**Biointerface Coatings for Biomaterials and Biomedical Applications**

Editors

**Hsien-Yeh Chen Peng-Yuan Wang**

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

*Editors* Hsien-Yeh Chen National Taiwan University Taiwan

Peng-Yuan Wang Chinese Academy of Sciences China

*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 *Coatings* (ISSN 2079-6412) (available at: https://www.mdpi.com/journal/coatings/special issues/ biointerface coat biomater biomed appl).

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LastName, A.A.; LastName, B.B.; LastName, C.C. Article Title. *Journal Name* **Year**, *Volume Number*, Page Range.

**ISBN 978-3-0365-2243-2 (Hbk) ISBN 978-3-0365-2244-9 (PDF)**

© 2022 by the authors. Articles in this book are Open Access and distributed under the Creative Commons Attribution (CC BY) license, which allows users to download, copy and build upon published articles, as long as the author and publisher are properly credited, which ensures maximum dissemination and a wider impact of our publications.

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## **Contents**

Reprinted from: *Coatings* **2020**, *10*, 427, doi:10.3390/coatings10040427 ................ **109**

## **About the Editors**

## **Hsien-Yeh Chen**

Hsien-Yeh Chen's research focuses on vapor-based polymers for biomaterials and biomedical devices. These polymers have been successfully applied in bone and tooth implants, drug-eluting stents, intraocular lenses, micro- and nanocolloids, nanomedicine carriers, and bioelectronic devices. Polymer technology enables biomaterials to be produced with properties including antifouling and antibacterial activities, cell differentiation manipulations, superhydrophobic interfaces, and surface patterns. Prof. Chen was the receiver of the Ta-You Wu Memorial Award in 2017, the National Award for the Innovation in Biotechnology of Taiwan in 2016, and the Green Chemistry Application and Innovation Awards in 2021.

## **Peng-Yuan Wang**

Peng-Yuan Wang received his Ph.D. from the Department of Chemical Engineering, National Taiwan University, in 2011. Dr Wang is currently the Principal Investigator at Oujiang Laboratory, Zhejiang, China. He was the Director of Shenzhen Key Laboratory, the Principal Investigator at the Shenzhen Institute of Advanced Technology (SIAT), Chinese Academy of Sciences (CAS), and the lecturer at Swinburne University. Dr Wang has been awarded several prestigious awards, including the Thousand Talent Fellowship in China; the Discovery Early Career Researcher Award (DECRA) and the Veski Victoria Fellowship in Australia. His research focuses on manipulating stem cell behavior using biophysical cues such as nanostructured patterns and cell compression, which provides new insights in mechanobiology, biomaterials, and regenerative medicine. Dr Wang has authored over 70 peer-reviewed SCI publications, 3 book chapters, and applied for/been issued over 10 patents worldwide. His works provide an essential framework for understanding biophysical effects on stem cells and the development of next-generation biomaterials and tools for regenerative medicine.

## **Preface to "Biointerface Coatings for Biomaterials and Biomedical Applications"**

In addition to meeting the minimal requirements for biocompatibility, advanced biomaterials have acquired functions allowing them to directly or indirectly influence specific biological environments. Modifications to biomaterials are elegantly achieved by establishing a biointerface layer to deliver the desired functions, but the definition and configuration required to balance such an interface between the exploited materials or devices and the encountered biological microenvironment pose challenges. The design of a successful biointerface usually depends on criteria such as controlled presentation of functional biomolecules on the surface, low nonspecific protein adsorption, responsive actions toward external stimuli, multifunctionality, compatibility with micro- to nanofabrication, surface morphology or microstructures, biodegradability, and physical or chemical gradients. Many promising approaches have been reported using existing surface modification technologies based on both physical and chemical methods for rendering fabricated coatings on biomaterials, from basic self-assembly of molecules to top-down construction of bulk materials. The development of technologies in this field is pursued by exploring complimentary and/or combinatorial strategies based on both existing and newly discovered methods, and these technologies are paving the way for advanced and effective functional biointerfaces for prospective biomaterials. This eBook provides insights into current obstacles in the development of biointerfaces for biomaterials and discuss current progress with examples of solutions to interface problems, and this work is suitable for audiences interested in biology and materials science who seek solutions to biointerface problems and challenges. The Editors are grateful for the contributing authors and their input regarding research and prospective insights for the eBook, the funding support (detailed in the Editorial) leading to the publication of the eBook, and finally the publication resources provided by MDPI Coatings.

• Surface coatings are useful for changing surface properties according to the requirements for numerous marine and biological applications. Surface coatings modified by vapor phase, plasma, dry-/wet-etching, and solutions also provide considerable advantages in forthcoming industrial technology.

• For this eBook, we invited experts from different fields and describe their latest technology in detail for material selection, deposition, characterization, and evaluation.

• In the editorial of the eBook, two Guest Editors summarize the latest technology and applications of surface coatings in biomedical engineering. They split the surface properties into three dominant factors, i.e., topography, chemistry, and stiffness, and discuss cell–material interactions at biointerfaces.

• In the review paper in this eBook, the surface coatings of orthopedic implants are discussed and summarized. The authors also highlight the currently available nanomaterial-based surface modification technologies available to augment the function and performance of these metallic bioimplants in a clinical setting.

• In the next two feature papers, bioceramic coatings fabricated by plasma spray for bone implants and microbicidal coatings for antimicrobial surfaces were investigated. The first paper shows that Sr–HT–G and HAp coatings both have good biocompatibility for bone marrow mesenchymal stem cells (BMSCs), whereas the second paper presents a new surface design for the production of durable microbicidal coatings. These two applications will be widely useful in clinics and beneficial in next-generation implants.

• In the next four research papers, the authors report different surface coating methods and characterizations, including chemical vapor sublimation and deposition (CVSD), micro-arc oxidation, magnetic poly(vinyl alcohol) (mPVA) gels, and matrix-assisted laser desorption/ionization-time of flight mass (MALDI-ToF mass) spectroscopy. These studies show the diversity of surface coatings and their applications.

• In the next two research papers, the authors focus on coatings for bone tissue engineering. The first paper suggests that the established NT-ICA-ASP/PLGA substrate is a promising candidate for functionalized coating materials in Ti implant surface modification. The second paper shows that PPy/PDA NW coating has better biocompatibility and bioactivity than pure PPy NWs, and PDA and has benefits for the adhesion, proliferation, and osteogenic differentiation of MC3T3-E1 cells cultured on the surface. In addition, PPy/PDA NWs can significantly promote the osteogenesis of MC3T3-E1 cells in combination with micro-galvanostatic electrical stimulation (ES).

• In the last two research papers, corrosion properties and anticorrosion coatings are studied and developed. The first paper reports that UMAO/PLGA/BR coatings have excellent biological activity, which can effectively solve the clinical problem of the rapid degradation of pure magnesium and easy infection. The second paper was inspired by mangrove leaves that have salt glands that can secrete excessive ions to control ion transport in and out. The authors fabricated bipolar, hydrophobic coatings by doping ion-selective resins and constructing surface structures, which restrict the transport of corrosive substances and electrochemical corrosion at the coating/metal interface. They showed that the bioinspired coatings effectively blocked and controlled the transport of both Na+ and Cl, and together with the hydrophobic surface, the coating system exhibited significantly improved anticorrosion properties, a more than three-fold decrease in corrosion current density compared with the control group (epoxy varnish).

• Surface coatings have a decisive impact on a material and are constantly evolving with new needs. It is expected that the above studies will be informative to the reader.

> **Hsien-Yeh Chen and Peng-Yuan Wang** *Editors*

## *Editorial* **Special Issue: Biointerface Coatings for Biomaterials and Biomedical Applications**

**Hsien-Yeh Chen 1,\* and Peng-Yuan Wang 2,3,\***

The success of recent material science and applications in biotechnologies should be credited to developments of malleable surface properties, as well as the adaptation of conjugation reactions to the material surface [1–4]. An article additionally addressing the progress and challenges of biointerface modifications using integrated nanomaterial bioimplants for orthopedic applications was also thoroughly reviewed in the current Special Issue by Ahirwar et al [5]. A biointerface is the region of contact between a biomolecule, cell, biological tissue, or living organism considered living with another biomaterial or inorganic/organic material. The surface property of a material is highly important, even compared to the bulk property [6]. For example, a hydrogel is soft (e.g., 1–100 kPa) but could have a stiffer thin layer on the top surface (>100 kPa) that modulates cell adhesion at the biointerface. On the other hand, rigid materials can also be modified with a softer surface coating to enhance biocompatibility. Therefore, surface modification or coatings have immense potential and wide application in biomedical devices and implants [7]. The crucial biointerface properties of using biomaterials include the ability to control the presentation of biomolecules with precisely defined chemical topology, the ability to control and suppress undesirable background noise from nonspecific biomolecules (e.g., protein) adsorptions, smart response with respect to environmental stimuli, multiple functions that are simultaneously activated, and surface gradients with gradual and cascade guidance from physical and/or chemical cues. These properties may subsequently lead to successful biodevice/material performance and efficacy. In the article "*Proteomic Analysis of Biomaterial Surfaces after Contacting with Body Fluids by MALDI-TOF Mass Spectroscopy*", Hirohara et al. developed an analysis method to evaluate the composition of protein absorbed on solid surfaces while preventing the effects of pipetting artifacts. The method includes denaturation, reduction, alkylation, digestion, and spotting of the matrix followed by matrix-assisted laser desorption/ionization-time of flight mass (MALDI-TOF mass) spectroscopy. The authors also developed an algorithm to evaluate the adsorbed proteins by collating the experimental and theoretical peak positions of fragmented peptides. Their results showed the mechanism of how the cell and tissues are affected after biomaterial contact [8].

Furthermore, biomolecular engineering technologies have enabled the successful use of these polymer materials. The use of physical approaches or chemical means to install biological functions onto polymer materials is of interest within this area of research. From a physical point of view, the ability to control biomolecules at the solid/liquid interface requires adequate knowledge and understanding of surface interactions, transport phenomena of interacting molecules, interactions with external stimuli, and surface functional groups. From a chemical point of view, conjugation reactions seek a fast reaction time, mild reaction conditions, and, more importantly, specificity to achieve successful conjugation in the vast array of functionalities present in biological microenvironments [9].

**Citation:** Chen, H.-Y.; Wang, P.-Y. Special Issue: Biointerface Coatings for Biomaterials and Biomedical Applications. *Coatings* **2021**, *11*, 423. https://doi.org/10.3390/coatings 11040423

Received: 31 March 2021 Accepted: 2 April 2021 Published: 6 April 2021

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

**Copyright:** © 2021 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/).

Recently demonstrated and promising concepts have focused on the creation of multiple surface functionalities on material surfaces [10–15]. The approaches consider the physical and chemical surface properties while delivering cascading and/or simultaneous activities to respond to sophisticated bioenvironments. Moreover, the need to precisely incorporate biomolecules at specific locations on a micro/nanoscale, i.e., in confined micro/nanodomains and to induce topographically derived responses toward biological environments, has also become essential. These concepts have fueled modern schemes for the design of prospective polymer materials for biological applications as well as more advanced developments in biointerface science. In the article "*Facile Route of Fabricating Long-Term Microbicidal Silver Nanoparticle Clusters against Shiga Toxin-Producing Escherichia coli O157:H7 and Candida auris*", Gangadoo et al. reported a facile coating technique using silver nanoparticles (Ag NPs) to fabricate a surface with long-term microbiocidal activity. The Ag NP coating was fabricated on copper surfaces via an ion-exchange and reduction reaction followed by a silanization step, resulting in high-aspect-ratio Ag NP clusters. Their results demonstrated the durability and high efficiency of microbiocidal activity against both Shiga toxin-producing *E. coli* and *C. auris* cells [16]. In the article "*One-Step Preparation of Nickel Nanoparticle-Based Magnetic Poly (Vinyl Alcohol) Gels*", Li et al. established a onestep synthetic method of magnetic poly (vinyl alcohol) (mPVA) gels. Using this method, multiresponsive gels were obtained by incorporating Ni nanoparticles (NPs) into a stimuliresponsive polymer. Ni-NP PVA gels are anticipated to be applied for controlled drug delivery systems, especially anticancer applications. This novel and facile method can be utilized for magnetic gels for biotechnology [17]. In the article "*Mangrove Inspired Anti-Corrosion Coatings*", Cui et al. discovered that well-controlled transport of corrosive substances is the critical key to anti-corrosion performance. The authors demonstrated a bipolar hydrophobic coating that can effectively block and control the transport of both Na + and Cl-, exhibiting significantly improved anti-corrosion properties and a more than three orders of magnitude decrease in corrosion current density. These bioinspired coatings may lead to outstanding and long-term anti-corrosion performance [18].

Surface topography, including nanostructure and roughness, is one of the critical parameters of materials at biointerfaces. Advanced technologies, including electron beam lithography, laser ablation, and electrochemical etching, have been developed to generate surface nanotopographies [19,20]. In the article "*Vapor-Stripping and Encapsulating to Construct Particles with Time-Controlled Asymmetry and Anisotropy*", Wu et al. demonstrated the fabrication method of particles with asymmetric and anisotropic structures by utilizing a time-controlled vapor stripping and encapsulating process. The results showed that the innovative process, chemical vapor sublimation and deposition (CVSD), enabled sensitive soybean agglutinin (SBA) protein tubes to be encapsulated in poly-p-xylylene particles. The SBA protein tubes retained their original morphology and could be used to construct particles with asymmetric and anisotropic structures. In addition, the size of the particles could be predicted and controlled. The CVSD process is a promising strategy for fabricating particles [21]. Nevertheless, these methods are material-dependent. For example, silicon has been commonly used in lithography and etching methods. Nanostructures have been used to control cell adhesion and differentiation. The outcomes depend on the size, geometry, arrangement, density, and, not surprisingly, cell type. In general, nanostructures are biocompatible, cell adhesive, low bacterial adhesive, low inflammatory, and biofunctional. For example, nanostructures on a bone implant should induce a mild inflammatory effect and good bone cell affinity and bone integration. In the article "*Icariin/Aspirin Composite Coating on TiO2 Nanotubes Surface Induce Immunomodulatory Effect of Macrophage and Improve Osteoblast Activity*", Ma et al. used an aspirin (ASP)/poly (lactic–co–glycolic acid) (PLGA) coating on icariin (ICA)-loaded TiO2 nanotubes (NT-ICA-ASP/PLGA). Compared to those cultured on the Ti surface, macrophages on the NT-ICA-ASP/PLGA substrate displayed decreased M1 proinflammatory and enhanced M2 pro-regenerative gene and protein expression, which implied an activated immunomodulatory effect. Moreover, when cultured with conditioned medium from macrophages, osteoblasts on the NT-ICA-ASP/PLGA

substrate showed improved cell proliferation, adhesion and osteogenic gene and protein expression compared with those on the Ti surface. These results suggested that the NT-ICA-ASP/PLGA substrate is a promising candidate for functionalized coating material in Ti implant surface modification [22]. In the article "*The Bioactive Polypyrrole/Polydopamine Nanowire Coating with Enhanced Osteogenic Differentiation Ability with Electrical Stimulation*", He et al. used a two-step method to construct a functional conductive coating of polypyrrole/polydopamine (PPy/PDA) nanocomposites for bone regeneration. The PPy/PDA NW coating exhibited better biocompatibility and bioactivity than pure PPy NWs and PDA and was beneficial for the adhesion, proliferation, and osteogenic differentiation of MC3T3-E1 cells cultured on the surface. In addition, PPy/PDA NWs significantly promoted the osteogenesis of MC3T3-E1 cells in combination with micro galvanostatic electrical stimulation (ES) [23].

Surface chemistry, including material chemistry, biomolecule grafting, and chemical coatings, is another critical parameter of materials at biointerfaces [24–26]. Surface chemistry represents functional groups, electrostatic properties, and wettability on the material surface. These properties are crucial in interactions with proteins, cells, and tissues. For example, the carboxy group can stimulate osteogenic differentiation of mesenchymal stem cells (MSCs). Uniform coatings such as spin coating provide a thin layer of polymer (~submicrons) on materials. The adhesive force between the thin layer and substrate is essential for the long-term application of the materials. Nonuniform coatings can be achieved using dewetting methods, which in turn generate nano or micrometer-scale features. In the article "*Corrosion Behavior and Biological Activity of Microarc Oxidation Coatings with Berberine on a Pure Magnesium Surface*", Mu et al. reported a coating named ultrasonic microarc oxidation/polylactic acid and glycolic acid copolymer/berberine (UMAO/PLGA/BR) on a pure magnesium substrate for bone materials. Different amounts of berberine (BR) can seal the corrosion channel to different extents. These coatings have a self-corrosion current density (Icorr) at least one order of magnitude lower than that of the UMAO coatings. When the BR content was 3.0 g/L, the self-corrosion current density of the UMAO/PLGA/BR coatings was the lowest (3.14 × <sup>10</sup>−<sup>8</sup> A/cm2), and the corrosion resistance was improved. UMAO/PLGA/BR coatings have excellent biological activity, which can effectively solve the clinical problem of rapid degradation of pure magnesium and easy infection [27]. In the report "*Chemical and Biological Roles of Zinc in a Porous Titanium Dioxide Layer Formed by Micro-Arc Oxidation*" by Shimabukuro et al., zinc incorporated by microarc oxidation (MAO) in porous titanium dioxide was investigated under physiological saline conditions. The time transient state from zinc to zinc oxide led to early stage release in 7 days and antibacterial ability after 28 days of incubation. Additionally, there was no interruption of osteogenic cell proliferation and calcification in zinc specimens. In conclusion, timetransient zinc not only gives antibacterial properties but also shows great compatibility with osteogenic cells and has great potential in chemical and biological fields [28].

Surface stiffness in the material top surface (~submicrons) is crucial in controlling cell adhesion, the cytoskeleton, and cell differentiation [29]. In vivo, the stiffness of the extracellular matrix (ECM) ranges from a few kPa (e.g., brain) to GPa (e.g., bone). Thus, it makes sense that surfaces with different stiffnesses could mimic the native ECM and guide cells to become a specific cell type. The top surface of materials is easily oxidized, altering the stiffness of the surface. Because the top surface has a more dominant effect on cell behavior than the bulk material, technology with higher resolution is needed to analyze the mechanical properties of the top layer. In the article "*Mechanical Properties of Strontium–Hardystonite–Gahnite Coating Formed by Atmospheric Plasma Spray*", Pham et al. measured the mechanical properties and tested the cell viability of a bioceramic coating, strontium–hardystonite–gahnite (Sr–HT–G, Sr–Ca2ZnSi2O7–ZnAl2O4), to evaluate the potential use of this novel bioceramic for bone regeneration applications contrasted to the properties of the well-known commercial standard coating of hydroxyapatite (HAp: Ca10(PO4)6(OH)2). The Sr–HT–G coating exhibited a more uniform distribution of hardness and elastic moduli across its cross-section compared to HAp. The Sr–HT–G coating also

revealed higher microhardness, nanohardness and elastic moduli than those shown for the HAp coating. The nanoscratch tests for the Sr–HT–G coating presented a low volume of material removal without high plastic deformation. Furthermore, the Sr–HT–G coating had a lower wear volume than the HAp. The Sr–HT–G coating had a slightly higher cell attachment density and spreading area of bone marrow mesenchymal stem cells (BMSCs) than the HAp coating [30].

Finally, biointerfaces often contain a combination of surface topography, chemistry, and stiffness. These three properties are sometimes difficult to distinguish. For example, the generation of surface nanotopography could change the surface wettability, and polymer coating could alter the surface nanotopography. Recently, a new family of twodimensional (2D) materials called colloidal self-assembled patterns (cSAPs), composed of different particles, has been developed [31,32]. The formed particle patterns can be hexagonal, close-packed, or randomly distributed. The topography of cSAPs can be tuned by controlling particle-particle interactions. It is easy to decorate biosignals on cSAPs using pre- or postmodification of particles. cSAPs can be used as substrates, coatings, and free-standing membranes, depending on the application. cSAPs can be the next-generation material at biointerfaces. As more stringent specifications are required for designing the surface properties of prospective materials and the development of new devices is pursued with complicated geometries and minimized sizes, the surface properties of such materials/devices now also require a more defined and flexible presentation of the chemical functionalities (e.g., multifunctional or gradient distribution) and the precise confinement of these chemical conducts in relevant locations of interest [33]. The emerging applications of the existing technologies and/or new technologies from two dimensions into more sophisticated three-dimensional regions [34] are challenging the field of biointerfaces science, and further research developments are expected on this path.

**Author Contributions:** Conceptualization, H.-Y.C. and P.-Y.W.; writing—original draft preparation, H.-Y.C. and P.-Y.W.; writing—review and editing, H.-Y.C. and P.-Y.W.; supervision, H.-Y.C. and P.-Y.W.; funding acquisition, H.-Y.C. and P.-Y.W. All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by the Ministry of Science and Technology of Taiwan (MOST 108-2221-E-002-169-MY3; 108-2218-E-007-045; 109-2314-B-002-041-MY3; 109-2634-F-002-042). Funding support also came from the Ministry of Science and Technology of China (2019YFE0113000); the National Natural and Science Foundation of China (31870988); the Chinese Academy of Sciences (172644KYSB20200002, 172644KYSB20200048, CAS-ITRI201902); the Science, Technology, and Innovation Commission of Shenzhen Municipality<PHONE_NUMBER>3048123); the Shenzhen Key Laboratory of Biomimetic Materials and Cellular Immunomodulation (ZDSYS20190902093409851).

**Conflicts of Interest:** The authors declare no conflict of interest. The funders had no role in the design of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript; or in the decision to publish the results.

#### **References**

## **Materials for Orthopedic Bioimplants: Modulating Degradation and Surface Modification Using Integrated Nanomaterials**

**Harbhajan Ahirwar 1, Yubin Zhou 2,3,\*, Chinmaya Mahapatra 4, Seeram Ramakrishna 5, Prasoon Kumar 6,\* and Himansu Sekhar Nanda 1,\***

Received: 24 January 2020; Accepted: 10 March 2020; Published: 12 March 2020

**Abstract:** Significant research and development in the field of biomedical implants has evoked the scope to treat a broad range of orthopedic ailments that include fracture fixation, total bone replacement, joint arthrodesis, dental screws, and others. Importantly, the success of a bioimplant depends not only upon its bulk properties, but also on its surface properties that influence its interaction with the host tissue. Various approaches of surface modification such as coating of nanomaterial have been employed to enhance antibacterial activities of a bioimplant. The modified surface facilitates directed modulation of the host cellular behavior and grafting of cell-binding peptides, extracellular matrix (ECM) proteins, and growth factors to further improve host acceptance of a bioimplant. These strategies showed promising results in orthopedics, e.g., improved bone repair and regeneration. However, the choice of materials, especially considering their degradation behavior and surface properties, plays a key role in long-term reliability and performance of bioimplants. Metallic biomaterials have evolved largely in terms of their bulk and surface properties including nano-structuring with nanomaterials to meet the requirements of new generation orthopedic bioimplants. In this review, we have discussed metals and metal alloys commonly used for manufacturing different orthopedic bioimplants and the biotic as well as abiotic factors affecting the failure and degradation of those bioimplants. The review also highlights the currently available nanomaterial-based surface modification technologies to augment the function and performance of these metallic bioimplants in a clinical setting.

**Keywords:** bioimplants; orthopedic; metallic biomaterials; degradation; surface modification; coatings; nanomaterials

## **1. Introduction**

Orthopedic bioimplants play a significant role in improving the quality of human life [1]. In this regard, bone–implant interface greatly influences bone healing through an osseointegration process [2]. The appropriate surface properties of orthopedic bioimplants include modulation of the differentiation of mesenchymal stem cells to express osteogenic phenotype [3,4]. Besides, surface modification of these bioimplants can also facilitate the biodegradation process [5,6], improve the mechanical properties commensurate with the native bone and improve integration with host tissue nearby. Furthermore, surface modification can provide antibacterial properties to avoid any post-surgery infections [7]. The above requirements of surface modifications can be adapted by metallic materials that have inherent bulk properties to be used in orthopedic applications.

Metallic bioimplants are manufactured by innovative manufacturing technologies to meet the ever-increasing demand for orthopedic applications [1]. Among the various bioimplants that have been developed for orthopedic, bioimplants used to assist bone fracture are currently in demand [8,9]. Amid the reported metallic biomaterials, materials such as stainless steel (SS 316L), titanium alloy (Ti-6Al-4V) and cobalt-chromium (Co-Cr) alloy have been investigated a lot owing to their suitable bulk properties [10]. Other non-metallic materials that have been explored in bone fracture fixation that include alumina (Al2O3), nylon 6/6, polymethyl methacrylate (PMMA), etc. [8,9,11,12]. Efforts have also been directed towards the modification of bulk properties of metallic biomaterials to render them with mechanical properties commensurate with that of a native bone, which could reduce stress shielding at an interface of tissue and bioimplant [12,13]. The excellent biocompatibility, hemocompatibility and high fatigue strength have positioned the metallic biomaterials as most suitable materials for orthopedic applications [14]. However, these bioimplants are prone to problems of wear and corrosion in a blood/tissue milieu [15,16]. The leachates of corrosion and infection labile nature of metallic surface may trigger immunological reactions [17,18], leading to the need for heavy intake of immunosuppressant or other treatments causing deleterious effects on the patient's health [19]. The major limitations of metallic biomaterials are their non-integration with host tissue owing to the above-mentioned inappropriate surface properties, vulnerability to post-surgery microbial infections, and other related risks that may demand a revision surgery and removal of bioimplants [8,12,20].

Although efforts have been directed to overcome the challenges associated with bulk and surface properties of these metallic bioimplants, nanomaterials have of late emerged as an alternative to alter the surface properties for their better integration with host tissue, reducing the immunogenicity, providing an infection-free surface and enabling the drug loading [21–26].

Surfaces of the orthopedic bioimplants serve as the site of interaction for surrounding living tissue. Hence, it is imperative to enhance the biological performance of these bioimplants using bioactive nanomaterials [27–32]. Surface engineering using nanomaterials and other suitable coating technologies aims to design and develop the bioimplants with improved osseointegration for orthopedic applications [28,33,34]. The commercialized surface treatment strategies for several of the orthopedic metallic bioimplants have already been developed; those include grit-blasting followed by washing with non-etching acid and distilled water, spark anodization in the presence of calcium phosphate, sandblasted and acid-etched (SLA) treatment for the generation of macro/micro scale topography, laser-lok technology to generate structured grooves and channels, chemical treatment for the creation of nanoscale topography, molecular impregnation of calcium phosphate on a surface, plasma treatment, and acid etching [35–42]. Some of the common commercial generated surfaces include tiunite surface on titanium implants (Nobel Biocare Implant System) [43], SLA surfaces on Roxolid™—an alloy of titanium and zirconium (Straumann Implant System) [44], Microchannelled surfaces through laser-lok technology (Bio-horizons Implant System) [45], micro textured surface created by grit-blasting on titanium (Zimmer Implant System) [46], NanoTite™ surfaces on bone implants (3i Dental Implants) [47], and others. Surface modification processes such as grit blasting are achieved by bombardment of bioimplant surfaces by means of silica, hydroxyapatite, alumina, or TiO2 nanomaterials, and thereafter the surface is treated with a non-etching acid and distilled water to clean the un-bonded nanomaterials [37,39,42]. Acid-etching treatments are generally performed using strong acids such as hydrofluoric, nitric, or sulphuric acid to create micro/nanoscale roughness on the surface of the bioimplants [35,37,38]. Several other techniques such as deposition techniques including dip coating, laser technology that creates hydrophilic and chemically active surface that promote osseoconductivity, surface patterning by micro/nanochannels/groves for cellular infiltration, thermal spraying, biomimetic deposition of calcium phosphate and hydroxyapatite for better bone integration, and sol-gel deposition have also been developed to alter the surface of the metallic bioimplants [40,48]. These techniques have already been clinically accepted and have been used as a surface treatment solution for several orthopedic bioimplants.

The current review incorporates the description of the metals and metal-alloys that are commonly used in manufacturing of different orthopedic bioimplants. The review critically discusses the biotic as well as abiotic factors responsible for the degradation and failure of metallic/metallic-alloy bioimplants. It also highlights the currently available advanced nanomaterial-based surface modification technologies to enhance the function and performance of these metallic bioimplants. To further stimulate the exchange of ideas among the experts in the field, the opinion from some of our experts is also included to make this review more interesting and appealing for future readers, expecting more practical and mature orthopedic bioimplants to be explored to improve human health.

## **2. Materials for Orthopedic Bioimplants**

Appropriate selection of materials to support fracture and healing is critical for the long-term success of orthopedic bioimplants. The choice of a bioimplant is primarily driven by the intended application, amenability to manufacturing, and the potential market size. Among the different types of biomaterials, metallic biomaterials including metals and metal alloys are widely used for manufacturing of orthopedic bioimplants because of their biocompatibility, low cost, rich in resources, and appropriate mechanical properties such as high tensile strength that provides strength to the fractured bones [1]. Some of the key metals and metal alloys used for orthopedic bioimplant manufacturing are discussed. Table 1 summarizes the different biomedical metals, their properties and the manufactured bioimplants from these metals and their applications.

**Table 1.** Biomedical metals, properties, manufactured bioimplants, and their applications.

## *2.1. Titanium (Ti) and Ti-Alloys*

Titanium (Ti) and its alloys have the characteristics of low density, high mechanical strength and excellent biocompatibility [73]. Ti in combination with other metals forms biocompatible Ti-alloys, which are widely used in bioimplant manufacturing. One of the most commonly used Ti-alloy is Ti-6Al-4V. It occupies approximately 45% of the total industrial production of Ti-based bioimplants [74]. The Young's modulus of Ti alloys is in the range of 55–110 GPa, which is higher than that of a native bone [52]. Therefore, the stress shielding effect remains an issue that can only be reduced, but currently impossible to avoid [74]. Ti and its alloys are non-toxic and inert at in vivo environment due to their corrosion resistant properties [73]. However, some studies have reported that aluminum and vanadium ions released from Ti-6Al-4V alloys can potentially damage vital organs [75]. Vanadium

ions are cytotoxic and known to cause poor osteointegration, while aluminum ions are known to cause neurological disorders [76,77]. These concerns have led to the development of Al/V-free α + β type Ti alloys with improved mechanical, tribological, and biological properties. These Ti-alloys have almost similar characteristics as that of Ti-6Al-4V. However, elastic modulus is still much higher than that of cortical bones (30 GPa). Thus, the generated stress shielding effect may further loosen the implanted bioimplants [53]. Other new generation β-type Ti-alloy includes Ti35Nb2Ta3Zr, Ti-Nb-Ta-O, Ti-Nb-Ta-Zr, Ti-35Zr-5Fe-6Mn, and Ti-33Zr-7Fe-4Cr, which have shown their respective advantages for manufacturing of orthopedic bioimplants [51,54,57–61]. β type Ti-alloys are known to consist of β-stabilizing elements such as Nb, Mn, Sn, Ta, and Zr. These elements are considered as safe for human health and hence the alloys are considered as biocompatible in nature.

To overcome the challenges of infection and support better osseointegration, the surface of Ti-alloys were modified through nanotexturing by the surface mechanical attrition treatment (SMAT) process [50,56,78]. Other methods such as coating technologies using micro-arc oxidation (MAO) have also been used to enhance the surface characteristics, biocompatibility, and osseointegration of the bioimplants manufactured from Ti-alloys, especially of β type. For example, TiO2 doped calcium-phosphate coating (Ca-P) and calcium-phosphate-strontium coating (Ca-P-Sr) were used to improve the surface properties of Ti35Nb2Ta3Zr to enhance the in vitro and in vivo performance of the bioimplant [79].

Titanium or their alloys are used to develop the tibial tray that accommodates tibial polyethylene component in prosthesis for total knee replacement (TKR), femoral stem of endoprosthesis for total hip replacement (THR), and bone screws and plates to fix fractures and plates/screws for maxillofacial applications in the cranio-facial and mandibular areas [55,80,81].

#### *2.2. Stainless Steel (SS)*

Stainless Steel (SS) is one of the most widely used metallic biomaterials in orthopedics because of their ease of manufacturing, low cost, and wide resource availability. SS contains a minimum of 10.5% chromium and varying amounts of other elements such as iron, carbon, etc. [82]. As a result of chromium addition, the surface of SS develops a thin and relatively passive metal oxide layer that protects the surface against corrosion. In addition, at least 0.03% carbon in stainless steel (SS 316L) increases its mechanical strength and maximizes the corrosion resistance properties and improves the overall tribological performance of the bioimplants [82]. The SS 316L is inexpensive, reliable, and widely used in orthopedic bioimplant manufacturing [12]. It has a lower carbon content than SS 316 (a stainless-steel grade having 0.08% carbon) and offers excellent toughness to the overall bioimplant. SS 316L exhibits relatively good biocompatibility compared to SS316 [49,82]. It has much higher elastic modulus (about 200 GPa) than that of a typical human femur cortical bone (10–30 GPa) [11]. This may result in high stress-shielding at bioimplant–tissue interface leading to the failure of the implanted bioimplant [13,82]. In addition, SS bioimplants also succumb to fatigue damage due to their low fatigue strength [62]. SS-based bioimplants either need revision surgery or to be used as a permanent bioimplant after bioactive coating or surface modification using bioactive nanomaterials. The modification of SS with bioactive hydroxyapatite (HA) improves the osseointegration and bio-integration properties of an orthopedic bioimplant [63,83,84]. Typical applications include bone plates, medullary nails, screws, pins, sutures, and steel threads used in fixation of fractures [14,85–90].

## *2.3. Cobalt (Co) Alloy*

Co alloys are wear, corrosion, and heat-resistant metallic materials used in bioimplant manufacturing [64]. In vitro and in vivo tests confirmed that these alloys as biocompatible and appropriate materials for manufacturing of surgical bioimplants such as orthopedic prostheses for the knee, shoulder, and hip as well as fracture fixation devices. Typical Co-based alloy (Co-Cr-Mo alloy) in conjunction with an ultra-high molecular weight polyethylene (UHMWPE) is used in prosthetic knees and ankles [91]. The major alloying elements include Co, Cr, Mo, and Ni. Although these

are essential trace elements in a human body, they have been well proven to be toxic when leached out in a body due to corrosion of cobalt alloys. The excessive presence of these trace elements (Co, Cr, and Mo) has been reported to damage organs such as the kidney, liver, lungs, and also blood cells [92]. The elastic modulus and ultimate tensile strength of the Co-alloys are 200–230 GPa and 430–1028 MPa, respectively, which is approximately 10 times higher than that of a human bone [49]. Hence, the bioimplants manufactured from these materials may result in stress-shielding effect at a bioimplant–tissue interface [13]. The surface modification of Co-Cr-Mo alloy implants could be achieved by low temperature plasma treatment, where surface get alloyed with nitrogen and carbon through S-phase transformation [65]. This process improves the hardness, corrosion, and wear resistant properties of Co-Cr-Mo alloys.

## *2.4. Biodegradable Metals*

Biodegradable metal-based orthopedic bioimplants eliminates the complications associated with the long-term presence of bioimplants in human body. Once these materials degrade, the degradation products can be metabolized to fulfill the elemental requirements of the metabolic pathways [20,66]. Among the different metals, magnesium (Mg) shows a great promise as a biocompatible and biodegradable material [66,69]. The attractive characteristics of Mg are its high strength, elastic modulus, and a close resemblance to the modulus of a human bone. The properties such as high mechanical strength can reduce the amount of bioimplant material needed for an applied load and hence to manufacture the bioimplant. The reduced elastic modulus of the bioimplant can prevent the modulus mismatch between a bone and Mg-based bioimplant, leading to the reduction in the stress shielding at bone–bioimplant interfaces [66]. The mechanical properties of the Mg alloys can be enhanced by alloying with aluminum and other alloying elements [68]. Current investigations are centered on identifying the new Mg-alloys with no or low cytotoxicity. Various biomedical Mg-alloys such as Mg-Y-Nd [93] and Mg-Ca [94] have been studied for the development of biodegradable Mg-alloy-based orthopedic bioimplants. Alloying metals have to be carefully selected to avoid metal-related toxicity and corrosion [20]. The major limitation of Mg and Mg-alloys is their low corrosion resistance. Low corrosion resistance results in rapid release of the degradation products due to fast in vivo degradation. These necessitate the surface modification of these materials too [71]. Mg-alloys are also being explored for tissue engineering (3-D scaffold design) for bone tissue regeneration [67,68,70,95].

### **3. Degradation of Orthopedic Bioimplants**

Degradation is one of the major considerations in bioimplant design, processing, and application. Biodegradable implants are expected to degrade progressively over a period of time to assist in the healing process and compensate for the clinical need. Bioimplants are designed either to degrade or remain inside a body rather than their removal after their function is served. Degradation of the bioimplants is desirable in several cases such as absorbable sutures, drug delivery system, and tissue engineering [3,96,97].

#### *3.1. Metallic Bioimplant Degradation: Role of Biological Factors*

Metallic materials used for the manufacturing of orthopedic bioimplants are high strength and corrosion-resistant metals and metal alloys. The prospective applications of these materials are to provide long-term mechanical support to the biological structures, while remaining inert and interacting minimally with the neighboring biological tissues [98]. However, these metallic materials undergo degradation following a time-dependent kinetics after being in contact with biological moieties for a long period of time (Figure 1a). Biological components are also chemically active, generating various ionic species during their metabolism. These chemical moieties are known to interact with the surface of metal/metal alloy-based bioimplants [99]. After initial phases of adsorption and surface oxidation, the proteins slowly interact with the implant surface in a size-dependent manner, with smaller proteins interacting being the initiator. Johnson et al. reported the degradation effects of fetal bovine serum

(FBS) on a Mg-alloy bioimplant [100]. The investigation demonstrated an oxidized Mg-Yttrium (MgY) was much more resistant to degradation in FBS compared to a native Mg-Y. The response of host body chemistry to the implant such as inflammation-dependent release of reactive oxygen species (ROS) creating an oxidative environment determines the degradation pattern of metal and metal-alloy [101].

**Figure 1.** (**a**) Graphic illustration of the degradation kinetics of a typical metal bioimplant and (**b**) cyclic representation of metallic bioimplant degradation pathway [99]. Adapted with permission from [99]; 2018 Springer.

## *3.2. Time Dependent Degradation E*ff*ects*

The degradation process starts immediately after the bioimplants are implanted inside the body. The overall degradation is a time dependent process, as illustrated in Figure 1a. The degradation pattern of a metallic bioimplant (for example, Mg-alloy-based bioimplant) undergoes an initial linear degradation that includes the oxide layer formation followed by a metal-protein-based degradation mechanism. The process is further followed by an encapsulation of a bioimplant by a fibrous tissue causing a complete isolation of a bioimplant from the surrounding tissue site. In the final step, the chronic inflammation starts clearing up the metallic implant materials via secretion of alkaline enzymes and ROS [102]. The transition from an early phase of metal-alloy chemistry to late stage of macrophages engulfing results in the erosion of metal particles, elevating the inflammation, and rapid degradation of the bulk material. Long-term studies on metallic degradation (6–24 weeks) have shown a time dependent decrease in the metal ion release in the tissue compared to the incremental degradation observed in nature. This type of degradation was also found to pass through a stabilization phase during a period of 6 to 24 weeks as shown in Figure 1a. Furthermore, it is also evident that some proteins can bypass the normal route of metal-alloy interaction and directly bind to the metal surface without the formation of an oxide layer as shown in Figure 1b. In addition, initiation of localized corrosion at the load-bearing site contributes towards the metal erosion to a greater extent than the other parts [99].

## *3.3. Degradation Mechanism*

Metal and metal alloy-based bioimplants are prone to corrosion in a tissue and blood milieu. The corrosion process is a hallmark of a degradation process. It usually starts with the redox reactions and hydrolysis at the interface involving electron or proton exchanges (Figure 2a). These reactions yield hydroxides that precipitate over the metal surface. These precipitates are known to cause foreign body reactions, further resulting in a severe inflammation and long-term fibrous tissue development. Metals such as silver (Ag) and gold (Au) are more resistant to the corrosion as compared to iron (Fe) or aluminum (Al). These precious metals are sometimes used as a coating to modulate the corrosion rate of the bulk metals used in manufacturing of the load bearing bioimplants for an orthopedic application. Iron or related materials usually precipitate as metal-protein complexes in the subsequent steps (Figure 2b). Steinemann et al. have reported that the produced metal–protein complexes are

insoluble in a body fluid and determine the stability of the hydrolyzed products [103]. In addition, metal oxidation depends on the physical properties of a metal such as crystalline or amorphous nature of a metal [104]. It has been observed that the rate of degradation is lower for the amorphous metal oxides compared to the crystalline oxides [104]. The slow degradation of any amorphous oxides is due to the higher capacity of amorphous metal oxides to form organic complexes compared to crystalline metal oxides in an aqueous medium dominated by hydroxyl ions. The observed effect could also be due to weaker metal-oxide bonding with each other owing to larger inter-atomic distances in amorphous metal oxides. This weaker bonding could be easily leveraged by the macromolecules such as proteins, giving rise to the complexes (metal–proteins) that protects the underlying metal layer [105]. In this contest, metals such as Ti form the most stable oxide layer. The metal alloy-based orthopedic bioimplants manufactured from titanium–aluminum–vanadium (TiAlV) have further shown promising resistance to corrosion and metal ion release in the presence of proteins compared to the bioimplants made from metal-alloys such as SS 316L [73]. In addition, the physical imperfections and long-term wear and tear are also responsible for the corrosion of metal implants. Santos et al. have reported the initiation of degradation in metallic screws and plates at the site of wear/tear that have undergone specific corrosion at the load bearing areas [106]. In the implants such as screws and plates that bear regular load, there is a greater tendency to accumulate maximum precipitate at the grain boundaries of these metal/metal alloys implants (Figure 2c). Deposition of a larger precipitate at the grain boundary in combination with the regular abrasive motion on the surfaces deems them much susceptible to corrosion [106,107]. Furthermore, physical imperfections have also been known to contribute towards the localized corrosion, precipitation, and aggregation of a large quantity of metal–protein complexes at the specific site, thereby weakening the load-bearing structures. The physical imperfection also leads to fretting corrosion that is a combined form of local imperfections and micro motion, observed specifically in screw plate models of orthopedic bioimplants (Figure 2d).

**Figure 2.** Mechanism of metallic bioimplant degradation (**a**,**b**) [15,16] and fretting corrosion tests on orthopedic plates and screws made of ASTM F138 stainless steel (**c**,**d**) [106]. Adapted with permission from [15]; 2008 John Wiley & Sons; Adapted with permission from [16]; 2003 Springer.

Following ionic interactions, proteins are known to get adsorbed relatively quickly within seconds to hours, even on a metal's surface with similar charges. The differences in the kinetics of adsorption of the different proteins leading to an intense competition for binding to the oxide layer on the metallic surface. Moreover, this oxide layer formation is itself influenced by the surface properties like roughness, surface energy, and others [108]. The formation of these metal oxide–protein complexes accelerates the degradation process. Several other factors in an in vivo environment such as ionic strength of the local environment, pH, and physical stress on the oxide layer further govern the degradation kinetics [108]. For example, albumin with high concentration in synovial fluid strongly binds to the metal surfaces. The presence of any defect further enhances these phenomena generating metal–protein complexes, which in turn reduce the rate of erosion of a metal surface [109].

## *3.4. Metal Self-Induced Biological Responses*

Metal–protein conjugates can act as signaling molecules for the secondary inflammation that is more intense than the primary inflammation [102]. The debris of the materials can sensitize the immune system depending upon the size of debris (nano or macro size), which induces increased cellular responses causing rapid degradation [110]. In in vivo environment, bioimplants subjected to mechanical stresses are more vulnerable to degradation. Hence, metal-alloy bioimplants that are mostly used in the high stress and load bearing conditions need to be evaluated for their long-term fatigue and friction-based degradation [111]. Oxidative stress in the tissue due to an increased ROS generation, Fenton chemistry-based corrosion in combination with fatigue-based degradation, significantly promotes the degradation rate [112].

#### **4. Surface Modification E**ff**ects**

Surface modification is usually carried out to minimize bacterial adhesion, inhibit biofilm formation, and provide effective bacterial extermination to protect the implanted biomaterials [7,28,113,114]. Furthermore, their role is to provide stealth properties to a bioimplant for any immunogenic reactions, better integration with the surrounding tissue, and enable appropriate cellular responses. For example, Ag nanoparticle coating over the Ti surface has been carried out to construct an improved bioactive and biocompatible surface (Figure 3a–c) [115]. Ti-based bioimplant and tissue integration could be promoted by surface decorated with Ag nanoparticles, through the promotion of H2S production. The production of H2S upregulates the expression of sulphur containing proteins such as albumin that is highly beneficial for the enhanced metal–protein interactions. Further the tissue could well interact with the Ti-Ag complex, resulting in an enhanced bone regeneration as shown in the Figure 3b. Controlled degradation behavior of Ti-Ag complex could cause the local release of Ag<sup>+</sup> ions, resulting in an enhanced tissue integration, anti-bacterial effect, osteoconductivity, and long-term low toxicity (Figure 3c). These effects could demonstrate multiple therapeutic effects of Ag-nanoparticle coating on a bioimplant surface. Similarly, different strategies are employed to prevent infection on a metallic bioimplant surface using antibiotics, antimicrobial peptides, inorganic antibacterial metal elements, and antibacterial polymers [116]. Certain surface modifications could also cause changes in the mechanical and biological properties of the bioimplant material. The physical/bulk modification is performed to obtain an optimum shape and size of a biomaterial with an appropriate mechanical behavior, while chemical modification is carried out to render bioactivity to the surface of a material. These alterations to materials empower them with improved cell adhesion, attachment, and eventually proliferation [115,117]. The tissue interacts predominantly with these materials at an interface where nanomaterials decorate the surface of a bioimplant. As shown in the Figure 3d, Ag-nanoparticle could act as a focal point of tissue adhesion and bone growth through interaction with calcium that regulates the biological responses. Nanomaterials are the advanced materials that could be effectively utilized to improve the surface and bulk properties of orthopedic bioimplants. A class of nanomaterials in response to electric/magnetic fields or polarization exhibit anti-bacterial properties without affecting the surface chemistry of bioimplants [118]. Nanotechnologies could improve the antibacterial response of the prosthetic bioimplants, which include compositional modification, surface chemistry alteration, as well as the application of properly tuned external stimuli [7,118]. Various desirable surface properties such as protein adsorption, osteoblast attachment, osteoblast differentiation, antibacterial activity, biocompatibility with living tissues is achieved through nanomaterials. Further, corrosion resistance of the bone–bioimplant interface has been achieved by modifying the surface of the bioimplants through

nano-structuring and functional nanocoating. A number of antibacterial agents, such as Ag, Au, zinc oxide (ZnO), zirconium nitrate (Zr(NO3)4), zirconium oxide (ZrO2), titanium oxide (TiO2), have been incorporated in a hydroxyapatite (HA) matrix to develop HA-based antibacterial coatings for orthopedic metallic bioimplants, which allows no bacterial growth on the bioimplant substrate and improves bio-integration properties [25,119–123].

**Figure 3.** (**a**) Ti-based bioimplant–tissue integration promoted by Ag nanoparticles; (**b**) H&E staining at 2, 6, and 12 weeks of implantation showing enhanced bone growth around the bioimplant. (**c**) Controlled release of Ag-nanoparticle (Ag NPs) from Ti surface causing enhanced tissue integration, anti-bacterial effect, osteoconductivity, and long-term low toxicity; (**d**) Schematic showing AgNPs acting as focal point of tissue adhesion and bone growth through interaction with calcium that regulates the biological responses [115]. Adapted with permission from [115]; 2017 American Chemical Society.

## **5. Improving the Surface Properties of Bioimplants Using Integrated Nanomaterials**

A proper design of a bioimplant material is aimed to provide durability, functional stability, and an appropriate biological response. Durability and functionality depend on the bulk properties of the material, whereas biological response depends on the surface chemistry, surface topography, and surface energy of a biomaterial. Surface modifications of bioimplants play a vital role in matching the complexities of the biological system and improving the performance of the bioimplant materials [25]. In this context, nanomaterials could be effectively utilized to improve the surface properties of several orthopedic bioimplants [26,124].

## *5.1. Surface Coating Using Ag-Based Nanocomposites*

Silver (Ag) possesses an inherent antibacterial property and low toxicity to human cells, rendering it as an appropriate antibacterial agent for biomedical applications [125]. Ag can be used in the form of ions and compounds to destroy the bacterial cells [26,121,125]. Ciobanu et al. introduced a method for synthesizing Ag-doped nanocrystalline hydroxyapatite (HA) [126] in which Ag doped nanocrystals of HA was synthesized at 100 ◦C in deionized water. The Ag-doped nano-HA materials demonstrated an excellent cell adhesion and cell proliferation resulting in the synthesis of bone-related proteins and deposition of calcium. These hybrid nanomaterials could be used as a promising candidate for the coating and the surface modification of orthopedic bioimplants. There are two major mechanisms primarily responsible for the antibacterial response of Ag in several nanomaterials [127]. Primarily, it forms Ag<sup>+</sup> ions during its oxidation, which is highly reactive with bacterial cells. These ions in the form of nanoparticles can bind to DNA, RNA, and proteins in bacterial cells to further inhibit the growth of bacteria [128]. Ag<sup>+</sup> ions are also responsible for bringing the structural changes in the

bacteria and promote cell distortion. The antibacterial activity of Ag-based nanocomposite could also be attributed to the formation of ROS, which includes free radicals such as super oxides, hydroxyl radicals, etc. These super oxides and hydroxyl-free radicals are responsible for the antibacterial response from these surface modified orthopedic metallic bioimplants [121]. In order to induce antibacterial properties to Ti bioimplants, Yang et al. utilized friction stir processing (FSP) to embed silver nanoparticles (AgNPs) in a Ti-6Al-4V (TC4) substrate. Here, silver nanoparticles placed in the preformed grooves on the surface of TC4 when subjected to FSP, they get homogenously distributed in the surface of TC4 matrix. The distribution profile of AgNPs is dependent on the depth of the preformed grooves [129,130]. On examination, it was observed that silver-rich NPs with a size ranging from 10 to 20 nm were diffused into the substrate. Thus, both FSP and the addition of silver increase the corrosion resistance and reduce the infection rate. The antibacterial effect is independent of Ag<sup>+</sup> ion release and is likely due to the number of embedded silver NPs on the surface. TC4/Ag metal matrix nanocomposite is a potential nanocomposite that embeds AgNPs on a biomaterial surface for creating balance between the antibacterial effect and biocompatibility. The modified surface possesses an antibacterial properties [129,130].

## *5.2. Surface Coating Using Nano-TiO2 and TiO2-Based Metal Nanocomposites*

Titanium oxide (TiO2) nanomaterials have an excellent biocompatibility and chemical stability for which these nanomaterials have been used as coating over the metallic bioimplants [131]. In presence of light, TiO2 oxidizes to produce free radicals (e.g., hydrogen peroxide, superoxide and hydroxyl free radicals). These free radicals have already demonstrated to elicit antibacterial responses [132]. TiO2 coating on metallic bioimplants could be activated using direct organic coating like spray coating of polymers where doped antibacterial metal ions (Ag+) are released as an "antibiotic" providing antibacterial property to a bioimplant surface. In this process, Ti-based bioimplant surface is doped with Ag<sup>+</sup> through hydrothermal treatment on polyethylene glycol (PEO), which significantly modulates the surface chemistry of the metallic bioimplant. The remodeled surface undergoes a patterned and slow degradation releasing Ag<sup>+</sup> ions in a controlled manner, leading to increase in the expression of ROS that is detrimental to the bacterial cell wall integrity (Figure 4a) [121]. The other method could be the unique redox photochemical induced mechanism that promotes enhanced bone–bioimplant integration [123]. The inorganic coating onto the bioimplant serves as a site for the redox photochemical reaction. The photochemical reaction results in the development of an electrolytic process by release of ions from the bulk material surface upon exposure to ultraviolet rays. The redox photochemical-based deposition of TiO2 demonstrated highly active surface-induced antimicrobial activity by its capacity to generate cations (Figure 4b). The later mechanism also creates an excellent corrosion resistance TiO2 layer over the bioimplant surface that protects an orthopedic bioimplant from bacterial attack as well as corrosion. TC4 surface modification is being explored by FSP to create a nanocomposite of TiO2 and TC4. It is a method where an intense, localized plastic deformation is produced on the surface of a TC4. This results in the formation of nanocrystalline and amorphous TiO2 on the surface of TC4. The presence of nanocrystalline and amorphous TiO2 improved the surface properties like surface microhardness, biocompatibility, and resistance to corrosion [133]. The FSP also resulted in the uniform incorporation of TiO2 particles to the surface of TC4 matrix. Due to the grain refinement and phase transformation, the surface microhardness and corrosion resistance properties of modified TC4 was improved. In vitro studies demonstrated an enhanced cell adhesion and proliferation capability of the TC4 substrate and modulated the biocompatibility of TC4 substrate [133,134].

**Figure 4.** (**a**) Schematic demonstrating stepwise fabrication of antibacterial responsive coating of Ag<sup>+</sup> doped TiO2 [121]; (**b**) Schematic showing the deposition and antibacterial action of TiO2 coated bioimplant via redox photochemical method [123]; (**c**) Schematic illustration of antibacterial response of ZnO-based nanomaterial coating [118]; (**d**) Schematic showing ion beam assisted coating of antibacterial Ag-CeSZ-based nanomaterial coating [120]. Adapted with permission from [118]; 2018 ACS Publications; Adapted with permission from [120]; 2019 Elsevier; Adapted with permission from [123]; 2011 SAGE Publications.

## *5.3. Surface Coating Using ZnO-Based Nanocomposite*

The surface modification using HA-ZnO nanocomposite can reduce ions leaching from a metal alloy and prevents the bacteria colonization over a bioimplant surface (Figure 4c) [118]. The experimental investigation suggested that the number of bacterial colonies could be reduced to 13% from 50.45% when ZnO content was increased from 1.5% to 30% (wt) in a HA-ZnO nanocomposite. The antimicrobial responses of ZnO-based composites are due to the formation of ROS and release of Zn2<sup>+</sup> ions as shown in Figure 4c [119]. ROS are toxic to gram-negative bacteria, while Zn2<sup>+</sup> ions are responsible for killing of gram-positive bacteria. ROS reacts with lipid layer of the cell wall (gram-negative bacteria) leading to the distortion of the bacterial cell wall. Such distortion destroys the cell wall and eventually leads to bacterial cell death. Zn2<sup>+</sup> ions diffuse inside the cell (Gram-positive bacteria) and disrupt the amino acid metabolism and enzymes, resulting in a cell death [119].

#### *5.4. Surface Coating Using Ag-CeSZ Nanocomposite*

The surface modification using silver-ceria stabilized zirconia (Ag-CeSZ)-based nanomaterials have well proven to offer better mechanical properties and fracture toughness to the bioimplant compared to a conventional yttrium stabilized zirconia [120]. Three source electron beam physical vapor deposition (EBPVD) is used for the deposition of these coatings over several orthopedic bioimplants (Figure 4d). Silver block (99.99% purity) and CeSZ sintered pellets are taken in two separate graphite crucibles and kept separately in a water-cooled copper hearth. The electron beam is then generated and controlled with an accelerating voltage of 8 kV. The filament current is varied between 30 and 60 mA using a 30 kV TT controller. In a case of Ag-CeSZ nanocomposite coating, both Ag and CeSZ are evaporated separately using two E-beam guns. The substrates are kept at

the temperature of 673 K and the thickness of the coating can be controlled in a range of 2 μm. Ti bioimplants when coated with Ag-CeSZ nanocomposite coatings show improved mechanical and biological properties. The mechanical properties of Ag-CeSZ nanocomposite coatings are due to its crystalline nature. The coating also demonstrates excellent cell adhesion, antibacterial activity and resistance to sodium fluoride (2%), showing its promising multifunctional importance in orthopedic coating technologies [135].

## *5.5. Surface Coating Using Ti*/*SiC Metal Matrix Nanocomposite*

The metal matrix composites offer increased stiffness, strength, and wear resistance over monolithic matrix materials. Nanocomposites based on silicon carbide (SiC) have exhibited enhanced mechanical properties. Recent reports on SiC have indicated that its biocompatibility is comparable to that of HA, with respect to the long-term osteogenic properties [136]. The crystalline SiC surface promotes adhesion, proliferation, and differentiation of the primary cultured osteoblast cells. Interestingly, SiC also improves the wear resistance and hardness of the bioimplant on which it is coated. During FSP, the metallic surface undergoes a plastic deformation, leading to an effective grain refinement. This ultrafine-grained metal substrate produced by the application of plastic deformation provides superior cell substrate attachment and biocompatibility. The surface nanocomposites produced by FSP exhibit excellent bonding with the underlying metallic substrate. Zhu et al. have fabricated a novel Ti/SiC metal matrix nanocomposite (MMNC) using FSP and investigated its microstructure and mechanical properties. Additionally, the proliferation and osteogenic differentiation properties of rat bone marrow stromal cells (BMSCs) on the sample surface were investigated and it was demonstrated that the modified surfaces supported the cell attachment and osseointegration [137].

#### **6. Conclusions and Future Recommendation**

Manufacturing of bioimplants often involves the integration of processes of material selection, design, and fabrication of bioimplants, and surface modifications through micro/nano texturing or nanomaterial coating. Engineering native metals by converting them into alloys amalgamate best properties of different metals in a single formulation. This provides the flexibility in tailoring the bulk properties of metals as per the orthopedic requirements. However, the surface properties of bioimplants based on alloys/metals require appropriate modification to elicit favorable biological responses. The surface modification of bioimplants through nanocomposites materials have the potential to enhance the host response in the long-run. These nanomaterials play a key role in minimizing the bacterial adhesion to further inhibit biofilm formation to protect the implanted biomaterials from microbial attack. They also play a vital role in eliciting appropriate cellular responses like cell migration through contact guidance on patterned deposition of nanomaterials, cellular differentiation, and gene expression through modulation of stiffness/hydrophobicity of the surfaces, initiation of degree of immunogenicity, delayed surface erosion, and degradation and composition of microenvironment at or in the vicinity of the bioimplant site. Advanced nanomaterials can also serve as a reservoir of drugs to be delivered at the bioimplant site. Thus, coated nanomaterials have the potential to alter the surfaces of various metallic materials for their adoption in orthopedic applications. However, care must be taken during the preparation and deposition of the nanomaterials on the surface of bioimplants like control over the size distribution of the nanomaterials, bonding of the nanomaterials with the bioimplant surface, thickness of the deposited nanomaterial, and eventually, the scalability of the process being used during nanomaterial deposition. In summary, it is imperative to say that the modulation of the degradation process and surface modification using emerging nanomaterials is going to generate a plethora of bioimplants for orthopedic applications in the near future. Therefore, the application of nanotechnology would be critical for the future success of orthopedic bioimplants.

**Author Contributions:** H.A., Y.Z., and C.M. wrote the paper. H.S.N. and P.K. formulated the work plan and designed the manuscript layout. H.S.N., P.K., and Y.Z. edited the entire manuscript and added significant discussions. S.R. provided critical comments on which H.S.N., P.K., and Y.Z. have worked and revised further. All authors have read and agreed to the published version of the manuscript.

**Funding:** The authors would like to acknowledge the funding support from the Faculty Initiation Grant to Biomedical Engineering and Technology (BET) Laboratory, Indian Institute of Information Technology Design and Manufacturing (IIITDM) Jabalpur and Guangdong Medical University Scientific Research Foundation (4SG19003Ga).

**Acknowledgments:** The authors would like to thank Mamta Anand (Assistant Professor in Humanities and Management, IIITDM Jabalpur) for assisting us in English language editing of the revised manuscript.

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

## **References**

© 2020 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 (http://creativecommons.org/licenses/by/4.0/).

## *Article* **Mechanical Properties of Strontium–Hardystonite– Gahnite Coating Formed by Atmospheric Plasma Spray**

**Duy Quang Pham 1, Christopher C. Berndt 1, Ameneh Sadeghpour 2, Hala Zreiqat 3, Peng-Yuan Wang 4,5 and Andrew S. M. Ang 1,\***

Received: 1 October 2019; Accepted: 12 November 2019; Published: 15 November 2019

**Abstract:** In this work, we measured the mechanical properties and tested the cell viability of a bioceramic coating, strontium–hardystonite–gahnite (Sr–HT–G, Sr–Ca2ZnSi2O7–ZnAl2O4), to evaluate potential use of this novel bioceramic for bone regeneration applications. The evaluation of Sr–HT–G coatings deposited via atmospheric plasma spray (APS) onto Ti–6Al–4V substrates have been contrasted to the properties of the well-known commercial standard coating of hydroxyapatite (HAp: Ca10(PO4)6(OH)2). The Sr–HT–G coating exhibited uniform distribution of hardness and elastic moduli across its cross-section; whereas the HAp coating presented large statistical variations of these distributions. The Sr–HT–G coating also revealed higher results of microhardness, nanohardness and elastic moduli than those shown for the HAp coating. The nanoscratch tests for the Sr–HT–G coating presented a low volume of material removal without high plastic deformation, while the HAp coating revealed ploughing behaviour with a large pileup of materials and plastic deformation along the scratch direction. Furthermore, nanoscanning wear tests indicated that Sr–HT–G had a lower wear volume than the HAp coating. The Sr–HT–G coating had slightly higher cell attachment density and spreading area compared to the HAp coating indicating that both coatings have good biocompatibility for bone marrow mesenchymal stem cells (BMSCs).

**Keywords:** Sr–HT–Gahnite; hydroxyapatite; atmospheric plasma spray; nanoindentation; mechanical properties

## **1. Introduction**

Substitutional materials for bone have been active research areas since the 1970s and ceramics account for the majority of these replacements [1]. These bioceramics include three groups: (i) bioinert, (ii) bioactive, and (iii) biodegradable or bioresorbable ceramics [2]. In brief, bioinert ceramics have no or very little interaction with surrounding tissues when they are placed into the host body. However, bioactive ceramics can form chemical bonds with surrounding bones and these biodegradable ceramics are gradually replaced by endogenous tissues after implantation into the host body [2]. Calcium phosphate ceramics are commonly used for bioactive applications of bone tissue repair and augmentation because of their similar chemical composition to minerals in bone [3].

Hydroxyapatite (HAp:Ca10(PO4)6(OH)2) is the most widely accepted biocompatible and bioactive calcium phosphate ceramic with respect to tissues [4]. In bulk form, HAp presents excellent physio-chemical properties as a bioceramic; however, these attributes are deteriorated by an intrinsic brittleness, low impact resistance and low tensile strength limit when implanted for load-bearing applications [5]. One method to improve the mechanical properties while maintaining the bio-adaptive properties is by forming a HAp coating onto a metallic biomaterial [6] in applications where the bioceramic is under compressive loads. Implants that are based on a HAp coated metal can provide excellent bioactivity and biocompatibility to the host bone, while retaining the good strength and ductile properties of the metallic substrate [4].

Although a HAp coating on a metallic substrate has demonstrate good clinical and commercial outcomes [7], the demand for improving the mechanical properties of a HAp coating is still imperative. For example, the high-velocity suspension flame spray method has been used to coat a nano-structured HAp/Ti composite onto 316L stainless steel (SS) substrates [8], and the suspension plasma spray (SPS) method has been employed to produce a nano-diamond reinforced hydroxyapatite composite coating onto a titanium substrate [9]. The advantages of these two methods using suspension-based feedstock are the ability to deposit materials in submicron sizes or create new composite materials for thermal spray coating [10]. However, the feedstock feeder system and torch re-configurations are required in order to adapt to the suspension feedstock, which is considered costly and complicated. Using flame spray as a coating method has been the preferred method in the last decades because it is a low cost system, portable and easy to use, but a high porous coating with oxidation and low jet temperature with low particle velocity are the major concerns of the flame spray [11]. In contrast, atmospheric plasma spray (APS) is the most commonly used coating method for orthopedic implants and has been approved by Food and Drug Administration [12]. The APS technology allows high jet temperature and velocity so it can be used to create a denser coating with the same material [11].

Introducing substitute bioceramics that may enhance mechanical properties and biocompatibility is an alternative approach addressed in this research. It is known that silicon is a trace element in the human body that occurs naturally; e.g., 100 ppm in bone and 200–600 ppm in cartilage and other connective tissues [13]. Silicon performs as a biological linkage to form the structure of connective tissues and engages in the process of biomineralization to promote bone growth [14]. The development of Si-containing biomaterials is an active area of research such as developing bioglasses and Si-doped bioceramics [15]. Calcium silicate (Ca–Si) based ceramics, in particular CaSiO3 and Ca2SiO4 are currently used in orthopaedic implants since they are bioactive materials that induce a bond with bone to form an apatite layer and promote the bone growth [16–18]. However, the main concerns of these bioceramics are low chemical stability with a high ionic dissolution and high degradation rates; This instability affects the mechanical strength and osseointegration at high pH levels [19], i.e., pH = 8.34 after soaking CaSiO3 25 days in a simulated body fluid (SBF) solution [20]. One method to control the dissolution is by incorporating metal ions into Ca-Si based bioceramics, which often uses zinc or/and strontium metals [19]. Zinc supports bone formation, increases bone protein and alkaline phosphatase activity in osteoblasts differentiation and mineralization, and enhances cell proliferation osteoconductivity [21]. Strontium is associated with bone through two main mechanisms which are (i) incorporation of Sr into bone crystal lattice–surface exchanging and (ii) ionic exchange with bone–ionic substitution [19].

Therefore, strontium doped hardystonite (Sr–Ca2ZnSi2O7), named Sr–HT, is a Ca-Si based ceramic that incorporates zinc and strontium, and which shows improvements in cellular activity and chemical stability compared with CaSiO3 [19]. Sr–HT in the form of a scaffold (pore interconnectivity: 99%, pore size: 300 to 500 μm, and porosity: 78%) has shown exceptional bioactivity as a scaffold to repair large-sized bone defects. It presented a high compressive strength of 2.16 ± 0.52 MPa, which is a favorable property for load-bearing applications. By mixing ~80 wt % Sr–HT with ~20 wt % gahnite

(ZnAl2O4) [22], the Sr–HT–G ceramic had high fracture toughness and compressive strength properties that were superior to those of cortical bone [22]. A Sr–HT–G scaffold (pore interconnectivity: 100%, pore size: 500 μm, and porosity: 85%) presented higher compressive strength of 4.1 ± 0.3 MPa than solely Sr–HT [22]. Sr–HT–G scaffolds are able to repair the large bone defects, which are outstanding against the β-tricalcium phosphate/hydroxyapatite (TCP/HA) scaffolds [22]. The Sr–HT–G scaffolds can promote the attachment of adipose derived stem cells and the alkaline phosphatase activity of these cells when contrasting with β-TCP/HA scaffolds [23]. These scaffolds also enhance the migration, proliferation and differentiation of human umbilical vein endothelial cells without adverse effect on viability cells [23]. The 3D-printed Sr–HT–G scaffolds in the study of Li et al. have shown a significant effect in repairing bone defects in sheep tibia for 3 and 12 months, where the bone formation significantly bridged the defects [24].

In this study, the atmospheric plasma spray (APS) technique was employed to deposit Sr–HT–G powders onto Ti–6Al–4V substrates in order to evaluate the possibility of using this coating for applications of orthopedic implants. Commercial hydroxyapatite powders were used to produce control samples. The study compares the mechanical and chemical properties of the Sr–HT–G to HAp coating and includes: (i) surface characteristics of the coatings, (ii) chemical analysis and phases of coatings, (iii) Vickers microhardness, (iv) distribution of hardness and elastic moduli via nanoindentation, (v) nanoscratch and nanowear behavior, and (vi) stem cell culture study.

## **2. Materials and Methods**

### *2.1. Powder Preparation*

The Sr–HT–G powders were prepared by Allegra Orthopaedics Limited (Sydney NSW, Australia) following the method that is described in the study of Roohani-Esfahani et al. [22]. Briefly, Hardystonite (Ca2ZnSi2O7) powders were prepared through the sol–gel method from tetraethyl orthosilicate ("TEOS" of formula (C2H5O)4Si), zinc nitrate hexahydrate (Zn(NO3)2·6H2O) and calcium nitrate tetrahydrate (Ca(NO3)2·4H2O). Then, strontium ions (from Sr(NO3)2, 5 wt %) were added to substitute calcium ions to form Sr–HT. Finally, alumina (Al2O3) powder (15 wt %) was added into the Sr–HT system. Initially, the Sr–HT–G powder was of submicrometer size and was not suitable for APS due to poor powder flow characteristics. The flow behavior was improved by initially mixing the Sr–HT–G powder with 5 wt % polyvinyl alcohol (PVA, Sigma Aldrich, St. Louis, MO, USA). This mixture was consolidated into bulk form by heating in an oven at 70 ◦C for 24 h. The consolidated mass was then ground in a mortar and pestle before sieving to 45–106 μm. The HAp powder, from Medicoat SAS (Etupes, France), exhibited particle sizes of 45 to 125 μm.

## *2.2. Plasma Coating Set Up and Operation*

A Metco 9MB plasma torch with a 7 mm-Metco GH nozzle (Sulzer Metco, Westbury, NY, USA) was used to deposit Sr–HT–G and HAp powders onto Ti–6Al–4V substrates. The coatings were deposited at room temperature, using a six-axis robot (YR-SK16-J00; Motoman Robotics, Miamisburg, OH, USA) to control the torch position with the velocity of between 100 to 150 mm/s. The same coating parameters were applied for both Sr–HT–G and HAp powders (Table 1).

**Table 1.** Setup parameters for plasma coating of Sr–HT–G and HAp powders.

### *2.3. Coating Characterisations*

Phase compositions of powders and coatings are analysed by the Bruker D8 Advance XRD system (Bruker Corp., Billerica, MA, USA) with Cu Kα radiation. The XRD system operates at 40 kV and 30 mA to scan over the range of 20◦–70◦ with a step size of 0.08◦ and 2 s of dwell time. The XRD results were indexed using Diffracplus EVA software (EVA V5, Bruker Corp., Billerica, MA, USA) and then compared with database PDF2018-PDF-2-Release 2018 RDB from the International Center for Diffraction Data (ICDD, Newtown Square, PA, USA). The crystallinity level is calculated based on the intensity of peaks, which is described in Equation (1), where *A*c and *A*a are total intensity of peaks and total intensity of amorphous phase [25].

$$\text{Crystallimity} = A\_{\text{c}} (A\_{\text{c}} + A\_{\text{a}}) \times 100\text{\textdegree } \tag{1}$$

where *A*c and *A*a are total intensity of peaks and total intensity of amorphous phase.

Chemical analysis of coatings was conducted using an X-ray photoelectron spectrometer (XPS) with a Kratos Axis Nova system (Kratos Analytical, Manchester, UK) and Raman microspectroscopy with an In-Via Raman Microscope system (Renishaw®, Gloucestershire, UK). The XPS system has the X-ray source of Al Kα with hν = 1486.69 eV that operates at 150 W. XPS data were analysed by CasaXPS software (Version 2.3.15, CASA Software Ltd., Cheshire, UK) after a Shirley background subtraction with the correction of binding energy for carbon at 285 eV. The XPS characterisation was done on the as-sprayed samples, which is how it would be used in actual orthopedic applications. Raman microspectroscopy used 785 nm laser wavelength with 10% of the 150 W laser power and 10 s for exposure time. High-resolution electron micrographs were obtained using a field emission SEM (SUPR 40 VP, Carl Zeiss, Oberkochen, Germany) at 5 kV. In all cases, the coating surfaces were cleaned by ethanol and then dried by compress air. After that, the samples were place in a vacuum environment till it underwent characterisation.

Coatings were cut and prepared following ASTM E1290-03 "Standard Guide for Metallographic Preparation of Thermal Sprayed Coatings" [26] to get cross sections of coatings for analysis. A Vickers microhardness tester (Micromet 2103 Microhardness tester, BUEHLER, Lake Bluff, IL, USA) was setup to perform the microhardness test on cross sections of coatings at the load of 300 gf and dwell time of 15 s. Forty readings was recorded on each sample and the Weibull distribution analysis was performed to analyse coating properties and reliability. Surface roughness, *R*a, was measured using a portable surface roughness tester SJ-210 (Mitutoyo, Kawasaki, Japan) with a 2 μm stylus diameter together with a 3D optical profiler system (Bruker Corporation, Billerica, MA, USA).

A nanoindenter (Hysitron TI Premier, Bruker, Eden Prairie, MN, USA) with a Berkovich diamond tip was used to perform the nanohardness and elastic moduli. A cono-spherical 1 μm diameter diamond tip was used in the nanoindentation system to performe nanoscratch and nanowear tests. The resolutions of the nanoindentation system are 1.0 nN and 0.006 nm for the applied load and tip displacement, respectively. The distributions of hardness and elastic moduli of coatings were investigated using accelerated property mapping (XPM) with the support of the scanning probe microscopy (SPM) method. The nanoindentation test were performed following the testing procedure, which was described in the previous study of Pham et al. [27]. Briefly, the tests were done on areas of 30 <sup>×</sup> 30 <sup>μ</sup>m<sup>2</sup> with 400 indents in each area (1.5 <sup>μ</sup>m gap between indents). Fifteen areas was selected along the cross section of each coating, and the applied loads were from 500 to 3000 μN. Elastic moduli and hardness were determined by measuring penetration depths after indenting at the applied load using the Oliver–Pharr method [28].

Nanoscratch tests were performed at the loads from 1000 to 2000 μN with the scratch length of 10 <sup>μ</sup>m. Nanowear test was conducted by applying a load of 200 <sup>μ</sup>N onto the area of 10 <sup>×</sup> 10 <sup>μ</sup>m2 and scanned over the section three times for each selected areas at the scanned frequency of 0.8 Hz. The wear volumes were calculated by measuring the different heights between the before and after of the scanned areas, and multiplying the height difference with the scanning size. The results from

nanoscratch and nanowear tests were analysed by TriboView software (Version <IP_ADDRESS>, Hysitron TI Premier, Bruker, Eden Prairie, MN, USA).

## *2.4. Stem Cell Culture*

Human bone marrow mesenchymal stem cells (BMSCs) were purchased from Cyagen Biosciences (HUXMF-0101, Suzhou Inc., Suzhou, China) at passage 2 and maintained in a complete growth media (HUXMF-90011, Cyagen Bioscience). The complete growth media of BMSC was composed of 440 mL basal medium, 50 mL fetal bovine serum, 5 mL glutamine, 5 mL penicillin-streptomycin. It was used as purchased without the addition of extra growth factors. The growth media is only used for BMSC maintenance. For the biocompatibility test, cells were seeded at a density 5000 cells/cm<sup>2</sup> and fixed after 24 h. Phalloidin-Rhodamine B (cat# P1961, Sigma, Tokyo, Japan) and DAPI (cat# D1306, Sigma) were employed for F-actin filament and cell nucleus staining, respectively. Immunostained samples were washed thoroughly and examined under an inverted fluorescence microscope (IX73 Olympus, Tokyo, Japan) at 20 magnification. Cell density (cells/mm2) and cell spreading area (μm2) on three different substrates including as-received Ti, HAp coating, and Sr–HT–G coating, which were quantitatively analysed based on the obtained images. The cell number and cell spreading area were analysed from the DAPI stained and F-actin stained cells via NIH ImageJ software (Version 1.52). DAPI stained images were processed via threshold color and the number of nuclei was automatically determined via the plugin of cell counter in Image J software (Version 1.52); Similarly, F-actin stained images were processed via threshold color and the occupied area was automatically measured via Image J software (Version 1.52).

## **3. Results and Discussions**

## *3.1. Feedstock Morphology*

Morphologies of Sr–HT–G (after of mixing with PVA, consolidating, grinding and sieving) and HAp powders are indicated in Figure 1a,b, respectively. Sr–HT–G powders were manually ground and sieved so the powders are described as angular and blocky, while HAp powders are spherical-like shape. The morphology of Sr–HT–G powders provide higher ratio of surface area to volume than that of HAp powders, which can lead to a better degree of heating in plasma plumes [29]. HAp may present better flow transportation characteristics than Sr–HT–G because spherical powders are preferable morphology in plasma spray. These HAp particles can easily enter into the center of the plasma plume to be well deposited, while Sr–HT–G powders may not be injected ideally into the plasma flame due to irregular sizes [11,29].

**Figure 1.** Feedstock morphology of strontium–hardystonite–gahnite (Sr–HT–G) (**a**) and hydroxyapatite (HAp) (**b**).

#### *3.2. Coating Surface Analysis*

#### 3.2.1. Phase Compositions

Figure 2 presents the XRD results of powders and coatings for Sr–HT–G and HAp. Sr–HT–Ga powders consist of two distinct phases in Figure 2a, which are Sr–HT (PDF 01-072-1603) and gahnite (PDF 01-070-8181). The gahnite phase in starting powders is formed by the reaction between zinc in Sr–HT and alumina during the sintering process at high temperatures to produce powders [22]. Similar to the staring powders, the coating of Sr–HT–G also presents two phases, which are Sr–HT and gahnite. The peaks of Sr–HT–G coating are significantly less sharp and are broaden after the plasma coating, which indicate the reduction in the level of crystallinity. However, there is no new peaks formed in the Sr–HT–G coating, but they are slightly shifted in the peaks of the coating comparing with its powders. The reason for the peak shifts could be the internal or residual stress after coating, which causes the change in lattice parameters. On the other hand, HAp powders present high purity with phase of HAp, but its coating presented the phase of Ca10(PO4)6(OH)2 (PDF 00-064-0738) and a new phase of tricalcium phosphate-TCP (Ca3(PO4)2) (PDF 01-072-7587) in Figure 2b. This new phase TCP was formed by the reaction of HAp in the high temperature region of plasma plume [30].

**Figure 2.** XRD results of Sr–HT–G powder and coating (**a**) and HAp powder and coating (**b**).

Both coatings illustrate lower intensity levels or lower crystallinity compared with their starting powders. Significant lower intensity of the peaks in a coating comparing to its starting powder can be seen in the result of Sr–HT–G, of which the level of crystallinity from Equation (1) are 74.6% and 32.7% for powders and the coating, respectively. On the other hand, levels of crystallinity in HAp powders and coating are not significantly different, which are 84.8% and 75.5% for powders and the coating, respectively. The lower crystallinity in the coating is a feature of the plasma process, which feedstock are exposed and heated by a 10,000 ◦C plasma flame, then rapidly cooled at about 10<sup>6</sup> degrees per second, where the feedstock do not have enough time to recrystallize [29]. It can be seen that the Sr–HT–G coating presents lower levels of crystallinity than that in the HAp coating. Sr–HT–G is a Ca-Si ceramic with the sorosilicate (Si2O7) group that has a more intricate crystalline structure, so the process of recrystallize will be more complex than that of HAp [31]. In contrast, HAp powders with larger particle sizes are well deposited with a larger quantity in the coating because they are fed into the center of the plasma flame. As the result, a greater amount of bulk crystalline can be found in the coating, which leads to a higher percentage of crystallinity [32].

#### 3.2.2. XPS on Coating Surfaces

Figure 3a presents the XPS analysis result of Sr–HT–G coating with the presence of Sr, Ca, Zn, Si, O, Al on the coating surface. Two peaks of Zn in the spectrum are Zn 2*p* and Zn 2*p*1/2 at 1022.9 and 1045.9 eV, respectively. Furthermore, the presenting of the peaks Al 2*p* at 74.0 eV and Al 2*s* at

119.0 eV together with the peak O 1*s* at 531.5 eV indicates the formation of ZnAl2O4 on the coating surface [33,34]. In a high-resolution scan presented in Figure 3c, Si presented the peak of Si 2*p* at 102 eV, which refers to the presence of a silicate group [35]. This silicate group is Si2O7, which is in the structure of Sr–HT. Formation of Sr–HT and gahnite phases are further supported by results from the analysis in XRD. The Sr–HT and gahnite composite ceramic has shown an improvement in biocompatibility and bioactivity, which has been evaluated by its capability in osteoconduction and osteoinduction [36]. Thus, a plasma coating with the present of Sr–HT and gahnite phases will potentially produce good biocompatibility and bioactivity.

Chemical elements on the surface of HAp coating under XPS analysis are shown in Figure 3b. The surface of HAp coating consists of Ca, P, and O as the main elements in its starting powders. The peak of P 2*p* is 132.8 eV, while the peak of O 1*s* is 530.8 eV. This result indicates the present of the PO4 <sup>3</sup><sup>−</sup> group [35], which is in the structure of hydroxyapatite Ca10(PO4)6(OH)2. Calcium has four peaks, which are Ca 2*s*, Ca 2*p*, Ca 3*s*, and Ca 3*p*. The Ca 2*s* presents a peak at 438.6 eV, while the Ca 2*p* consists of a doublet with values of 346.9 and 350.4 eV for Ca 2*p*3/2 and Ca 2*p*1/2, respectively. The peaks of Ca 3*i* and Ca 3*p* are 44.0 and 25.3 eV that are illustrated in Figure 3d. The results of calcium peaks indicate the formation of Ca3(PO4)2 [37]. These HAp phases have been shown to support the in vitro and in vivo formation of bone.

**Figure 3.** XPS analysis of Sr–HT–G coating (**a**) and HAp coating (**b**). The peaks of Si 2*p* (**c**), Ca 3*i* and Ca 3*p* (**d**).

#### 3.2.3. Morphologies of Coating Surfaces

Surfaces of the Sr–HT–G and the HAp coatings are shown in Figure 4a,b, respectively, where both coatings show well-melted splats. However, the Sr–HT–G coating shows more uniform surface morphology with spherical-like splat shapes, while HAp coating presents irregular splats with a big splash at the rims. Furthermore, some fragmentations can be found within splats of the HAp coating, whilst the Sr–HT–G coating shows more consistent structure on the surface with more uniform splats. Under the melting and re-solidify processes in APS, well-melted splats or partially melted particles will spread onto the substrate and form layers of splats [38]. Molten particles of Sr–HT–G show higher spreading ability than the HAp, which could be closely related to its viscosity characteristics and eventually reveals more regular Sr–HT–G splats [39].

**Figure 4.** SEM of coating surfaces of Sr–HT–G (**a**) and HAp (**b**).

## *3.3. Coating Cross Section Analysis*

## 3.3.1. Coating Microstructures

Figure 5 indicates the microstructure of the Sr–HT–G and the HAp coatings at cross sections. Both coatings present features of thermal spray coating including cracks, pores, and unmelted particles [38]. Different phases can be found in the HAp coating as the result of different contrasts, which are shown in bright gray and dark gray. The amorphous phase is easily removed during the polishing process, which causes the height to be lower than the crystalline areas. As the result, the amorphous phase appears to be darker. This observation has been re-confirmed with Raman spectroscopy. Similar results have been presented in the works of Gross et al. [30,40]. However, there is not much difference in contrast of phases in the Sr–HT–G coating as its microstructure presents a similar contrast to the source of ignition during the SEM process. It can be seen that the Sr–HT–G coating has less cracks in microstructure, while the HAp coating presents more cracks. Moreover, due to the reaction of alumina with zinc in Sr–HT, a more uniform structure is formed [36], where the Sr–HT–G coating has a denser structure and a more consistent microstructure compared with the HAp coating. Less pores in the structure of the Sr–HT–G coating could come from the larger distribution of powder sizes, where different sizes of Sr–HT–G splats are likely to fill up the pores in the microstructure during the process of deposition [11]. In addition, the interface of coating-substrate in the Sr–HT–G coating is quite consistent with crack-free features, while the HAp coating shows minor cracks at the coating-substrate interface.

**Figure 5.** Microstructure of Sr-HT-G (**a**) and HAp (**b**) on coating cross section.

## 3.3.2. Vickers Microhardness

The results of Vickers microhardness are shown in Figure 6, where the hardness of the Sr–HT–G and the HAp coatings are 330.4 ± 54.4 HV300 and 120.2 ± 24.3 HV300, respectively. Vickers microhardness in the coating of Sr–HT–G is significantly high due to the addition of Zn into the Ca-Si structure [19,21], which is almost triple the hardness of the HAp coating. In addition, the formation of a glass phase between Sr–HT grains and the existence of a strong restraint surrounding the Sr–HT grains from the glass-ceramic phase of gahnite also leads to the improvement in mechanical properties of the Sr–HT–G coating [36]. The coefficient of variant (CoV = standard deviation divided by the mean) of the Sr–HT–G is coating slightly lower than that of the HAp coating, which are 16% and 20%, respectively. The coating of Sr–HT–G has lower CoV values in microhardness because the Sr–HT–G coating possesses a more uniform microstructure compared to those in the HAp coating. The Weibull analysis of microhardness of the Sr–HT–G and HAp coatings are presented in Figure 6, where both coatings present good regression fits with *R*<sup>2</sup> values of 0.97 and 0.96, respectively [41]. The Weibull moduli m, which presents the scattering behavior of results within the distribution [41], show that both the Sr–HT–G and HAp coatings have low values of Weibull modulus but they are the typical values in thermal spray. The Sr–HT–G coating has a slightly higher value of Weibull moduli than the HAp coating, which indicates less variation in the microhardness results in the Sr–HT–G coating.

**Figure 6.** Weibull regression analysis of the Vickers hardness test of HAp and Sr–HT–G coatings.

#### 3.3.3. Nanohardness and Elastic Moduli

The cross section of the Sr–HT–G coating shows a uniform structure in the coating, shown in Figure 7a. Figure 7b,d presents corresponding the SPM and SEM images of the same area after performing the nanoindentation. The indentation mapping results of the Sr–HT–G coating are presented by the distributions of elastic moduli and hardness in Figure 7c,e, respectively. The variant in the distributions of nanohardness and elastic moduli depends on the features of the selected area such as element compositions, phase structure, cracks, and pores. Generally, lower hardness and lower elastic moduli can be found in pore and crack areas [42]. Uniform distributions of nanohardness and elastic modulus can be achieved in the Sr–HT–G coating, which indicate the consistent of mechanical properties across the coating. This uniformity could be achieved by the presence of the gahnite phase in the coating that could enhance the mechanical properties across the coating [36].

**Figure 7.** Results of the nanoindentation test of the Sr–HT–G coating. (**a**) The SPM image of surface before applying the indents; (**b**) and (**d**) the SPM and SEM image of surface after applying the indents; (**c**) and (**e**) property map of reduced modulus and hardness of the coating.

The distributions of nanohardness and elastic moduli in a selected area from the HAp coating are found in Figure 8. Similar to the testing procedure of the Sr–HT–G coating, Figure 8a presents a selected area that was pre-scanned to capture the features before performing the nanoindentation. After performing the nanoindentation, the surface was captured again via SPM and SEM techniques, which are shown in Figure 8b,d, respectively. The difference in contrast of the SEM image present different phases that have formed, which are named as "A" and "B" regions. The results of elastic moduli and nanohardness of the HAp coating are shown in Figure 8c,e in the form of distribution maps, respectively. The marks of indents in the HAp coating at the "A" area present deeper and larger sizes than those in the "B" area. Raman microspectroscopy was used to analyze the functional groups of the highlighted Region A and B of the HAp coating. It can be seen that both regions have same functional groups as the result of similar positions of Raman peaks in Figure 9. However, sharper peaks with higher intensity can be found in Region B, which indicates the formation of crystalline HAp while Region A presents the formation of an amorphous phase [43]. Similar to the study of J. Wen et al. [44], Region B yielded better mechanical properties, which it showed higher results of nanohardness and elastic moduli than Region A.

This indicates there is less uniformity of coating phases in the HAp coating when comparing with the Sr–HT–G coating. As the result, both nanohardness and elastic moduli showed high variants in distribution of values across the HAp coating. The average values of elastic moduli and nanohardness of the Sr–HT–G coating were 93.0 ± 16.6 and 7.2 ± 2.1 GPa, respectively. On the other hand, lower results of both elastic modulus and nanohardness can be found in the HAp coating, which the average values were 77.6 ± 41.5 and 4.4 ± 3.2 GPa, respectively. The average results of both nanohardness and elastic moduli in the Sr–HT–G coating showed significantly better results than that of the HAp coating, which were in accordance with the results in the Vickers microhardness tests. Furthermore, the Sr–HT–G coating presented lower values of CoV in both hardness and elastic moduli with 17.8% and 29.2%, respectively, while these values of hardness and elastic moduli in the HAp coating were higher at 53.4% and 72.7% due to the lack of uniform phases.

**Figure 8.** Results of the nanoindentation test of the HAp coating. (**a**) The SPM image of surface before applying the indents; (**b**) and (**d**) the SPM and SEM image of surface after applying the indents; (**c**) and (**e**) property map of reduced modulus and hardness of the coating.

**Figure 9.** Raman microspectroscopy of the regions in the HAp coating.

#### 3.3.4. Nanoscratch and Nanoscanning Wear Performance

Results of nanoscratch tests under the load of 1000 μN are presented in Figure 10, where Figure 10a,d shows the surfaces of the selected areas after the scratching of Sr–HT–G and HAp coatings, respectively. It can be seen that the Sr–HT–G coating depicts tearing behavior only with shallow depths of penetration, while deeper and more plastic flow across the edge of the scratch direction can be found in the HAp coating. The 3D maps in Figure 10b,e render more details of the behavior of these coatings under scratch tests. The HAp coating experienced more plastic deformation with large pileup of materials that represents ploughing behavior under the scratch test.

**Figure 10.** Nanoscratch test results of the Sr–HT–G coating (**a**,**b**); and the HAp coating (**d**,**e**); friction coefficient (**c**); and normal displacement (**f**).

Figure 10c,f reveals the coefficient of friction (CoF) and depth of the scratch (displacement) of the Sr–HT–G and the HAp coatings, respectively, which show the details of Scratch 1, 2, and 3 from the Figure 10a,c. In the nanoscratch test, the actual applied load started from the 27th second to the 42nd second, while other times for loading and unloading periods are not considered. From the results of the nanoscratch tests, CoFs and displacements of the Sr–HT–G coating show high consistency with low values of CoFs and low penetration depths in all the scratches. On the other hand, CoF values and displacements in the HAp coating fluctuate continuously, i.e., values of CoF and displacement significantly decreased from 29.5 to 35 s in the scratches due to higher scratch resistance in this region of the HAp coating. This result could have originated from the significant difference in mechanical properties of different phases in the HAp coating, which has a similar trend with the distribution of hardness and elastic moduli.

The average depths of scratch are 60.9 ± 9.7 nm (CoV = 16%) and 245.0 ± 150.4 nm (CoV = 61%) for the Sr–HT–G and the HAp coatings, respectively. Lower scratch depth with less removal of materials in the Sr–HT–G coating indicate that the Sr–HT–G coating is more resistant to scratches. Furthermore, the average CoF of the Sr–HT–G coating is 0.14 ± 0.02 (CoV = 14%), which is much lower than the average CoF in the HAp coating, 0.69 ± 0.14 (CoV = 20%). Similar to the analysis in the nanoindentation, the lower result of CoV in the scratch test shows the coating microstructure of the Sr–HT–G coating presents more uniform features than the HAp coating. This results in the Sr–HT–G coating shows better scratch resistance than the HAp coating.

The wear behavior of the Sr–HT–G and the HAp coating under the applied load of 1000 μN are presented in Figure 11, in which the pre-wear test, and post-wear test images were evaluated. The black square areas indicate the regions that were selected to perform nanowear tests. The surface before performing wear tests are presented in Figure 11a,d,g; while Figure 11b,e,h and Figure 11c,f,i illustrate the surface after the wear test of SPM images and their 3D views, respectively. Similar to the result from nanoindentation and nanoscratch tests, the Sr–HT–G coating presents uniform distribution with consistent results of wear volume in Figure 11b,c. On the other hand, the HAp coating reveals more variance of wear behavior due to a significant difference in the region of "A" and "B" as amorphous and crystalline regions, respectively, in Figure 11e,h. It can be seen that amorphous areas present less wear resistance than crystalline areas, showing similar trends in nanoindentation and nanoscratch tests. The wear volume for Sr–HT–G and HAp coatings are 0.11 <sup>±</sup> 0.03 and 5.52 <sup>±</sup> 2.90 <sup>μ</sup>m3, respectively. This enhanced mechanical properties of the Sr–HT–G coating was achieved by doping zinc into a Ca-Si

ceramic together with the formation of gahnite in the structure [19,36]. This result indicates that the Sr–HT–G coating possesses superior wear resistance than the HAp coating.

**Figure 11.** Nanowear test results of the Sr–HT–G coating (**a**–**c**) and the HAp coating (**d**–**i**).

## *3.4. Results of Stem Cell Culture*

After cell culture for 24 h, BMSC adhered on all the substrates indicating a good biocompatibility, which can be seen in Figure 12a. Cell density on Ti substrates was relatively higher than samples of the Sr–HT–G coating and the HAp coating in Figure 12b. In addition, BMSCs have a relatively larger cell spread on Ti substrates than that of the Sr–HT–G and HAp coatings, which can be seen in Figure 12c. This strong cell attachment on the Ti substrate is related to the differences in surface roughness profiles; both coatings had similar and higher surface roughness than the Ti. The surface roughness (*R*a) of the Sr–HT–G coating and the HAp coating were 13.7 ± 0.76 and 10.2 ± 0.7 μm, respectively; while the roughness of the as-receive Ti was about 0.1 μm. Nonetheless, the results indicate both coatings have good biocompatibility for BMSCs, in which the Sr–HT–G coating presented slightly higher cell density and spreading area of cell attachment compared to the HAp coating.

**Figure 12.** Biocompatibility tests of different substrates using bone marrow stem cells. (**a**) BMSC morphology on Ti, Sr–HT–G, and HAp surfaces. Cell nucleus (blue) and F-actin (red) were stained. (**b**) Cell density and (**c**) spreading area on various substrates. \*\* *p* < 0.01 and \*\*\* *p* < 0.001 compared to Sr–HT–G and HAp surfaces.

#### **4. Conclusions**

The deposition of Sr–HT–G powders onto Ti–6Al–4V as a novel coating via an atmospheric plasma spray technique has demonstrated significant improvement in mechanical properties compared with the current commercial relevant hydroxyapatite coating. The Sr–HT–G coating showed more uniform microstructures with consistent distributions of hardness, elastic moduli, compared to the HAp coating. The as sprayed Sr–HT–G coating presented phases of Sr-HT and gahnite without formation of new phases. Vickers microhardness of the Sr–HT–G coating was significantly higher than results from the HAp coating; 330.4 ± 54.4 and 120.2 ± 24.3 HV300, respectively. Nanoindentation tests revealed consistent distributions of nanohardness and elastic moduli in the Sr–HT–G coating with the average values of 93.0 ± 16.6 and 7.2 ± 2.1 GPa for hardness and elastic moduli, respectively. On the contrary, the HAp coating showed lower and high variation in the distribution of hardness and elastic moduli, especially between the distinct areas of crystalline and amorphous phases. The nanohardness and elastic moduli of the HAp coating were 77.6 ± 41.5 and 4.4 ± 3.2 GPa, respectively. The Sr–HT–G coating also showed better resistance to scratch and wear compared to the HAp coating under the same tests. In addition, the Sr–HT–G coating presented slightly higher density and spreading area of cell attachment in contrast to the HAp coating, indicating that the Sr–HT–G coatings possess good biocompatibility with good attachment of bone marrow mesenchymal stem cells. This study indicates that atmospheric plasma sprayed Sr–HT–G can be a viable approach for orthopaedic implants.

**Author Contributions:** Conceptualization, A.S., H.Z., C.C.B. and A.S.M.A.; methodology, D.Q.P., A.S.M.A. and C.C.B.; software, D.Q.P. and Y.-P.W.; validation, C.C.B., A.S.M.A. and H.Z.; formal analysis, D.Q.P.; investigation, D.Q.P.; resources, A.S., D.Q.P. and A.S.M.A.; data curation, D.Q.P., A.S.M.A.; writing—original draft preparation, D.Q.P., P.-Y.W.; writing—review and editing, D.Q.P., A.S.M.A., C.C.B. and P.-Y.W.; visualization, D.Q.P., A.S.M.A.; supervision, A.S.M.A., C.C.B., H.Z.; project administration, A.S.M.A., C.C.B.; funding acquisition, A.S., H.Z. and C.C.B.

**Funding:** This research is funded by the Centre for Innovative BioEngineering under the Industrial Transformation Training Centre (ITTC) scheme via the Australian Research Council (ARC) Award IC170100022. PYW thanks the support from the National Key Research and Development Program of China (2018YFC1105201), the National Natural Science Foundation of China, and the International cooperative research project of Shenzhen collaborative innovation program<PHONE_NUMBER>3048123).

**Acknowledgments:** The authors also would like to thank Deming Zhu for his assistance in operating the XPS and Andrew Moore for expert assistance in the preparation of samples. We deeply appreciate Ping Du from SIAT who performed the biocompatibility test.

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

## **References**

© 2019 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 (http://creativecommons.org/licenses/by/4.0/).

*Article*

## **Facile Route of Fabricating Long-Term Microbicidal Silver Nanoparticle Clusters against Shiga Toxin-Producing** *Escherichia coli* **O157:H7 and** *Candida auris*

**Sheeana Gangadoo 1, Aaron Elbourne 1, Alexander E. Medvedev 2, Daniel Cozzolino 1, Yen B. Truong 3, Russell J. Crawford 1, Peng-Yuan Wang 4, Vi Khanh Truong 1,5,\* and James Chapman 1,\***

Received: 4 November 2019; Accepted: 17 December 2019; Published: 1 January 2020

**Abstract:** Microbial contamination remains a significant issue for many industrial, commercial, and medical applications. For instance, microbial surface contamination is detrimental to numerous aspects of food production, infection transfer, and even marine applications. As such, intense scientific interest has focused on improving the antimicrobial properties of surface coatings via both chemical and physical routes. However, there is a lack of synthetic coatings that possess long-term microbiocidal performance. In this study, silver nanoparticle cluster coatings were developed on copper surfaces via an ion-exchange and reduction reaction, followed by a silanization step. The durability of the microbiocidal activity for these develped surfaces was tested against pathogenic bacterial and fungal species, specifically *Escherichia coli* O157:H7 and *Candida auris*, over periods of 1- and 7-days. It was observed that more than 90% of *E. coli* and *C. auris* were found to be non-viable following the extended exposure times. This facile material fabrication presents as a new surface design for the production of durable microbicidal coatings which can be applied to numerous applications.

**Keywords:** copper nanoparticles; silver nanoparticles; nanostructure; nanocluster; antifungal; antibacterial; escherichia coli; candida auris

## **1. Introduction**

The surface colonization of bacteria and fungi on abiotic substrates is commonly referred to as a biofilm formation and significantly contributes to healthcare and industrial concerns [1,2]. This issue is further impacted by the current rise in antibiotic resistance amongst microbial species, which has caused a significant increase in persistent infections and related deaths [3]. Recent economic projections have estimated that bacterial infections could be responsible for approximately 10 million deaths per annum by 2050 if new antimicrobial surface therapies are not developed [4,5].

In particular, *Escherichia coli (E. coli)* and *Candida auris (C. auris)* are common pathogenic microbes responsible for recent outbreaks [6,7]. Shiga toxin-producing *E. coli* (STEC) O157 has emerged as

a public health threat following its initial identification as a pathogen in 1982, which was initiated when an outbreak of illness was associated with the consumption of undercooked ground beef [8]. A 2018 outbreak of the *E. coli* strain O157:H7 was reported across 36 states, infecting 210 people [6]. *C. auris* is another emerging pathogen currently responsible for invasive disease in healthcare facilities around the world [7,9] and the fungal infection carries an astonishingly high mortality rate of 60% among infected patients, representing a significant healthcare issue [7,9]. This high mortality rate has occurred due to the simultaneous emergence of multidrug-resistant *C. auris* isolates across three separate continents [9], raising pressing concerns regarding the identification and detection of invasive candidiasis isolates [10].

Silver nanoparticles (Ag NPs) are highly effective microbiocidal agents against both bacteria and fungi [11]. Ag NPs have been used in forms including solutions, thin-film coatings, or embedding in polymers [12–14]. While there have been some reported cases in which bacteria and fungi were found to be resistant against Ag NPs [15,16], recent reviews have shown alternative ways in eradicating bacterial cells via mechanical rupturing with nanostructure modifications [13,17–19]. There is a lack of significant research within the ability of these nanostructures to rupture fungal cells, in particular with the combination of Ag NPs and nanostructures not implemented as treatment strategies towards both bacterial and fungal cell surfaces.

In this work, a facile route was used to fabricate Ag NPs coatings on copper (Cu) surfaces, in which Ag NPs would assemble into high-aspect-ratio clusters. These surfaces were assessed for their long-term microbiocidal activity against both Shiga toxin-producing *E. coli* and *C. auris* cells. The fabricated surfaces present a new direction in the design of durable microbicidal surface coatings.

## **2. Materials and Methods**

### *2.1. Fabrication of Hydrophobized Ag NP Coatings on Cu Surfaces*

Cu surfaces were cut into 2 <sup>×</sup> 2 cm2 squares, pre-cleaned, and washed with HCl (Merck Pty Ltd, VIC, Australia) and MilliQ water (Synergy®UV Millipak Express 20, Merck Pty Ltd, VIC, Australia). The surfaces were then submerged with 0.1 M AgNO3 for 5 min, where Ag NPs formed on the Cu surface via a series of reduction–oxidation reactions. Surfaces were again washed carefully with MilliQ water, then with 100% ethanol (Merck Pty Ltd, VIC, Australia), and with MilliQ water as a last wash, and were blown dry with compressed nitrogen. This study adopted the fabrication method used in the previous work [20]. The Cu surfaces were further submerged with 0.1 M dichloromethyl silane (Sigma–Aldrich, Castle. Hill, NSW, Australia) in dichloromethane (Sigma–Aldrich Pty Ltd, NSW, Australia) for 15 min and were then washed with ethanol and MilliQ water to remove all chemical residues. In this study, pristine 'copper surfaces' are noted as 'Cu surfaces' and 'silver nanoparticles embedding on copper surfaces' were noted as 'Ag NPs-Cu surfaces'.ImageJ version 1.8.0 (https://imagej.nih.gov/ij/) was implemented to analyze SEM images. Color Threshold and Analyze Particles Plugins were used to determine the dimensions of non-spherical Ag NPs. A total of 90 data points per system were used over three SEM micrographs to determine the distribution of NP clusters.

## *2.2. Bacterial Strains, Growth Conditions, and Sample Preparation*

The antimicrobial efficacy of the surfaces was investigated using the strains *E. coli* O157:H7 and *C. auris,* which were obtained from the American Type Culture Collection and SA Pathology Lab, respectively, and were chosen as representatives of two major emerging pathogen outbreaks. The *E. coli* bacterial cultures were grown on Luria–Bertani (LB) agar (BD Difco, VIC, Australia) and fungal *C. auris* cultures were grown on Potato Dextrose Agar (PDA) (Sigma-Aldrich Pty Ltd, NSW, Australia) overnight at 37 ◦C. Bacterial and fungal cells were collected at the logarithmic stage of growth. To determine similar numbers of cells, despite variations in cell densities following collection, the density of the bacterial and fungal suspensions was adjusted to OD 600 = 0.1 at the logarithmic stage of cell growth. To quantify cell numbers in the adjusted bacterial suspensions before attachment

experiments, a hemocytometer was used as suggested previously [21]. Pristine Cu and Ag NPs-Cu surfaces were cut into squares of 0.5 <sup>×</sup> 0.5 cm2. The surfaces were pre-sterilized with 70% ethanol, dried within a sterilized laminar flow cabinet for 24 h and placed in a sterilized 24-well plates (Thermo Fischer Scientific Australia Pty Ltd, VIC, Australia). Bacterial and fungal suspensions were prepared as above with OD600 = 0.1 and 1 ml of each suspension was added into 24-well plates containing the sterilized surfaces. The plates were incubated at 25 ◦C in dark conditions, to avoid any effects of light on cell viability. The samples were incubated in static conditions for 1 and 7 days without any disturbance from media addition or exchange. This procedure was adopted to avoid any effects of fresh media and shaking on antimicrobial assays. Two technical replicates were done.

## *2.3. Confocal Laser Scanning Microscopy (CLSM)*

The surfaces were washed with 10 mM PBS (pH = 7.4) prior to the CLSM imaging conducted using ZEISS LSM 880 Airyscan upright microscope (Zeiss, Oberkochen, BW, Germany). To assess their viability, adhered cells were stained using a LIVE/DEAD® BacLightTM Viability Kit (including SYTO® 9 and propidium iodide) (Molecular ProbesTM, Invitrogen, Grand Island, NY, USA). In this kit, SYTO® 9 binds to nucleic acids in both intact and damaged cells, while propidium iodide (PI) predominantly enters cells with a damaged membrane considered non-viable. The bacterial and fungal cells on surfaces were stained according to the manufacturer's protocol [22]. The proportions of live and dead cells on pristine Cu and Ag NPs-Cu surfaces were then evaluated using a ZEISS LSM 880 Airyscan upright microscope (Zeiss, Oberkochen, BW, Germany). To determine the percentage of non-viable cells, CellC Cell Counting Software (https://sites.google.com/site/cellcsoftware/) was used as previously instructed [19]. An analysis was done over 5 representative micrographs over 2 technical replicates. A student *t*-test (Microsoft Excel) was conducted to compare the antimicrobial performance of these samples.

## *2.4. SEM Characterization*

Prior to the SEM imaging, bacterial and fungal cells on Cu and Ag NPs-Cu surfaces were fixed using 3% glutaraldehyde and were dehydrated using a series of ethanol concentrations (30%, 50%, 70%, 80%, 90%, 100%). The samples were removed and dried under the laminar flow of a Biosafety Cabinet Class II for 2–4 h. Fixed cells were coated with a thin film of gold prior to imaging. SEM images were taken using a field-emission scanning electron microscope (FE-SEM) (FEI Verios, FEI company, OR, United States) at 5 kV, where imaging of the systems uses methods that have been previously described [23].

## *2.5. Focussed Ion Beam-Scanning Electron Microscopy (FIB-SEM) Characterization*

Surfaces were affixed and dehydrated using the identical dehydration steps described above. A cross-sectional analysis was carried out to show the interaction of *E. coli* and *C. auris* with the Ag NP cluster substrates was performed using a FEI Scios Dualbeam FIB-SEM. The cells were first coated with platinum (Pt) using an ion gun at 16 kV/0.15 nA to prevent further damage, followed by a sequential ion slicing at 16 kV/0.15 nA. The imaging was carried out using a standard secondary electron (ETD) and upper in-lens column (T2) detectors at 2 kV/0.1 nA.

## **3. Results**

## *3.1. Fabrication and Characterisation of Hydrophobized Ag NP-Coated Cu Surfaces*

Ag NP clusters were fabricated directly onto Cu surfaces via a single-step electrochemical synthesis. A further surface silanization step was added to create a hydrophobic Ag NP cluster on the Cu surfaces. The fabricated Ag NPs-Cu surfaces were then characterized using SEM and energy dispersive X-ray (EDX) spectroscopy. In Figure 1A,B, the SEM micrographs show that the Ag NPs exhibit sizes that are consistently less than 100 nm in diameter. These Ag NPs were found to form clusters, with wide

sizes ranging from 70 nm to 1200 nm and an average size of ~350 nm per cluster. The aspect-ratio of the clusters was estimated using the ratio of length (nm) to width (nm). This measurement revealed that the clusters possess an aspect-ratio (length/width) ranging from 1 to 14. With the addition of silane to the clusters, the surfaces exhibit water contact angles of ~120◦, meaning that the surfaces are hydrophobic in nature. Silane coatings are well-known corrosion-inhibitors [24] and therefore provide a strong potential in preventing the corrosion of Ag–Cu surfaces. In Figure 1C, chemical mapping using Energy Dispersive X-ray Analysis (EDX) revealed that the Ag NPs are formed via replacing Cu(0) using a series of electroless galvanic reactions (see Equations (1) and (2)) [25,26]:

$$\text{Red: } -2\text{Ag}^+ + 2\text{e}^- \to 2\text{Ag}^0 \tag{1}$$

$$\text{Ox: Cu} \rightarrow \text{Cu}^{2+} + 2\text{e}^- \tag{2}$$

Galvanic replacement reactions are commonly known to produce Ag NPs via this route [27,28]. Importantly, low level of oxidation was detected on the Ag NPs-Cu surfaces. Therefore we provide a simple method to fabricate Ag NP clusters which have demonstrated several advantages: simplicity, facility, and a low fabrication cost. This galvanic replacement reaction method can be used on an industrial scale, which highlights the utility of this synthetic route [29].

**Figure 1.** Surface characterization of pristine Cu and Ag nanoparticle coatings (Ag NPs-Cu). (**A**) SEM images showing the surface topography of pristine Cu (left) and Ag NPs-Cu (right) surfaces (Scale bar 20 μm, inset 3 μm). (**B**) High resolution SEM image showing Ag NPs cluster domains (scale bar 4 μm). The inset shows the cluster size distribution. (**C**) EDX spectroscopy showing the distribution of Cu, Ag, and O elements across the Ag NPs-Cu surfaces (scale bar 1 μm).

#### *3.2. Microbicidal Performance of Ag NPs-Cu Surfaces over 1- and 7-day Incubations*