Based on the SEM observations, the delamination, fiber pull-out, and matrix cracking were the dominant damages in the 3D-printed tensile-tested specimens. The molecular chains of the polymer changed their structure into an amorphous one, and only local motions of stretching occurred when the specimens were exposed to stable heating (prolonged). In the case of thermal cycling, the strain gradients were accumulated in the matrix material, and the local stresses increased as a result of the reheating and re-cooling exposure of the polymeric composites; the molecular motion of the long-range polymer structure was reactivated several times. Micro-cracking in the matrix occurred as a result of internal stresses, which led to material failure and a reduction in the mechanical properties. The temperature had a considerable influence on the matrix material, and connecting together the CFRP components reduced the material's strength. As a result, it is critical to comprehend temperature-induced reactions in composite materials.

The findings presented in this work offer pathways to investigate further the degradation of 3D composites under several temperature conditions to study the integrity of the structure and the damage mechanism. In future work, similar thermal exposure modes (thermally stable and thermal cycle) will be performed at different magnitudes to provide an extensive horizon of thermal effects on the mechanical characteristics of 3D-printed CFRP composites. Furthermore, investigating the effect of different thermal cycling parameters such as heating and cooling rates and dwell time on the materials' properties would be beneficial for industrial applications. The thermal degradation of the polymer material used in the additive manufacturing of a composite should be considered to be investigated through rheological and thermogravimetric measurements to understand the thermal behavior of the material during the heating conditions.

Ultimately, understanding the thermal influence at different conditions on the morphological structure of FDM printed materials and their mechanical properties can show the first step to inspecting and eliminating them. These thermal effects may incite researchers to develop techniques of thermal treatment of printed specimens that withstand their mechanical characteristics to be potential applications for the operation of space under various thermal conditions.

**Author Contributions:** Conceptualization, I.I.M. and M.M.; methodology, I.I.M., M.M., R.R. and M.R.; software, I.I.M. and M.M.; validation, M.M.; formal analysis, I.I.M. and M.M.; investigation, I.I.M., M.M., R.R., N.M. and M.R.; resources, M.M., R.R. and M.R.; writing—original draft preparation, I.I.M., M.M. and N.M.; visualization, I.I.M., M.M. and N.M.; supervision, M.M. and R.R.; project administration, M.M.; funding acquisition, M.M. All authors have read and agreed to the published version of the manuscript.

**Funding:** The research was financed by National Science Centre, Poland, within the project entitled Thermal degradation processes of additive manufactured structures (2019/35/O/ST8/00757), and the Polish National Agency for Academic Exchange within the Foreign doctoral fellowship program (PPN/STA/2021/1/00006).

**Institutional Review Board Statement:** Not applicable.

**Data Availability Statement:** Not applicable.

**Acknowledgments:** The authors of this work would like to thank to researchers from the Kaunas University of Technology, Lithuania, for the sample manufacturing.

**Conflicts of Interest:** The authors declare that they have no competing financial interest that could have influenced the work reported in this study.

### **References**


## *Article* **Ballistic Limit of UHMWPE Composite Armor under Impact of Ogive-Nose Projectile**

**Li Ding 1,2,\* , Xiaohui Gu <sup>1</sup> , Peihui Shen <sup>1</sup> and Xiangsheng Kong <sup>2</sup>**


**Abstract:** The ballistic response of armor has been widely used to evaluate its feasibility and advantages as a protective structure. To obtain the ballistic performance and ballistic limitations of composite armor, a type of ultra–high molecular weight polyethylene (UHMWPE) composite armor is designed, which is composed of UHMWPE laminates and steel face sheets of Q235. The total thickness of the armor is 53 mm, with an in–plane dimension of 300 mm × 300 mm. Then, an experimental study of the ballistic impact response of composite armor subject to a typical ogive–nose projectile was carried out. In the velocity range of 501.1 to 1026.1 m/s, the 14.5 mm caliber armor– piercing projectile could penetrate through the composite armor. At the velocity of 433.3 m/s, the A–P projectile was embedded in the armor, leaving a bulge mark on the back sheet. Therefore, 467.2 m/s is taken as the ballistic limit of the armor under the impact of the ogive–nose projectile. In addition, a corresponding numerical simulation model is also established to predict the ballistic limit of the projectile. The numerical predictions are consistent with the experimental results. The ballistic limit obtained from the numerical simulation results is 500 m/s, which is acceptable with a relative error of 7.02%. The failure mechanism of the composite armor is also obtained. Petaling is the main dominant failure mode for both face sheets, while delamination and shear failure dominate the penetration process of UHMWPE laminates. Finally, the perforation mechanism of composite armor under the impact of an A–P projectile is analyzed with theoretical models to predict the residual velocity, the work performed during the perforation, and the resisting stress of *σ*s in the cavity–expansion model. The experimental and numerical simulation results can provide necessary data in the analysis of the composite structure's dynamic response under the impact of sharp head penetrators. The research results present the ballistic performance, failure mechanism, and ballistic limit of the composite armor under the impact of a typical ogive–nose projectile, which can be significant in the design of composite armor in the areas of ship shield, fortifications protection, and bulletproof structures against threats from sharp head penetrators.

**Keywords:** ballistic limit; composite armors; ogive–nose projectile; theoretical analysis; ultra–high molecular weight polyethylene

## **1. Introduction**

Multilayered armor systems have been widely used in numerous ballistic and bulletproof applications, including bulletproof helmets, vests, and other armor parts, providing an acceptable range of protection for soldiers and structures [1–6]. Due to the advantages of high stiffness and low density, composites reinforced with ultra–high molecular weight polyethylene (UHMWPE) fibers are increasingly being used in the field of national defense as parts of lightweight armor systems to protect fortifications and structures from ballistic impacts [7–9]. The ballistic performance of monolithic UHMWPE composite under the impact of a blunt projectile has been studied [10–14], and the failure mechanisms analyzed. A common design of composite armor is a sandwich structure, which is made of thin face sheets and low–density non–metal cores. In addition, the ballistic impact response

**Citation:** Ding, L.; Gu, X.; Shen, P.; Kong, X. Ballistic Limit of UHMWPE Composite Armor under Impact of Ogive-Nose Projectile. *Polymers* **2022**, *14*, 4866. https://doi.org/10.3390/ polym14224866

Academic Editors: R. A. Ilyas, S. M. Sapuan, Emin Bayraktar, Shukur Abu Hassan, Nabil Hayeemasae and Khubab Shaker

Received: 13 October 2022 Accepted: 8 November 2022 Published: 11 November 2022

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2022 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

of a sandwich structure consisting of UHMWPE under a blunt projectile has also been investigated [15–19]. Deflection and bulging [4] consisting of shear plugging, formation of a transition plane, and bulging are the failure modes of UHMWPE [11], which allow it to have an excellent ability to resist the penetration of blunt projectiles.

It is reported that the bulletproof properties of composite materials are greatly discounted under the ballistic impact of projectiles with a sharp head [17,20,21]. However, little is reported on the ballistic performance of the UHMWPE or composite armors containing UHMWPE laminates under the impact of sharp head projectiles. Whether the ballistic performance of the composite armor can be predicted using numerical simulation has not been discussed yet. In addition, to evaluate the resistance of composite armors, especially with the sandwich structure containing UHMWPE laminates, the failure mode and failure mechanism of UHMWPE should also be further studied. An experimental study is one of the basic methods of scientific research, which provides the most direct evidence to help understand physical phenomena. However, the penetration process of the ballistic impact can hardly be observed in experiments. Therefore, due to the limits of testing conditions and the costs of the investment, it is difficult to carry out the relevant research that depends on experiments. Numerical simulations are similar to experimental research that is based on theoretical research [10,14,22]; they are efficient but irrelevant in external conditions. With the help of computer realization, the laws of penetration and dynamic response of the target can be obtained with accurate simulation models. The intermediate process in the penetration process could also be observed, and the crucial parameters could be extracted in the post–processing section to help understand the mechanism of the physical experiment better.

In this paper, a sandwich structure of UHMWPE composite armor is designed, which is made of two pieces of UHMWPE laminates in the middle and Q235 steel face sheets. The ballistic performance of the composite armor is studied systematically both from experiment and numerical simulation. The ballistic limit of UHMWPE composite armor under the impact of an ogive–nose penetrator is obtained, and the failure mechanism of the armor is analyzed. Finally, typical theoretical models are used to predict the residual velocity, the work performed during the perforation, and the resisting stress of *σ*<sup>s</sup> in the cavity–expansion model to help better understand the penetration process of the composite armor under the impact of sharp head projectiles.

#### **2. Configuration of the Armor and Projectile**

#### *2.1. Design of the Armor*

As shown in Figure 1, the UHMWPE composite armor is made up of two pieces of UHMWPE laminates in the middle and Q235 steel face sheets. Typical UHMWPE laminate with a material grade of FDB4-HW-S1 is selected. Each piece of UHMWPE laminate has a thickness of 20 mm, and each Q235 steel face sheet has a thickness of 6 mm. Each layer of armor has the same in–plane dimensions of 300 mm × 300 mm, with a total thickness of 52 mm. The thin binder layer is replaced in each panel, then followed by pressing to obtain the overall panel structure of the composite armor. Due to the existence of a binder layer, the total thickness of the composite armor may increase from 1 mm to 53 mm.

**Figure 1.** Structures of UHMWPE composite armor. (**a**) Assembly drawing. (**b**) Component draw-**Figure 1.** Structures of UHMWPE composite armor. (**a**) Assembly drawing. (**b**) Component drawing.

The material properties of Q235 steel are presented in Table 1, which is provided by The material properties of Q235 steel are presented in Table 1, which is provided by the manufacturer of Wuhan Iron & Steel Co., Ltd., Wuhan, China.

the manufacturer of Wuhan Iron & Steel Co., Ltd., Wuhan, China.

**Table 1.** Material properties of Q235.

ing.

out first.


#### *2.2. Structure of the Projectile*

*2.2. Structure of the Projectile* The structure of the standard Chinese projectile of 14.5 mm caliber armor–piercing (A–P) projectile is shown in Figure 2. The main geometric parameters of the projectile are presented in Figure 2a, with a total diameter of 14.93 mm and a length of 66.7 mm. The projectile is mainly composed of the brass jacket, steel core, and lead filler. The steel core with an ogive–nose head is made of a non–deformable hardened steel core, which will mainly contribute to the penetration performance of the projectile. The A–P core has a diameter of 12.48 mm and a length of 53.4 mm, with a mass of 40.2 to 41 g. Generally, the jacket is made of gliding brass, and the filler is made of lead, which serves to protect the barrel from the core, enhance the sealing effect and provide an optimized shape for flight in the air. They have a relatively small effect during the penetration, and they are not The structure of the standard Chinese projectile of 14.5 mm caliber armor–piercing (A–P) projectile is shown in Figure 2. The main geometric parameters of the projectile are presented in Figure 2a, with a total diameter of 14.93 mm and a length of 66.7 mm. The projectile is mainly composed of the brass jacket, steel core, and lead filler. The steel core with an ogive–nose head is made of a non–deformable hardened steel core, which will mainly contribute to the penetration performance of the projectile. The A–P core has a diameter of 12.48 mm and a length of 53.4 mm, with a mass of 40.2 to 41 g. Generally, the jacket is made of gliding brass, and the filler is made of lead, which serves to protect the barrel from the core, enhance the sealing effect and provide an optimized shape for flight in the air. They have a relatively small effect during the penetration, and they are not modeled in the following numerical simulation [23–25].

modeled in the following numerical simulation [23–25]. The ballistic limit or limit velocity is the velocity required for a particular projectile to reliably penetrate a particular piece of material. In other words, a given projectile will not pierce a given target when the projectile velocity is lower than the ballistic limit [26]. In addition, it is also important to evaluate the resistance of the armor. Due to the unknown perforation mechanism and the unknown ballistic limit of composite armor under the ballistic impact, especially the sharply pointed projectile, an experiment of the The ballistic limit or limit velocity is the velocity required for a particular projectile to reliably penetrate a particular piece of material. In other words, a given projectile will not pierce a given target when the projectile velocity is lower than the ballistic limit [26]. In addition, it is also important to evaluate the resistance of the armor. Due to the unknown perforation mechanism and the unknown ballistic limit of composite armor under the ballistic impact, especially the sharply pointed projectile, an experiment of the UHMWPE composite armor under the ballistic impact of the A–P projectile was carried out first.

UHMWPE composite armor under the ballistic impact of the A–P projectile was carried

**Figure 2.** Structure of a 14.5 mm caliber armor–piercing projectile. (**a**) Geometric structure. (**b**) Brass jacket. (**c**) Steel core. (**d**) Lead filler. **Figure 2.** Structure of a 14.5 mm caliber armor–piercing projectile. (**a**) Geometric structure. (**b**) Brass jacket. (**c**) Steel core. (**d**) Lead filler.

#### **3. Experimental Details and Results 3. Experimental Details and Results**

#### *3.1. Design of the Experiment 3.1. Design of the Experiment*

Figure 3 shows the 14.5 mm caliber A–P projectile and the state of the assembly of the projectile in the cartridge. By adjusting the quantity of the propellent in the cartridge, the pre–set velocities of the A–P projectile can be acquired. Figure 3 shows the 14.5 mm caliber A–P projectile and the state of the assembly of the projectile in the cartridge. By adjusting the quantity of the propellent in the cartridge, the pre–set velocities of the A–P projectile can be acquired.

*Polymers* **2022**, *14*, x FOR PEER REVIEW 5 of 22

**Figure 3.** State of the projectile in the experiment. (**a**) 14.5 mm caliber A–P projectile. (**b**) Assembly of the projectile in the cartridge. **Figure 3.** State of the projectile in the experiment. (**a**) 14.5 mm caliber A–P projectile. (**b**) Assembly of the projectile in the cartridge. (**b**) **Figure 3.** State of the projectile in the experiment. (**a**) 14.5 mm caliber A–P projectile. (**b**) Assembly of the projectile in the cartridge.

Figure 4 shows the states of the composite armors used in the experiment. The armors were clamped to the rear base on the steel shelf. The rear base was a triangle–shaped bracket, which was fixed to the steel shelf with bundles of iron wire. Figure 5 presents the electronic time–measuring instrument with six channels to capture the signal when the penetrator perforates through the tinfoil target. The electronic time–measuring instrument is of high sensitivity and is able to record the minimum time difference of 1 μs. The layout of the ballistic impact experiment is presented in Figure 6. The 14.5 mm caliber smooth–bore gun was mounted on a rigid platform. The distance between the gun muzzle and armor is about 4 m. To measure the impact velocity of the projectile, two tinfoil targets were placed in front of the armor. The composite armor was placed on a steel shelf at the same height as the ballistic gun. Figure 4 shows the states of the composite armors used in the experiment. The armors were clamped to the rear base on the steel shelf. The rear base was a triangle–shaped bracket, which was fixed to the steel shelf with bundles of iron wire. Figure 5 presents the electronic time–measuring instrument with six channels to capture the signal when the penetrator perforates through the tinfoil target. The electronic time–measuring instrument is of high sensitivity and is able to record the minimum time difference of 1 µs. The layout of the ballistic impact experiment is presented in Figure 6. The 14.5 mm caliber smooth–bore gun was mounted on a rigid platform. The distance between the gun muzzle and armor is about 4 m. To measure the impact velocity of the projectile, two tinfoil targets were placed in front of the armor. The composite armor was placed on a steel shelf at the same height as the ballistic gun. Figure 4 shows the states of the composite armors used in the experiment. The armors were clamped to the rear base on the steel shelf. The rear base was a triangle–shaped bracket, which was fixed to the steel shelf with bundles of iron wire. Figure 5 presents the electronic time–measuring instrument with six channels to capture the signal when the penetrator perforates through the tinfoil target. The electronic time–measuring instrument is of high sensitivity and is able to record the minimum time difference of 1 μs. The layout of the ballistic impact experiment is presented in Figure 6. The 14.5 mm caliber smooth–bore gun was mounted on a rigid platform. The distance between the gun muzzle and armor is about 4 m. To measure the impact velocity of the projectile, two tinfoil targets were placed in front of the armor. The composite armor was placed on a steel shelf at the same height as the ballistic gun.

**Figure 4.** Photograph of the armor in the experiment. (**a**) front view. (**b**) side view. **Figure 4.** Photograph of the armor in the experiment. (**a**) front view. (**b**) side view.

**Figure 4.** Photograph of the armor in the experiment. (**a**) front view. (**b**) side view.

**Figure 5.** Photograph of the electronic time–measuring instrument. **Figure 5.** Photograph of the electronic time–measuring instrument. **Figure 5.** Photograph of the electronic time–measuring instrument.

(**b**) (**b**)

**Figure 6.** Configuration of the ballistic impact experiment. (**a**) 14.5 mm caliber smooth–bore gun. (**b**) Layout of the targets. **Figure 6.** Configuration of the ballistic impact experiment. (**a**) 14.5 mm caliber smooth–bore gun. (**b**) Layout of the targets. **Figure 6.** Configuration of the ballistic impact experiment. (**a**) 14.5 mm caliber smooth–bore gun. (**b**) Layout of the targets.

#### *3.2. Experimental Results 3.2. Experimental Results 3.2. Experimental Results 3.2. 3.2. Experimental Results*

REVIEW

The perforation results of composite armor by the A–P projectile are listed in Table 2. For the unknown perforation mechanism and the unknown ballistic limit of composite armor under the ballistic impact of the sharply pointed projectile especially, the impact velocities *v*<sup>i</sup> of the ogive–nose projectile were set at 1026 m/s initially to much lower velocities. In the velocity range of 501.1 to 1026.1 m/s, the A–P projectile could penetrate through the composite armor. While at the velocity of 433.3 m/s, the A–P projectile was embedded in the armor, leaving a bulge mark on the back sheet. The diameter of the entrance hole on the front sheet was about 14–16 mm, while the diameter of the outlet on the back sheet was about 13–24 mm. The entrance dimension discrepancies are small on the front sheet, and the divergence increased with the increase in impact velocity. The perforation results of composite armor by the A–P projectile are listed in Table 2. For the unknown perforation mechanism and the unknown ballistic limit of composite armor under the ballistic impact of the sharply pointed projectile especially, the impact velocities *v*<sup>i</sup> of the ogive–nose projectile were set at 1026 m/s initially to much lower velocities. In the velocity range of 501.1 to 1026.1 m/s, the A–P projectile could penetrate through the composite armor. While at the velocity of 433.3 m/s, the A–P projectile was embedded in the armor, leaving a bulge mark on the back sheet. The diameter of the entrance hole on the front sheet was about 14–16 mm, while the diameter of the outlet on the back sheet was about 13–24 mm. The entrance dimension discrepancies are small on the front sheet, and the divergence increased with the increase in impact velocity. The perforation results of composite armor by the A–P projectile are listed in Table 2. For the unknown perforation mechanism and the unknown ballistic limit of composite armor under the ballistic impact of the sharply pointed projectile especially, the impact velocities *v*<sup>i</sup> of the ogive–nose projectile were set at 1026 m/s initially to much lower velocities. In the velocity range of 501.1 to 1026.1 m/s, the A–P projectile could penetrate through the composite armor. While at the velocity of 433.3 m/s, the A–P projectile was embedded in the armor, leaving a bulge mark on the back sheet. The diameter of the entrance hole on the front sheet was about 14–16 mm, while the diameter of the outlet on the back sheet was about 13–24 mm. The entrance dimension discrepancies are small on the front sheet, and the divergence increased with the increase in impact velocity. The perforation results of composite armor bytheA–<sup>P</sup> projectile are listed in Table 2. For the unknown perforation mechanism and the unknown ballistic limit of composite armor ballistic sharply pointed projectileespecially, velocities <sup>i</sup> the ogivenose at 1026 m/s initially much velocities. velocity range m/s, Acould through the While the of embedded in the armor, leaving a bulge mark on the back sheet. The diameter of the entrance on the front sheet was 1416 mm, while the on back was about entrance dimension are small on the front sheet, and increase velocity. results composite the mechanism the unknown armor the impact of the sharplyprojectile especially, impact <sup>i</sup>projectile wereset 1026 initially to lower locities. In the velocity range of 501.1 to 1026.1 m/s, the A–Pprojectile could penetrate through the composite armor. While at the velocity of 433.3 m/s, the A–<sup>P</sup> projectile was armor, bulge mark on back sheet. of the mm, of outlet sheet about 13The entrance on the front sheet, and the divergence increasedwith the increase inimpact velocity.

**Table 2.** Perforation results in the experiment. **Table 2.** Perforation results in the experiment. **Table 2.** Perforation results in the experiment. **Table 2.** Perforation results in the experiment. the


#### **Table 2.** *Cont. Polymers* **2022**, *14*, x FOR PEER REVIEW 8 of 22 *Polymers* **2022**, *14*, x FOR PEER REVIEW 8 of 22 *Polymers* **2022**, *14*, x FOR PEER REVIEW 8 of 22 *Polymers* **2022**, *14*, x FOR PEER REVIEW 8 of 22 *Polymers* **2022**, *14*, x FOR PEER REVIEW8 of 22 *Polymers* **2022**,*14*, x FOR PEER REVIEW8of 22


#### **Table 2.** *Cont. Polymers* **2022**, *14*, x FOR PEER REVIEW 9 of 22 *Polymers* **2022**, *14*, x FOR PEER REVIEW 9 of 22 *Polymers* **2022**, *14*, x FOR PEER REVIEW 9 of 22 *Polymers* **2022**, *14*, x FOR PEER REVIEW 9 of 22

*Polymers* **2022**, *14*, x FOR PEER REVIEW 9 of 22

*Polymers* **2022**, *14*, x FOR PEER REVIEW 9 of 22
