**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 (20180921173048123); 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 10.0.0.2, 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 (20180921173048123).

**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*

Two representative species of major global disease outbreaks in recent years, Shiga toxin-producing *E. coli* O157:H7 and *C. auris*, were selected to study the relative material and surface interactions, with the use of glass surfaces employed as control substrates (Figures S1 and S2). The data in this study reveal the native cell morphology and viability of the respective bacterial and fungal pathogens incubated on both pristine and Ag NPs-Cu surfaces for periods of 1 and 7 days. To assess their viability, LIVE/DEAD Fluorescent Kits (Molecular ProbesTM, Invitrogen, Grand Island, NY, USA) were used to stain the viable and non-viable cells [22]. SYTO® 9 dye enter the viable cells and propidium iodide would enter the non-viable cells with compromised cell membranes. Control glass surfaces demonstrated over 90% cell viability for both *E. coli* and *C. auris* over a 7-day incubation (Figure S2) while pristine Cu surfaces showed slight killing activity against the cells as shown in Figure 1, despite the inherent antimicrobial activity native of Cu surfaces. *E. coli* specifically formed a well-established biofilm on the pristine Cu surfaces. In comparison, Ag NPs-Cu surfaces inhibited the growth of *E. coli* and *C. auris* (Figure 2), with 75% and 98% of non-viability observed respectively, after a 1-day incubation only. Greater results were obtained after a 7-day incubation on the Ag NPs-Cu surfaces with more than 90% and ~100% non-viability observed for *E. coli* and *C. auris* respectively. Further investigation of the interactions between bacteria/fungi and Ag NP clusters was carried out using FIB-SEM and clear deformations of *E. coli* and *C. auris* were examined and shown by tearing of the cell membrane and presence of multiple hole features in the cell morphology, as shown in Figure 3A,B. Past studies have reported similar examinations [18,30].

The antimicrobial behavior of Ag NPs is known to occur via several mechanisms: (1) physical damage by direct contact [31], (2) the release of silver ions [32,33], and (3) reactive oxygen species production [34,35]. The breadth of information concerning the antibacterial mechanism of Ag NPs is beyond the scope of this article; however, the interested reader is directed towards several important reviews in the field [11,34]. In previous work, Ag NPs have been reported as effective antimicrobial agents against both bacteria (*E. coli* and *Staphylococcus aureus*) and fungi (e.g., *Candida albicans*) [11,12,16,36], but past investigations have failed in showing the long-term effect of silver coatings against both bacteria and fungi over a 7-day incubation, as shown in this study. Furthermore, this study is the first showing the effects of antifungal coating towards *C. auris*. There are a number of studies demonstrating that high-aspect-ratio nanostructures can be used to capture bacteria and disrupt bacterialmembranes [18,37–40]. Surface nanostructures have been considered to be "promising methods" to stop bacterial adhesion and proliferation; still, there are a number of questions that should be answered before being applied in the actual environments [41]. However, the synergy between both surface nanostructure and chemistry should be considered to be the good candidate in the development of antimicrobial approaches.

**Figure 2.** Assessment of Shiga-toxin-producing *E. coli* O157:H7 and *C. auris* on copper and Ag NPs-Cu surfaces over 1- and 7-day incubations. CSLM micrographs showing the viability of (**A**) *E. coli* and (**B**) *C. auris* on pristine Cu and Ag NPs-Cu surfaces (green indicating viable cells; and red indicating non-viable cells). CLSM images are 150 μm × 150 μm. (**C**) Quantification of cell viability on Cu and Ag NPs-Cu (\* indicating *p* < 0.05 comparing with Cu surfaces on day 1 and 7, respectively, *n* = 10).

**Figure 3.** Morphology of *E. coli* and *C. auris* on control and Ag NPs-Cu surfaces (Scale bar 5 μm). (**A**) Top-view SEM micrographs showing the cell deformation under the Ag NP clusters (scale bar 5 μm). (**B**) SEM cross-sections using FIB-SEM reveal the interfacial interaction between cell surfaces and nanostructures (left scale bar 1 μm; right scale bar 4 μm).

## **4. Conclusions**

Both *E. coli* and *C. auris* have been reported to be the main cause of recent outbreaks of disease. These two microbes were chosen to investigate the antimicrobial activity of durable Ag NP cluster coated Cu surfaces. Microbiocidal composite Ag NP cluster–Cu coatings were fabricated via an ion-exchange reduction reaction. The durability of microbicidal efficacy of the surfaces was established against pathogenic fungal and bacterial species, specifically *E. coli* O157:H7 and *C. auris*, over exposure periods of 1 and 7 days. It was observed that more than ~90% of *E. coli* and ~100% *C. auris* were non-viable, after only 7 days of surface immersion. The surfaces reported in this study provide a facile fabrication route and a new design parameter to produce durable microbicidal coatings for numerous applications.

**Supplementary Materials:** The following are available online at http://www.mdpi.com/2079-6412/10/1/28/s1, Figure S1: Scanning electron micrographs of *E. coli* and *C. auris* on glass surfaces (scale bar 2 μm). The morphology of *E. coli* and *C. auris* were found to be intact and no damage on glass substrates, Figure S2: Confocal scanning laser microscopic images of *E. coli* and *C. auris* on glass surfaces (scale bar 10 μm). Most of the cells were found to be viable on glass substrates.

**Author Contributions:** For research articles with several authors, a short paragraph specifying their individual contributions must be provided. The following statements should be used "conceptualization, V.K.T. and J.C.; methodology, S.G., A.E., A.E.M., D.C., Y.B.T., R.J.C., P.W., V.K.T., J.C.; formal analysis, S.G., A.E.M., V.K.T., J.C.; writing—original draft preparation, S.G., A.E., V.K.T., J.C.; writing—review and editing, S.G., A.E., A.E.M., D.C., Y.B.T., R.J.C., P.W., V.K.T., J.C.; supervision, V.K.T., J.C. All authors have read and agreed to the published version of the manuscript.

**Funding:** The work is supported by the School of Science, RMIT University.

**Acknowledgments:** The authors thank both the Microscopy and Microanalysis Facility (RMMF) and the MicroNano Research Facility (MNRF) at RMIT University for the use of their facilities.

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

## **Vapor-Stripping and Encapsulating to Construct Particles with Time-Controlled Asymmetry and Anisotropy**

## **Ting-Ying Wu 1,**†**, Chendi Gao 2,**†**, Man-Chen Huang 3, Zhi Zhang 2, Peng-Yuan Wang 4, Hsun-Yi Chen 3,\*, Guosong Chen 2,\* and Hsien-Yeh Chen 1,5,6,\***


Received: 13 November 2020; Accepted: 14 December 2020; Published: 18 December 2020

**Abstract:** An innovative chemical vapor sublimation and deposition (CVSD) process was shown to produce nanoscale anisotropic hybrid materials. Taking advantage of controlled thermodynamic properties and the mass transfer of molecules, this process allowed for water vapor sublimation from an iced template/substrate and stagewise vapor deposition of poly-*p*-xylylene onto the sublimating ice substrate. In this study, the use of sensitive soybean agglutinin (SBA) protein tubes was demonstrated as an example to prepare the anisotropic hybrid material based on the CVSD process. The rationale of a timing parameter, Δ*t*, was controlled to program the sublimation of the SBA-ice templates and the deposition of poly-*p*-xylylene during the CVSD process. As a result of this control, a stripping stage occurred, during which SBA tubes were exposed on the particle surface, and a subsequent encapsulation stage enabled the transformation of the ice templates into a nanometer-sized anisotropic hybrid material of poly-*p*-xylylene as the matrix with encapsulated SBA tubes. The timing parameter Δ*t* and the controlled stripping and encapsulating stages during CVSD represent a straightforward and intriguing mechanism stemming from physical chemistry fundamentals for the fabrication of hybrid materials from sensitive molecules and with predetermined sizes and asymmetrical shapes. A simulation analysis showed consistency with the experimental results and controllability of the timing mechanism with predictable particle sizes.

**Keywords:** vapor sublimation; vapor deposition; nanoparticle; anisotropic material; timed control

## **1. Introduction**

Anisotropic micro- and nano-objects enable dissimilar and multiple properties in physical, chemical, and biological aspects, thereby offering unrivaled synergistic multifunctional properties compared to the simple functions usually found in conventional isotropic materials. These advanced

anisotropic materials have been shown to be useful in a wide range of applications as energy materials, optical materials, and biomaterials [1–3]. Many review articles have outlined recent progress in the fabrication processes and the promising applications of these anisotropy materials [4–9].

Exciting combinations used to access anisotropy across the building blocks of distinct materials have generated complexity with regard to a new dimension of combined physical (geometrical) properties and chemical (material) functionalities. Although successfully predicted [2], there are challenges regarding different thermodynamic complexities, such as intermolecular interactions between diverse molecules in composited systems [10,11], unfavorable and phase-separated boundaries, compromised functionality due to structural disorders, and limited experimental methods for the fabrication of sophisticated and sensitive anisotropic materials. Currently, existing methods for producing these materials include colloidal assembly [12], the application of postmachinal force to reshape structures [13,14], the application of regional modifications to create directional patches [15,16], and electrified jetting with combined jetting solutions (cojetting) [17,18]. Specific limitations, however, include restricted sizes and low yields, problematic intermolecular interactions, e.g., chemical reactions or physical bindings, causing various forms of aggregations, and uncontrollable diffusion or phase separation, resulting in undesired cross-boundary issues in each component composed of sensitive structures, which are subjected to irreversible deconstruction and denaturation during the fabrication process. These problems still limit the construction of anisotropic materials composed of sensitive structures and delicate functionalities, and novel approaches to produce these anisotropic materials easily with high yields are still in great demand. Chemical vapor sublimation and deposition (CVSD) is a unique and versatile technique to construct porous materials with sizes ranging from centimeters to nanometers [19,20]. This technique has been successfully employed for vapor-depositing a poly-*p*-xylylene polymer system on an iced template substrate. Compared to the stationary substrates used in conventional vapor deposition, the CVSD deposition occurred on a dynamic sublimation substrate, where shrinkage in the volume and perimeter of the sublimating substrate caused a mobile surface in the direction of shrinkage. The resulting poly-*p*-xylylene deposition on this mobile substrate enabled not only the planar coverage of poly-*p*-xylylene on the substrate surface, but also allowed for deposition into the third depth dimension due to the mobile substrate, resulting in a three-dimensional bulk material of poly-*p*-xylylene. Whereas past studies have exploited a refined CVSD approach to form nanometer-sized objects [19,21], we found that CVSD on a mixed solution of a solid template containing nonvolatile components can result in the construction of an anisotropic material final product (Figure 1). Two stages were proposed to fabricate the anisotropic object by CVSD and were controlled by a programmable time (Δ*t*) parameter to regulate the sublimation and deposition in the stripping stage and the encapsulation stage during the CVSD process, i.e., an iced template was subjected to sublimation to evaporate water molecules, and the same template was subjected to vapor-deposited poly-*p*-xylylene to transform the template into a poly-*p*-xylylene matrix for encapsulation of the SBA protein tubes. CVSD offers the same advantages as the conventional chemical vapor deposition process, i.e., no solvent, dry process, and conformal deposition with respect to the substrate topology and geometry [22–24]. Additionally, it provides (1) controlled unsteady-state mass transport by using a sublimation substrate, avoiding problematic intermolecular and phase separation issues, (2) a time-dependent control parameter to independently regulate the mass transport of each component in the sublimation phase and the deposition phase during CVSD, enabling the construction of the final anisotropic object with tunable bulk size and/or geometric shape, and (3) the predetermined distribution of the solute phases (one or multiple), eliminating stresses from interphase and intermolecular interactions due to changes in concentrations and rehydration, which frequently occur in conventional solution-based systems [25,26] and freeze-dried products [27,28], enabling the preservation of the delicate function and structure of sensitive molecules.

**Figure 1.** Illustration of the CVSD fabrication process to construct anisotropic nanomaterials. A time-dependent (Δ*t*) control parameter was used to control both the stripping stage and encapsulating stage during CVSD to tune the final hybrid products with a determined size and dimensional configuration.

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

#### *2.1. Preparation Process of SBA Protein Tubes*

First, 4-(2-Hydroxyethyl)-1-piperazineethanesulfonic acid (HEPES) buffer was prepared by dissolving 0.952 g HEPES and 0.468 g NaCl into 200 mL pure water ([HEPES] = 20 mM and [NaCl] = 40 mM). The pH of the buffer was adjusted by adding 1 mM NaOH solution until a value of 7.2 was achieved. The SBA protein, CaCl2, MnCl2 and the ligands were dissolved in buffer separately and kept at 4 ◦C for one day. All solutions were filtered through a 0.22 μm membrane before use. The SBA tube sample was generated by simply mixing these solutions and adjusting the final concentration of the sample to [SBA] = 0.1 mM, [ligand] = 0.1 mM, [CaCl2] = 5 mM, [MnCl2] = 5 mM), [HEPES] = 20 mM and [NaCl] = 40 mM. The formation of these SBA protein assemblies was induced by dual-supramolecular interactions: carbohydrate−protein interactions occurred first between the protein and N-acetyl-α-D-galactosamine on the ligand, followed by rhodamine dimerization between two ligands. The protein tubes were formed in 48 h, and their length gradually increased with increasing time. These protein tubes remained stable for more than one year at 4 ◦C in a mild neutral pH = 7.2 buffer solution.

## *2.2. Fabrication Process*

The synthesized SBA protein tube solution with a concentration of 0.1 mmol/L was used to prepare the ice-particle templates. Droplets of the solution were formed by spraying the solution onto a hydrophobic poly(tetrafluoroethylene) (PTEF) substrate. A solidification procedure using a liquid nitrogen bath was performed to transform the solution droplets into iced particles. These ice particles served as templates for the subsequent CVSD process. The ice particles were placed in a homemade vapor deposition chamber [23,24] for the sublimation and deposition processes. The processing conditions were a pressure of 150 mTorr and at a deposition temperature of 4 ◦C. For the vapor deposition of poly-*p*-xylylene, polymerization from a precursor dichloro-[2,2]-paracyclophane (Galxyl C, Galentis, Marcon, Italy) was performed by vaporizing the precursor at approximately 120 ◦C, followed by a higher temperature of 650 ◦C to pyrolyze the vapor precursor into quinodimethane radicals. The deposition and polymerization finally occurred in a cooled sample holder at 4 ◦C. Argon carrier gas at a flow rate of 15 sccm was used during the process to deliver the vapor precursors

and the radicals. A deposition rate of 0.5 Å/s was monitored by in situ quartz crystal microbalance (QCM) equipment (STM-100/MF, Sycon Instruments, East Syracuse, New York, NY, USA) mounted on the deposition chamber.

## *2.3. Characterizations*

The vapor compositions during the CVSD fabrication process were characterized by real-time mass spectrometry with a residual gas analyzer (RGA, HAL RC 511, Hiden Analytical, Warrington, UK). The operation conditions were a pressure of 10−<sup>7</sup> mbar, an electron ionization energy of 70 eV and an ionization emission current of 20 μA. The mass detection range was from 0 amu to 400 amu. SEM images were recorded with a NovaTM NanoSEM (FEI, Hillsboro, OR, USA) at a primary voltage of 10 eV and emission current of 204 μA with an Everhart-Thornley detector (ETD). Negative TEM images were obtained by staining the samples with 1 wt% uranyl acetate for 10 s. The TEM experiments were then performed with an H-7650 TEM (Hitachi, Tokyo, Japan) at 25 kV, and the images were captured by a Veleta TEM camera (Olympus, Tokyo, Japan). Cryo-TEM samples were obtained using Vitrobot (Thermo Fisher, Waltham, MA, USA) then transferred to Fischione 2550 Cryo transfer holder (Fischione, Export, PA, USA) and images were captured under 200 kV by a Tecnai TF20 TEM (FEI, Hillsboro, OR, USA). AFM experiments were acquired by a nanoscope IIIa (Veeco, Edina, MN, USA) in tapping mode using silicon nitride tips (Bruker, Billerica, MA, USA) with a tip radius of 12 nm and a spring constant of 0.04 N/m. FT-IR spectra of the fabricated particle samples were recorded by a Spectrum 100 spectrometer (PerkinElmer, Waltham, MA, USA) equipped with an advanced grazing angle specular reflectance accessory (AGA, PIKE Technologies, Fitchburg, WI, USA) equipped with a liquid nitrogen-cooled MCT detector. The recorded spectra ranged from 600 to 4500 cm−<sup>1</sup> with 16 scan times at 4 cm−<sup>1</sup> resolution. An IX71 fluorescence microscope (Olympus, Center Valley, PA, USA) equipped with a Lumen200 200 W fluorescence lamp (Prior, Cambridge, UK) was used to detect the ligand within the SBA proteins, and a VK-9500 3D profile microscope (Keyence, Osaka, Japan) was also used to analyze the morphology and height of the fabricated particle samples.

#### *2.4. Simulations*

A simulation of the sublimation of an ice particle was performed and constructed with COMSOL Multiphysics, and two-dimensional (2D) and three-dimensional (3D) finite element analysis were applied. The deformed geometry method was employed to trace the sublimation interface. The model framework was constructed based on the assumption of homogeneous and isotropic composition of the ice particle with a symmetric and spherical geometry. The velocity of the deformed interface was derived based on mass and heat balance and was given by:

$$N\Delta H\_{sub} = V\_S \rho\_{ice} \Delta H\_{sub} = Q\_S \tag{1}$$

where *N* is the mass flux of water vapor, Δ*Hsub* is the latent heat of sublimation of ice, which is determined by the Clapeyron equation [29] since we assumed that the frozen ice phase is in equilibrium with the water vapor at the ice-air interface, *VS* is the interface velocity, ρ*ice* is the density of ice, and *QS* is the heat flux at the interface. The interface velocity was determined by the Stefan condition [30] and arbitrary Lagrangian–Eulerian (ALE) formulation [31]. The saturation vapor pressure at the boundary was obtained from the real-time mass detector, which is measured to be 5 <sup>×</sup> <sup>10</sup>−<sup>8</sup> torr. The general and reasonable diffusion coefficient assumed a constant dependence on temperature [32]. The bulk and interface temperatures were room temperature (20 ◦C) and at a phase change temperature of water at 150 mTorr, respectively.

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

Delicate soybean agglutinin (SBA) protein assemblies were fabricated by the ligand strategy following previously reported procedures, and were used for the generation of the proposed anisotropic

object due to their unique structure, which is symmetric and tubular in three dimensions. Briefly, SBA assemblies were constructed utilizing the specific sugar-binding capability of its four monomers in the presence of Mn2<sup>+</sup> and Ca2<sup>+</sup> [33]. Through the assistance of a developed ligand composed of N-acetyl-α-D-galactosamine, ethylene oxide spacer and rhodamine B, a self-assembled SBA tetramer structure with a helical tube with a diameter of approximately 20 nm and a length of 100–2000 nm was generated. The synthesized SBA was characterized with dynamic light scattering (DLS) and Cryo-TEM (cryogenic transmission electron microscopy), and the results showed an SBA tube approximately 250 nm in size and 26 nm in diameter (Figure 2a,b). Suspension solutions of the synthesized SBA protein tubes with a defined composition were then prepared to form droplets on hydrophobic surfaces [21] and transformed into solidified protein solution particles by a temperature solidification process, e.g., liquid nitrogen or dry ice bath, resulting in a sublimation template for the CVSD process used in this study. The thermodynamic conditions (approximately 0.2 mbar and 20 ◦C) of the CVSD operation favored the transformation of the solid-phase water component (ice) of the solidified protein solution into water vapor, similar to the generation of CO2 vapor from dry ice in ambient conditions [34] or to the freeze-drying process to evaporate the water (or solvent) phase from a solution mixture used in many purification and processing operations [35,36]. Under the CVSD conditions, water molecules were evaporated by sublimation from the solid iced particles, whereas the nonsublimating protein molecules retained their controlled distribution on the substrate and were seamlessly encapsulated into the deposited polymer structure [20,21].

**Figure 2.** (**a**) DLS and (**b**) Cryo-TEM data showed that the fabricated SBA tubes were approximately 234.3 ± 26.3 nm in length and 26 nm in diameter. (**c**) Droplets of protein solution were formed on a hydrophobic surface. A measured water contact angle value of 137.3 ± 6.5 degrees was shown in the inset. (**d**) Mass spectrometric analysis of vapor composition during the CVSD process. Characteristic peaks at approximately 18 amu (water molecule), 104 amu (quinodimethane), and 139 amu (chloro-quinodimethane) were detected.

The construction of the proposed anisotropic objects was therefore established by CVSD with a controllable two-stage process: (i) stripping stage, (ii) encapsulating stage, where the time-dependent parameter, Δ*t*, was used to control the sublimation degree of the sublimated ice template, and the elapsed time was propositional to the sublimated volume of shrinkage. More specifically, in the stripping stage, the protein molecules were used as a nonsublimation-solute system in the ice particles with the sublimating water molecules (solvent system). The sublimation of water molecules resulted in a smaller volume and receding interface perimeters of the ice particles, where the protein molecules were located towards the outer particle surfaces, while in the encapsulating stage, the vapor deposition of poly-*p*-xylylene occurred stagewise (after the same programmed time), and the remaining volume (mass) of the sublimating ice particles was transformed into porous structures of poly-*p*-xylylene, which was used a matrix for the encapsulation of the protein molecules without affecting their configuration and structure. Compared to the stripping stage, during which mass transport flux only occurred in one (outward) direction through sublimation, the encapsulating stage additionally involved the mass flux of deposited poly-*p*-xylylene into the system onto the sublimating template/substrate, which led to a transformation through the replacement of vaporized water molecules with the deposited poly-*p*-xylylene molecules. The stagewise stripping and encapsulating processes finally resulted in the proposed anisotropic nanoparticle product, which was composed of poly-*p*-xylylene as the matrix and distributed SBA proteins within the poly-*p*-xylylene structure. In the experiments, symmetrical sphere ice particles of the described protein solution were prepared by forming droplets on hydrophobic surfaces (water contact angle approximately 137.3 ± 6.5 degrees, Figure 2c) followed by a solidification process using a liquid nitrogen bath. The ice particles served as sublimation templates, and the operation time (Δ*t*1, Δ*t*2, Δ*t*3, ... ) was controlled and resulted in emerging and eventually exposed nonvolatile components of the protein solutes by the time-dependent shrinkage of the ice perimeter due to sublimation in the two-stage process of the fabrication process. The controllable Δ*t* parameter regulated the timing of the sequences (i) and (ii) during the fabrication process and determined the aspect ratio and shape of the finalized (hybrid-) nanoparticle products. A real-time mass spectroscopic gas analyzer was used to monitor the vapor composition (Figure 2d), and the correspondent species including water vapor (18 amu), carrier gas of argon (40 amu), derivatives of quinodimethanes (104 amu and 139 amu) were detected in the studied vapor system of the CVSD fabrication process. A more systematic analysis of the changes in mass was further performed to verify the controllability of time during the stripping and encapsulating stages in the CVSD process. As revealed in Figure 3, the detection of an increased level of water vapor (18 amu) indicated the occurrence of sublimation in the stripping stage, and similarly, a vaporized chloro-quinodimethane precursor (139 amu) was detected to indicate the initiation of deposition in the encapsulation stage. The timing parameter, with values of Δ*t*<sup>1</sup> = 60 s, Δ*t*<sup>2</sup> = 120 s, and Δ*t*<sup>3</sup> = 185 s, was shown to control and program both the stripping stage and the encapsulating stage, resulting in changes in the vapor compositions at the selected and designated times. From a chemical point of view, characterizations for the resultant hybrid particles products were also performed by using FT-IR analysis. As shown in Figure 4, the recorded spectra showed the characteristic –C–Cl and –C–H bands from poly-*p*-xylylene, similar to other reported poly-*p*-xylylene systems. For the hybrid product comprised both SBA and poly-*p*-xylylene, despite possible overlapped peaks of –C=O and –C–O with the –C–Cl peaks, additional –N–H and –O–H bands from solely SBA were detected and confirmed the successful fabrication of the hybrid materials of poly-*p*-xylylene/SBA.

**Figure 3.** Real-time mass spectrometry analysis of the vapor compositions during the CVSD process showed controllable timing of Δ*t*1, Δ*t*2, Δ*t*3, with corresponding regulations of vapor compositions in the stripping stage and the encapsulating stage.

**Figure 4.** FT-IR spectra showed the chemical composition of the fabricated anisotropic composite particles with poly-*p*-xylylene as the matrix and the encapsulated SBA proteins. Spectra were also recorded for pure poly-*p*-xylylene and pure SBA for comparison.

*Coatings* **2020**, *10*, 1248

A simulation analysis was then performed to understand the timing mechanism during the CVSD and to rationalize the controllability of the geometrical properties vs. the Δ*t* parameters. Taking advantage of the similarity of a freeze-drying process with the ice-template sublimation process of the current system during the proposed stripping stage, a reported mathematical freeze-drying model was used in the simulations [37]. Our model framework was designed according to the proposed dimensional scale of an ice particle system, and the simulated dimensional and temporal information was collected. Briefly, a spherical ice particle was initiated and coupled through heat and mass transfer mechanisms in its sublimation process, and the moving sublimation boundary was treated as a sharp interface. The interfacial temperature was obtained by the saturation vapor pressure, while the interface velocity was determined by the normal heat flux difference at the interface. The processing pressure of the CVSD process was used as the pressure of the external environment of the ice sphere. Although the SBA tubes were described as nonsublimating solutes within the ice template, they were sparsely distributed and nonreactive during the CVSD process. Therefore, the SBA proteins should have a limited effect on the sublimation process in the simulation. The simulation results were consistent with the experimental results, for example, a Δ*t* = 60 s reduced the size of an ice particle from the initial 50 μm to 32 μm, 160 s resulted in a size decrease to 5 μm, 177 s to 500 nm, and 178 s to 250 nm, as shown in Figure 5a. After the stripping stage, with the determined dimension configurations according to the set Δ*t*, the subsequent encapsulating stage initiated the transformation by inducing mass transfer of poly-*p*-xylylene through vapor deposition (in contrast to retrograding mass transfer by sublimation) onto the sublimating ice particle. The final asymmetrical nanoparticles were obtained by replacement of the remaining ice/water molecules with the aforementioned polymerized poly-*p*-xylylene molecules. As revealed in Figure 5b, combinations of characterization techniques, including optical microscopy, scanning electron microscopy (SEM), transmission electron microscopy (TEM), and AFM, showed that the encapsulation and poly-*p*-xylylene deposition occurred by retrograding the outer surface of the ice particles, thereby exposing the protein to the particle surface, and a time-dependent configuration of the size and the exposed protein-assembly dimensions was found, which was consistent with the simulation results. Particles with 38.3 ± 8.6 μm size were fabricated when Δ*t* = 60 s, 6.3 ± 1.6 μm when Δ*t* = 160 s, 608.3 ± 56.8 nm when Δ*t* = 177 s, and 294.3 ± 38.6 nm when Δ*t* = 178 s; a first-order linear correlation of <sup>Δ</sup>*<sup>t</sup>* vs. length (√<sup>3</sup> Δ*V)* was found, which supported the simulation data and the reported results [19]. The fabricated anisotropic particle structure was preserved without compromising the delicate SBA tube structure, and a tube diameter of approximately 40 nm were found to be consistent with the prepared SBA tube dimension in the solution phase. The approximately 14 nm-increased diameter thickness (from 26 nm of a bare SBA tube to 40 nm) was found to be comparable to the deposited thickness of poly-*p*-xylylene and was due to the deposition on the nonsublimating SBA tube, analogous to conventional poly-*p*-xylylene depositions on a solid substrate [22,38]. To compare the results, the time-dependent particle sizes based on the simulation trajectory and the fabricated particles are plotted in Figure 6, and a recorded simulation video showing a sublimating ice particle template with respect to the elapsed time is included in the supporting information.

**Figure 5.** (**a**) Simulation results showed that Δ*t* = 60 s reduced an initial 50 μm ice particle to 32 μm, Δ*t* = 160 s to 5 μm, Δ*t* = 177 s to 500 nm, and a Δ*t* = 178 s to 250 nm. (**b**) Combinations of characterizations by optical microscopy (OM), SEM, TEM, and AFM of the fabricated anisotropic nanoparticles. Controlled sizes and geometries are shown with respect to the specified Δ*t*.

**Figure 6.** Controlled particle size based on the programmed time-parameter during the two-stage CVSD fabrication process. The simulation and experimental results were mutually consistent. The data points were expressed as the mean value and the standard deviation based on three independent experiments.

#### **4. Conclusions**

In summary, controlling material anisotropy and asymmetry is a fascinating technique adopted from nature which offers superior and diverse properties derived from both the chemical and physical compartments of the materials. The current study demonstrated the challenge of inducing a sensitive, self-assembled SBA protein into a vapor-deposited poly-*p*-xylylene structure that was fabricated via a two-staged stripping and encapsulating process. The reported process employed a vapor-phase construction process using water and ice as templates, and the controllable timing parameter with controlled thermodynamic conditions and mass transport, particularly in the vapor phases of sublimation and deposition, offering flexible control over the particle sizes and aspect ratios in the nanometer range. The sizes of the fabricated anisotropic materials were shown to range from approximately 57 μm to 294 nm, and the simulation results were able to accurately predict the size of fabrication (from 50 μm to 250 nm). With the ability to extend this straightforward CVSD process to other nonsensitive materials (functional materials), we foresee the development of prospective anisotropic and asymmetric materials with unlimited functional products and applications in materials sciences.

**Author Contributions:** Conceptualization, H.-Y.C. (Hsien-Yeh Chen), H.-Y.C. (Hsun-Yi Chen), P.-Y.W., and G.C.; methodology, H.-Y.C. (Hsien-Yeh Chen) and T.-Y.W.; software, T.-Y.W. and M.-C.H.; validation, T.-Y.W., H.-Y.C. (Hsien-Yeh Chen), and P.-Y.W.; formal analysis, T.-Y.W. and C.G.; investigation, T.-Y.W. and C.G.; resources, T.-Y.W., C.G., P.-Y.W., and Z.Z.; data curation, T.-Y.W.; writing—original draft preparation, T.-Y.W., C.G., H.-Y.C. (Hsien-Yeh Chen), H.-Y.C. (Hsun-Yi Chen), and G.C.; writing—review and editing, T.-Y.W. and H.-Y.C. (Hsien-Yeh Chen); visualization, T.-Y.W. and H.-Y.C. (Hsien-Yeh Chen); supervision, H.-Y.C. (Hsien-Yeh Chen), H.-Y.C. (Hsun-Yi Chen), and G.C.; project administration, H.-Y.C. (Hsien-Yeh Chen), H.-Y.C. (Hsun-Yi Chen), and G.C.; funding acquisition, H.-Y.C. (Hsien-Yeh Chen), H.-Y.C. (Hsun-Yi Chen), and G.C. 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, grant number MOST 108-2221-E-002-169-MY3, MOST 108-2218-E-007-045, and MOST 109-2314-B-002-041-MY3. This work was further supported by the "Advanced Research Center for Green Materials Science and Technology" from The Featured Area Research Center Program within the framework of the Higher Education Sprout Project by the Ministry of Education (107 L9006) and the Ministry of Science and Technology in Taiwan (MOST 109-2634-F-002-042). The work was also funded by the NSFC of China, Grant number 51721002, 21861132012, 91956127, and 21975047.

**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**


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## *Article* **One-Step Preparation of Nickel Nanoparticle-Based Magnetic Poly(Vinyl Alcohol) Gels**

## **Jun Li 1,\*, Kwang-Pill Lee 2,3 and Anantha Iyengar Gopalan <sup>3</sup>**


Received: 3 September 2019; Accepted: 29 October 2019; Published: 9 November 2019

**Abstract:** Magnetic nanoparticles (MNPs) are of great interest due to their unique properties, especially in biomedical applications. MNPs can be incorporated into other matrixes to prepare new functional nanomaterials. In this work, we described a facile, one-step strategy for the synthesis of magnetic poly(vinyl alcohol) (mPVA) gels. In the synthesis, nickel nanoparticles and cross-linked mPVA gels were simultaneously formed. Ni nanoparticles (NPs) were also incorporated into a stimuli-responsive polymer to result in multiresponsive gels. The size of and distribution of the Ni particles within the mPVA gels were controlled by experimental conditions. The mPVA gels were characterized by field emission scanning electron microscope, X-ray diffraction, magnetic measurements, and thermogravimetric analysis. The new mPVA gels are expected to have applications in drug delivery and biotechnology.

**Keywords:** one-step preparation; nickel nanoparticles; magnetic poly(vinyl alcohol) gels

## **1. Introduction**

In biological and pharmaceutical fields, polymers have been highly advanced by developing the synthesis methods, controlling the manufacturing steps, and designing the properties. Among such materials, hydrogels have been widely applied to meet versatile requirements [1]. Since PVA hydrogel was first formed by gamma rays in 1958, PVA hydrogels have been largely produced. Hydrogels were used in the biomedical and biotechnological fields in the 1990s. With high water content, elasticity, and biocompatibility, PVA hydrogels are widely used as biomaterials. However, there are limitations for PVA hydrogels: the characteristics of low permeability, stability, and fixation. In order to overcome the limitations, the properties of PVA hydrogels are enhanced by blending with other materials, such as other polymers, metals, and clays. Polyvinyl alcohols (PVA) are synthetic polymers widely used in industrial, medical, and food fields since the early 1930s [2]. PVA is relatively harmless when administered orally, because PVA does not accumulate in the body.

Magnetic metal nanoparticles have gained research interest in biomedicine for their chemical, electrical, and magnetic properties [3]. The crystalline structure, particle size, and magnetic properties of the magnetic metal nanoparticles can control and improve the polymers for applications in medical fields such as treatment of hyperthermia and drug delivery. The chemical composition and fabrication processes have great effects on the sizes and shapes of the particles. To achieve different properties, the synthesis routes are co-precipitation, sol-gel, ball-milling, gamma irradiation, laser irradiation, the photoinduced method, e1tc. The mediated polymers improve the quality of the particles for drug delivery. The idea of drug delivery was proposed by Paul Ehrlich [4]. To reduce the systemic distribution and the required drug dosage, magnetic carriers are used. In 1976, Zimmermann and Pilwat proposed magnetic erythrocytes for the drug delivery. Since then, researchers began to use

magnetic microparticles and nanoparticles (NPs) to target specific sites within the body. Some magnetic carriers are magnetic cores encapsulated in a biocompatible polymeric coating with drug loading capability. The magnetic cores are NPs of metal or metallic oxide moieties with sizes between 1 and 100 nm.

Magnetic gels aconsist of magnetic particles embedded within polymers [5]. The term is commonly applied to magnetic nanoparticles being immersed in a hydrogel. They are stimuli-responsive in an external magnetic field, but biological matter is tolerant to magnetic fields. They are promising candidates for controlled drug release due to the interplay between magnetic and elastic properties. The magnetogels combine the advantages of hydrogels and magnetic nanoparticles [6]. Hydrogels are similar to the cellular matrix with their high portion of water. Magnetic nanoparticles made of any transition metal (Fe, Ni, Co, Cr, or Mn) and its oxides allow for the control to a specific location under a magnetic field. The nanosystems improve target specificity, therapeutic effectiveness, magnetic resonance imageology, and hyperthermia for cancer therapy.

The magnetic drug carrier particles have been applied for over 40 years under an external magnetic field [7]. The magnetic nanoparticles can be removed to reduce risk of particle aggregation after therapy is completed. The nanoparticles can be synthesized in the presence of polyvinyl alcohol (PVA) or another substance to make them appropriate for in vivo applications. The polymers make the NPs stable [8,9]. One of the challenges is to improve the localization of cross-linker release to minimize in vivo toxicity. Controlling the density of cross-links in the gel matrix can tune the porous structure and the swelling properties of the hydrogels in the aqueous environment. Manufacturing processes still need to be developed for better-controlled hydrogels in terms of size, shape, distribution, and mechanics [10].

Ni-based nanomaterials can be used as magnetically-responsive therapeutic platforms for anticancer drugs [11]. Many different synthetic routes for magnetic nanoparticle synthesis have been reported [12] Some of them are one-step, while others are multi-step procedures. They all have advantages and disadvantages. But none of them provides a universal solution for all types of magnetic nanoparticles. Employing simpler and easier synthetic routes for the magnetic nanoparticles with desired characteristics remains highly challenging.

Doxorubicin (DOX) is a kind of cytotoxic anticancer drug. It is widely used in clinical therapy [13] To reduce the limitation of cardiotoxicity and improve the biocompatibility and efficiency, drug delivery systems are being developed [14]. The superparamagnetic nanodevice is designed to deliver DOX under a magnetic field, with good stability [15]. The main aim of this work is to design a new formation strategy of the magnetic polymer for the anti-cancer drug delivery system. An easy, one-step synthetic method for nickel nanoparticles based on poly(vinyl alcohol) gel is described. Nickel chloride was reduced by sodium borohydride in aqueous PVA solution. The Ni PVA gels were characterized by various techniques. The loading and release process was studied.

## **2. Experimental**

## *2.1. Materials*

PVA was purchased from Aldrich (Saint Louis, MO, USA) with a hydrolysis degree of 98%–99% and a molecular average weight of 85,000–124,000 g/mol. Nickel(II) chloride hexahydrated was from Junsei Chemical Co. (Tokyo, Japan). Sodium borohydride was from Kanto Chemical Co. (Tokyo, Japan). Doxorubicin hydrochloride (DOX) was obtained from Korea United Pharm. Inc. (Chungnam, Korea).

## *2.2. Synthesis of Ni-NP PVA Gel*

Ni-NP PVA gel was prepared by reducing nickel chloride in PVA solution with sodium borohydride as the reducing agent. The aqueous solution of 10 wt.% PVA was first prepared by dissolving PVA in deionized water at 80 ◦C for 3 h. An aqueous solution of nickel chloride was then added to the PVA solution in a beaker. NaBH4 solution was added drop wise, using a separation funnel, to the nickel

chloride solution, while the temperature was maintained between 80 ◦C with continuous stirring. The reaction mixture was allowed to stir for about 1 h at 80 ◦C, by which time black-colored solution was obtained, and at the same time a large amount of gas bubbles of H2 were generated. The water in this solution was evaporated at 70 ◦C. After evaporation of most solvent, the hot residue was poured into a Petri dish. The black mixture was dried overnight at room temperature. After that, the residue was dried in an oven at 70 ◦C, and thus a gel was generated. The resulting gel was washed with distilled water, and further dried in an oven at 70 ◦C for 24 h. The resulting product was but to small pieces for various tests with scissors.

## *2.3. Characterization*

The morphology and elemental composition of the structure, and microstructural characterization of the samples, were characterized using field-emission scanning electron microscopy (FE-SEM) (S-4200, Hitachi, Tokyo, Japan), operated at 3–10 kV electron potential difference and equipped with a semiconductor detector that allowed for the detection of energy dispersive X-rays (EDX). X-ray diffraction analysis (XRD) patterns on samples were measured on a model D8-Advanced AXS diffractometer (Bruker, Billerica, MA, USA) using Cu Kα radiation. Samples were supported on glass slides. Measurements were taken using a glancing angle of incidence detector at an angle of 2◦, for 2θ values over 10◦–80◦ in steps of 0.02◦. Magnetic properties of the samples were measured by using a super-conducting quantum interface device (SQUID) magnetometer (MPMS-X 1, Quantum Design, San Diego, CA, USA) under an applied magnetic field, at room temperature. Thermogravimetric analysis (TGA) of the sample was performed on a TGA 7/DX Thermal Analyzer (Perkin-Elmer, Waltham, MA, USA) with a scan rate of 10 ◦C/min by pursuing N2 gas as a carrier at a flow rate of 100 mL/min. The differential scanning calorimetric (DSC) experiment was carried out using a DSC2010 Differential Scanning Calorimeter (TA Instruments, New Castle, DE, USA) over a temperature range 20 to 1200 ◦C at a scan rate of 10 ◦C/min.

## *2.4. Swelling Studies*

For the swelling kinetics' measurement, the gel (about 0.5 mm thick) was cut to 1 cm × 0.5 cm. The gel was then immersed in distilled water. The rate of gel expansion was determined by measuring the change in gel length at various time intervals. The dimensional change was estimated from the length ratio of the swollen hydrogel sample after 24 h (at the equilibrium) compared to its original length. The change in dimensions of the swollen samples (*L*s) was calculated according to equation:

$$L\_{\sf s}(\%) = \frac{L\_{\sf t} - L\_0}{L\_0} \times 100\% \tag{1}$$

where *L*<sup>0</sup> is the original gel length, and *L*<sup>t</sup> is the equilibrium gel length at a given time.

The water uptake of a gel sample was measured at room temperature. A swollen sample was removed from the solvent. After the water on the surface was absorbed gently with filter paper, the sample was immediately weighed. Then, the sample was returned to the medium. The mass swelling ratios (*M*s) was calculated as follows:

$$M\_{\\$}(\%) = \frac{m\_t - m\_0}{m\_0} \times 100\% \tag{2}$$

where *m*t, *m*0, and (*m*<sup>t</sup> − *m*0) are the masses of a swollen sample at a given time, dry gel, and absorbed water, respectively.

The equilibrium water content *W*e in swollen samples was calculated as:

$$\mathcal{W}\_{\mathfrak{C}}\left(\%\right) = \frac{m\_{\infty} - m\_{0}}{m\_{0}} \times 100\% \tag{3}$$

where *m*<sup>∞</sup> is the weight of the swollen hydrogel at equilibrium [16].

## *2.5. Drug Loading and Release*

The drug release kinetics is important for the application of the gel to drug delivery. Doxorubicin hydrochloride (DOX) was chosen as an anti-cancer model drug. The DOX loading was carried out by dispersing 10 mg of PVA gel in 3 mL DOX solution at a concentration of 0.02 mg/mL at room temperature in the dark. The mixture of PVA gel in DOX was shaken (80 rpm) at room temperature for 48 h to facilitate DOX uptake. The optical density of residual DOX in the supernatant was measured at 479 nm by UV-vis spectrophotometer (Cary 50, VARIAN, Palo Alto, CA, USA). The drug loading was determined as the difference between the initial DOX concentration and the DOX concentration in the supernatant. After the supernatant was removed, the release profile was obtained by reimmersing the gel loaded with DOX in 3 mL of water under gentle stirring. The concentration of DOX in the particle-free solution was determined at fixed time intervals by UV-vis spectrophotometry.

## **3. Results and Discussion**

## *3.1. Reaction Mechanism of Ni-NP PVA Gel*

Liquid phase reduction has some advantages over other synthetic methods for magnetic nanomaterials. Using some popular reducing agents, liquid phase reduction was applied to reduce magnetic metal ions to magnetic metal [12]. In this kind of reaction, hydrides are usually used for reducing agents. It is difficult to handle the sensitivity of hydrides to the mild environment. No special laboratory condition is required for these strong reactants, except for the moisture. Hydrides are also penetrative to some polymers, so the particles can still be reduced even with protective coatings.

NaBH4 is a particularly powerful reducing agent in liquid phase reduction. It is soluble in both methanol and water. The mechanism of reduction for Ni using NaBH4 is complicated. In the reaction, a black metal powder of Ni-NPs formed due to an instantaneous Ni2<sup>+</sup> <sup>→</sup> Ni co-reduction reaction with NaBH4 [17]. The reduction of nickel ions using NaBH4 followed the equation

$$\text{NiCl}\_2 + 2\text{NaBH}\_4 + 6\text{H}\_2\text{O} \rightarrow \text{Ni} + 2\text{B(OH)}\_3 + 2\text{NaCl} + 7\text{H}\_2\uparrow$$

The byproducts of B(OH)3 and NaCl remained dissolved in the aqueous solution and were removed after the washing process in fresh water.

Poly(vinyl alcohol) (PVA) is a good stabilizer for small metal particles in chemical synthesis to prevent the agglomeration and precipitation [18]. The embedding of the particles is also advantageous for the PVA casting. The Ni-NP PVA hydrogel was formed by the embedding of Ni-NPs into PVA aqueous solution. The Ni-NP remained in gel form within the matrix, as sketched in Scheme 1. Figure 1a,b shows photographs of the obtained PVA gel and Ni-NP PVA gel, respectively. Figure 1c,d shows photographs of the Ni-NPs and Ni-NP PVA gel synthesized by the same reduction reaction in the magnet field. When the Ni-NPs and the Ni-NP hydrogels were exposed to the field of a permanent NbFeB magnet, they were be attracted to the external magnet. This simple analysis demonstrated the magnetic properties of the polymers containing Ni-NPs.

**Scheme 1.** One step preparation of PVA/Ni magnetic gel.

**Figure 1.** Photographs of (**a**) PVA gel; (**b**) PVA/Ni magnetic gel; (**c**) Ni-nanoparticle (NPs) in the magnet field; (**d**) PVA/Ni magnetic gel in the magnet field.

## *3.2. Structural and Morphological Properties*

To explore the sample morphology, FE-SEM micrographs are presented. Figure 2a–f shows Ni-NPs, PVA gel, and PVA gel with magnetite particles. The surface of Ni-NP PVA gel is rough (Figure 2c) compared to the surface of PVA gel (Figure 2b). The neat PVA gel showed a very smooth surface. Figure 2c shows a micrograph of a Ni PVA gel with pores of about 50 nm. Ni-NPs' formation was evident from the presence of spherical particles over the PVA gel (Figure 2e) compared to the pure Ni-NPs formed by the same reduction reaction in aqueous solution (Figure 2a). From the image in Figure 2c, it can be seen that the magnetic particles are distributed in a compact PVA matrix. The EDX spectrum of PVA gel containing 2.3% Ni (Figure 2d) confirmed the presence of Ni-NPs. The EDX results revealed the percentage of Ni in PVA to be 10.97 wt.%. It is, thus, presumed that Ni-NPs have good dispersion in PVA gel. Compare the FE-SEM micrograph of PVA gel containing 2.3% Ni (Figure 2c) with 4.5% Ni (Figure 2e); the higher the amounts of magnetic nanoparticles, the rougher the surface of the Ni-NP PVA gel is. The EDX results (Figure 2f) revealed the percentage of Ni in PVA to be 41.15 wt.%.

**Figure 2.** *Cont*.

**Figure 2.** FESEM micrograph of (**a**) Ni-NPs; (**b**) PVA gel; (**c**) PVA gel containing 2.3% Ni; (**d**) EDX spectrum of PVA gel containing 2.3% Ni; (**e**) PVA gel containing 4.5% Ni; and (**f**) EDX spectrum of PVA gel containing 4.5% Ni.

Figure 3 shows the XRD spectra of PVA gel without or with Ni-NPs of different concentrations, and pure Ni-NPs obtained by the reduction reaction of nickel chloride with sodium borohydride. The XRD pattern of Ni-NPs (Figure 3d) shows that the products are metallic nickel. The average crystallite sizes were calculated from the peak broadening of XRD patterns by the Scherrer equation. The average size of the Ni-NPs was 7.0 nm. The X-ray diffractograms indicate a crystalline structure with some amorphosity of the disordered surface layers. The XRD spectra of Ni-NPs for 2θ values from 10◦ to 80◦ is similar to the literature, despite that an additional peak, the largest, at 2θ = 11.98◦ (intensity *I*<sup>p</sup> = 100 u), was not mentioned [19]. The XRD Patterns do not correspond to face centered cubic (fcc) nickel with space group Fm3m, but are indexed for a tetragonal crystal structure with space group 14/mcm [17]. Those peaks at 2θ = 24.96◦ (*I*<sup>p</sup> = 7.0 u), 2θ = 34.06◦ (*I*<sup>p</sup> = 16.2 u), 2θ = 36.03◦ (*I*<sup>p</sup> = 3.7 u), 2θ = 45.04◦ (*I*<sup>p</sup> = 6.8 u), 2θ = 60.00◦ (*I*<sup>p</sup> = 8.7 u), and 2θ = 71.96◦ (*I*<sup>p</sup> = 2.2 u) are attributed to (110), (002), (200), (211), (310), and (321) facets of the tetragonal crystal structure of Ni, respectively. The lattice parameters are *a* = 0.4887 nm and *c* = 0.4516 nm for the nickel nanoparticle sample.

Figure 3a–c compares the X-ray diffractograms of PVA gel containing different amounts of Ni-NPs. The interplaner distance was calculated. XRD patterns of the pure PVA gel (Figure 3a) indicate the characteristic peak for poly(vinyl alcohol) at a 2θ value of 19.82 with the inter planner distance of 0.4476. The same phenomenon has been observed by Clémenson et al. [20]. PVA, known to be semicrystalline in nature, shows a single broad peak. The relatively sharp and broad peak centered on about 19 indicates that the semicrystalline nature of the polymer PVA contains crystalline and amorphous regions. It is clear from Figure 3b,c that the peaks representing the tetragonal crystal structure of Ni in XRD disappear, indicating the conversion of nickel nanoparticle to the amorphous Ni-NP PVA gel. For both PVA gel with 1.5 mol% Ni and 2.3 mol% Ni-NPs, a broad peak centered on a 2θ value of about 19 is observed. It suggests the characteristic peak of PVA gel, indicating that the Ni-NP PVA polymer intercalation occurs with the formation of Ni-NPs within the polymer matrix. The amorphous broad band decreased for PVA gel containing 2.3% Ni, indicating the metal interacts with the polymer chain [21]. From the Figure 3a–c, comparing the XRD spectra of the bulk Ni-NPs,

the smaller spectral intensity of the diffraction peak of PVA in the composite is due to the presence of a higher content of Ni-NPs in hydrogels.

**Figure 3.** XRD patterns for PVA gel containing (**a**) 0% Ni; (**b**) 1.5% Ni; (**c**) 2.3% Ni; and (**d**) pure Ni nanoparticles.

#### *3.3. Magnetic Properties*

The magnetic properties of PVA gel with or without Ni-NPs of different concentration, and pure Ni-NPs were measured. Figure 4 shows the magnetization curves versus external magnetic field for neat PVA gel, or PVA gel with Ni-NPs of different concentration and pure Ni-NPs. Magnetization as a function of increasing external field up to 50 kOe at temperature of 5 K has been measured for selected samples. It can be found that the magnetization of pure Ni-NPs is obviously higher than other samples. The magnetic properties of Ni-NP PVA gel significantly changed with the composite formation. The magnetization of Ni-NP PVA gel is far weaker than that of the pure Ni-NPs. The greater the amount of Ni the Ni-NP PVA gel contains, the higher the value the magnetic moment shows at the same magnetic field. In contrast, the curve of the neat PVA gel is mostly flat. It is well known that the polymer is antimagnetic [22]. As shown in Figure 4, the curve of Ni-NPs shows that the magnetic moment does not reach full saturation even at 50 kOe. Among all these samples, the maximum magnetic moment value is 68.6 emu/g for pure Ni-NPs. The magnetization of the Ni-NPs and Ni-NP PVA gel samples increases quickly as the intensity of applied magnetic field increased below 10 kOe. The magnetization of the samples containing Ni increases slowly to reach saturation when the magnetic field above 10 kOe increases. The saturated magnetizations of the PVA gel containing 2.3 mol% Ni and PVA gel containing 4.5 mol% Ni at 50 kOe are 5.1 and 12.9 emu/g, respectively. Generally, the saturation of the magnetic moment increases with a higher amount of Ni-NPs. This is in agreement with the results of SEM and XRD.

In order to better understand the magnetic behavior, the samples were measured in the zero-field-cooled (ZFC) and field-cooled states (FC). The ZFC and FC measurements were performed by cooling the nanoparticles at zero field or in the presence of an external field. To obtain the ZFC measurement, the samples were cooled to 5 K in the absence of an external field. An external magnetic field of 1000 Oe was then applied to the samples. The magnetization was measured as the temperature was increased. The samples were then immediately cooled to 5 K in the presence of a magnetic field of 1000 Oe for the FC measurement. Figure 5a–c compares the magnetization versus temperature measured under ZFC and FC conditions for PVA gel without or with Ni-NPs of different concentration, and pure Ni-NPs, respectively. The effect of cooling is clearly observed. ZFC and FC magnetization curves split below *T* = 84, 54, and 35 K for PVA gels containing 2.3% Ni, 4.5% Ni, and the pure Ni gel, respectively. The difference in the magnetic moment signals is due to a two times larger amount of Ni in sample of PVA gel containing 4.5% Ni with respect to sample PVA gel containing 2.3% Ni, and the pure Ni without PVA. Samples of Ni-NP PVA gels and pure Ni nanoparticles present a significant magnetic irreversibility. The blocking temperature (*T*B) of a magnetic nanoparticle can be measured where the FC and the ZFC curves diverge. The ZFC magnetization curve exhibits a peak around *T* = 15 K. This temperature indicates a collective freezing of magnetic moments. The signal of a magnetic moment is relatively weaker when the amount of Ni in the sample is lower. On the other hand, no temperature dependence is observed for the PVA without Ni sample. It is confirmed that pure PVA is diamagnetic. The results establish the role of Ni loading within the PVA polymer for the magnetic moments. Figure 5a–c indicates that for all the samples containing Ni, the temperature dependence of the magnetic moment does not follow Curie's law at temperatures above 100 K. The difference between PVA gel and Ni-NP PVA gel suggests that magnetic response of the Ni particles manifests a certain degree of magnetic interaction [23].

**Figure 4.** Magnetization curves of the PVA gel containing (**a**) 0% Ni; (**b**) 2.3% Ni; (**c**) 4.5% Ni; and (**d**) pure Ni nanoparticles.

**Figure 5.** *Cont*.

**Figure 5.** Field-cooling (FC) and zero-field-cooling (ZFC) magnetic moment curves measured for samples: (**a**) PVA gels containing 2.3% Ni; (**b**) 4.5% Ni; and (**c**) pure Ni nanoparticles. The lower curve of each sample represents the ZFC measurement and the upper one represents the FC measurement.

#### *3.4. Thermogravimetric Analysis*

To study the thermal changes in the PVA gels containing Ni, the samples were thermally characterized by TGA and DSC. A typical TGA curve of the PVA gel containing Ni is presented (Figure 6). The thermogram suggests that weight loss mainly occurs in three steps. First step corresponds to the loss of water molecules and slow removal of impurities from room temperature to 192 ◦C with a total weight loss of 14% in a PVA gel containing Ni. The second step corresponds to the removal of oligomers and the decomposition of PVA in the range of 266–294 ◦C with a weight loss of 29% [22]. The third step corresponds to the degradation of polymeric backbone in the range of 385–986 ◦C with a weight loss of 48% in the sample. At last, about 5% weight residue remained at a temperature above 986 ◦C, due to the Ni magnet. There is nearly no residue that remained above 986 ◦C in the PVA alone. This result is consistent with Khanna et al. [18]. Figure 6 shows the DSC thermogram of the PVA gel containing Ni [24].

**Figure 6.** TGA curve and DSC plot of Ni-NP PVA gel.

#### *3.5. Swelling Studies*

Swelling experiments can yield important information concerning the stability of PVA gels in solution [25]. It can also be used to trigger drug release by controlling the hydrogel swelling properties. Water uptake was measured to determine the swelling abilities of the hydrogels. This is most commonly

used to evaluate the effects of different synthesis methods of PVA on water transport. The swelling kinetic curves of the hydrogels are shown in Figures 7 and 8. Figure 7 shows the dimension changes as a function of the exposure time for different PVA gels containing nickel in the amounts of 0 mol%, 1.5 mol%, 2.3 mol%, and 4.5 mol%. The mass change as a function of time from 5 min to 24 h for different hydrogels in water is shown in Figure 7. From Figures 7 and 8, the samples attained their maximum swelling ratio at about 16 h. The results of the dimension changes are in agreement with the results of the mass changes. The equilibrium water contents in swollen samples (*W*e) were 364.60%, 312.36%, 230.16%, and 166.67% for neat PVA gel, and PVA gels with 1.5 mol%, 2.3 mol%, and 4.5 mol% Ni-NPs, respectively. From both Figures 7 and 8, it can be observed that the swelling ratio decreases as the amount of Ni-NPs in a PVA gel increases. The NPs have effect an on the swelling properties of the PVA gel by a simply physical method. The interaction of the PVA chains with Ni-NPs might induce the formation of low-mobility regions that can act as additional physical cross-linking points. The swelling ratio decreases with the presence of NPs, pointing to an increase of the physical crosslinking density as reported for the formation of PVA ferrogels through a freezing–thawing procedure [26].

**Figure 7.** Dimension change (*L*s) versus immersion time for PVA gel and PVA/Ni magnetic gel.

**Figure 8.** Water absorption (*M*s) as a function of the exposure time for PVA/Ni magnetic gel.

## *3.6. Drug Loading and Release*

To design the magnetic polymer for the anti-cancer drug delivery system, a well-known anti-cancer drug was needed. It was convenient to investigate the drug release properties in vitro. Furthermore, regarding swollen polymers—the drug is in an aqueous environment; its water solubility becomes an important consideration. As one of the most common anti-cancer drugs, DOX was chosen in this study to model the release action of anti-cancer drug from Ni-NP PVA gel. Most of the DOX was up-taken about in 24 h. A loading capacity of about 96% of the drug weight was determined. This result was used for the evaluation of the cumulative release of DOX. To investigate the drug delivery properties, Ni-NP PVA gels with a cargo amount of 5.8 mg/g were studied.

The release kinetics measured for Ni-NP PVA gel under two different conditions of 25 ◦C and 37 ◦C is shown in Figure 9. Both curves were measured by the UV-vis absorbance in water. From the figure, the curve with temperature of 25 ◦C shows a relatively low cumulative release of DOX. The corresponding cumulative release rapidly achieves the level of about 10% in 10 h. About 15% of the drug was released in 48 h. The release rate was much higher with the temperature of 37 ◦C. The corresponding cumulative release rapidly achieved the level of about 39% in 10 h. In total, 73% of drug was released in 48 h. This probably manifests through the fact that the Ni-NP PVA gel is made from simply physical cross-linking. This kind of gel is flexible when put in wet or very humid conditions under high temperatures. To study the release behaviors of the Ni-NP PVA gel, the cumulative release ratio (*M*t/*M*∞) was calculated based on the classic Korsmeyer-Peppas equation [27]. *M*t/*M*<sup>∞</sup> is the fraction of drug released after time t relative to the amount of drug released at infinite time. There are two stages in the drug release curve. The initial diffusion of drugs mainly happened in the outermost layer of the Ni-NP PVA gel. At that stage, the drug released rapidly. At the second stage, the release was under diffusion's control and the drug released relatively slowly.

**Figure 9.** Drug release profiles of Ni-NP PVA gel in water.

#### **4. Conclusions**

The magnetic hydrogels have sparked particular interest in the anti-cancer drug delivery applications. PVA is a type of biocompatible material used for biomedicine. In this work, a new kind of nickel nanoparticle-based magnetic PVA gel was synthesized by a one-step procedure. The nickel nanoparticles and cross-linked mPVA gels were formed simultaneously. The structural and morphological properties of Ni-NP PVA gel were measured by FE-SEM and X-ray diffraction. The surface of Ni-NP PVA gel was rougher than that of the neat PVA gel. The amount of Ni-NPs detected was consistent with the theoretical value. The peak of PVA in the XRD spectra was weaker

than that of Ni-NP PVA gel. The magnetic and thermal properties were also measured. With a higher amount of Ni-NPs, Ni-NP PVA gel had higher magnetic moments. There are noteworthy features for the Ni-NP PVA gel in terms of ferromagnetism and thermally stability. All measurements confirmed the good formation of the Ni-NP PVA gel. The release of DOX showed diffusion control in vitro. With a higher temperature, the release rate of DOX was higher. The Ni-NP PVA gels are expected to be applied for controlled drug delivery. The novel, simple synthetic method can be used to form other magnetic gels for biotechnology.

**Author Contributions:** Conceptualization, K.-P.L.; methodology, A.I.G.; data curation, J.L.; writing—original draft preparation, J.L.; visualization, J.L.; supervision, K.-P.L.; project administration, K.-P.L. and A.I.G.; funding acquisition, K.-P.L.

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

**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* **Proteomic Analysis of Biomaterial Surfaces after Contacting with Body Fluids by MALDI-ToF Mass Spectroscopy**

**Makoto Hirohara 1,**†**, Tatsuhiro Maekawa 1,**†**, Evan Angelo Quimada Mondarte 1, Takashi Nyu 1, Yoshiki Mizushita <sup>1</sup> and Tomohiro Hayashi 1,2,3,\***


Received: 22 October 2019; Accepted: 17 December 2019; Published: 22 December 2019

**Abstract:** We developed a method to identify proteins adsorbed on solid surfaces from a solution containing a complex mixture of proteins by using Matrix-Assisted Laser Desorption/Ionization-Time of Flight mass (MALDI-ToF mass) spectroscopy. In the method, we performed all procedures of peptide mass fingerprint method including denaturation, reduction, alkylation, digestion, and spotting of matrix on substrates. The method enabled us to avoid artifacts of pipetting that could induce changes in the composition. We also developed an algorithm to identify the adsorbed proteins. In this work, we demonstrate the identification of proteins adsorbed on self-assembled monolayers (SAMs). Our results show that the composition of proteins on the SAMs critically depends on the terminal groups of the molecules constituting the SAMs, indicating that the competitive adsorption of protein molecules is largely affected by protein-surface interaction. The method introduced here can provide vital information to clarify the mechanism underlying the responses of cells and tissues to biomaterials.

**Keywords:** protein identification; matrix-assisted laser desorption/ionization-time of flight; proteomics; quartz crystal microbalance; self-assembled monolayer

## **1. Introduction**

Scaffolding protein in extracellular matrices (ECM) plays a significant role in determining cellular responses. There are many situations where the scaffolding proteins govern the cellular responses—biodevices implanted in a body [1], patterning of cells on solid substrates [2], and general cell culturing in dishes, among others. For the thorough understanding of the mechanism underlying the cellular behavior, we need to clarify the properties of the layer of the scaffolding proteins [3,4].

Investigation of the adsorbed proteins has been done mostly in terms of the amount of adsorption. They are measured with surface plasmon resonance (SPR) spectroscopy [5], quartz crystal microbalance (QCM) [6–8], and surface acoustic wave spectroscopy (SAW) [9–11] for real-time adsorption kinetic measurements; and X-ray photoelectron spectroscopy (XPS) [12,13], Fourier transform infrared absorption spectroscopy (FTIR) [14–16], and fluorescence microscopy with labeling molecules [17,18] for samples in dry state after adsorption.

In contrast with the number of proteins in the scaffolding layer, the composition of proteins in the protein layers formed on biomaterials after contacting body fluids or cell-culturing media has not been

intensively studied even though this information is essential for the understanding of the mechanism underlying the responses of the adhered cells to materials. In general, the identification of proteins in solution is carried out by a peptide mass fingerprint method [19]. In this method, the proteins in the solution are separated first by electrophoresis [20]. Then, the separated proteins are fragmented into peptides by digestion with trypsin. Finally, the proteins can be identified from the pattern of the peaks in the Matrix Assisted Laser Desorption/Ionization-Time of Flight mass (MALDI-ToF mass) spectrum. However, in the case of the scaffolding proteins on the surface of biomaterials, especially on small surface areas, this procedure is difficult to employ because the total amount of adsorbed proteins is often not enough for a reliable and efficient separation by electrophoresis which, in general, requires ~20–30 μg of proteins.

Several works that overcame the problem have been reported. One is the usage of nanoparticles to acquire a large interfacial area between materials and proteins [21]. However, this approach is not universal because many biomaterials cannot be fabricated in the form of nanoparticles. Another is the proteomic analysis of membrane filters for hemodialysis wherein the proteins adsorbed on the membrane filters after hemodialysis were collected by an eluting solvent and then isolated by electrophoresis followed by MALDI-ToF mass measurements [22,23]. Like the first example, this approach also has the merit of having a larger interfacial area. Unfortunately, this method cannot be employed for biomaterials with a flat surface. Moreover, proteins with relatively high hydrophobicity can easily adhere to the walls of pipets resulting in the change in the composition of the protein and thus greatly affects the results especially for a small collection of proteins [24]. Therefore, it is important to perform all steps of the proteomic analysis directly on the surface of the materials. Although this approach was already performed by Kirschhofer et al., who have succeeded in the identification of the adsorbed proteins, the proteins used in their reports consist of only three kinds with known compositions [25].

In this work, we performed the full peptide mass fingerprint method on the biomaterial surfaces and obtained mass spectra by MALDI-ToF mass spectroscopy of the adsorbed proteins from the more complex protein serum. To extract the composition of the adsorbed proteins from the resulting complex spectra, we developed an algorithm to give assignments to the peaks by collating the experimental and theoretical peak positions of peptide fragmented from the serum proteins. Furthermore, we established a method to evaluate the molar ratio between proteins in the scaffolding layer. In general, the composition of proteins strongly depends on the analytical techniques used [26–29]. To solve this problem, we thereby employed a method to evaluate molar ratios between proteins by using the results obtained from original standard samples with a known molar ratio. We believe that our method may shed light on the role of the scaffolding layer in determining the fates of cells or tissues after being in contact with biomaterials.

## **2. Materials and Methods**

#### *2.1. Self-Assembled Monolayers and Their Substrates*

Metal substrates for self-assembled monolayers (SAMs) of alkanethiols were prepared by thermal evaporation under vacuum (base pressure: 4 <sup>×</sup> 10−<sup>6</sup> Pa). A 5-nm Ge (adhesion promoter) was first deposited on a glass plate (18 <sup>×</sup> 18 mm2, Matsunami glass inc. Ltd., Oosaka, Japan) followed by the deposition of 100-nm Au film. The glass plates underwent prior cleaning by sonication in ethanol, then in pure water, and lastly dried with nitrogen gas.

Fabrication of the SAMs was carried out by immersing the substrates in a thiol solution in ethanol (0.1 mM) for 24 h. We used five kinds of thiols in this work (Table 1) [Sigma-Aldrich (St. Louis, MI, USA) and ProChimia Surfaces (Gdansk, Poland)]. After the immersion, the samples were carefully rinsed with pure ethanol to remove the physisorbed thiol molecules from the surface.


**Table 1.** Thiols used in this work and their chemical structures.

#### *2.2. QCM Measurements*

The amounts of proteins adsorbed onto the SAMs from protein solution and fetal bovine serum (FBS) (Equitech-Bio Inc, Kerrville, TX, USA, lot number: SFBM30-2485) were measured by QCM (D300 Q-Sense, Gothenburg, Sweden). The sensors were cleaned by UV-Ozone treatment, followed by rinsing with ethanol and pure water. The fabrication of the SAMs on the Au-coated QCM sensors was done with the same procedure as the case of the Au/Ge/glass substrates. For the adsorption of bovine serum albumin (BSA) (Sigma-Aldrich, St. Louis, MI, USA) and fibrinogen (Biogenesis, Poole, England), we prepared the protein solution by dissolving the proteins in phosphate buffer saline (PBS) (Sigma Aldrich, St. Louis, MI, USA) at a concentration of 1 mg/mL. For the adsorption of proteins from the serum, the FBS was diluted to 20% in volume with PBS.

In the QCM measurements, the measurement cell was first filled with PBS for background measurements, then the protein or serum solution was injected. After the frequency shift was equilibrated, PBS was injected again for rinsing. All measurements were performed at 25 ◦C We calculated the amount of the adsorbed protein from the Sauerbrey equation (Equation (1))

$$
\Delta m = -\frac{\mathbb{C} \cdot \Delta f}{n},
\tag{1}
$$

where *C*, Δ*f*, *n* are the conversion constant (17.7 ng cm−<sup>2</sup> Hz<sup>−</sup>1), change in the resonant frequency and overtone number (*n* = 3 in this work), respectively.

#### *2.3. Formation of Protein Layer on SAMs from FBS for Peptide Mass Fingerprinting*

To form a layer of serum proteins on the SAMs, 50 μL of the 20% (v/v) FBS were placed on the SAMs for 24 h at room temperature. Then, the SAMs were rinsed with PBS buffer and dried in air.

#### *2.4. Digestion of the Serum Proteins Adsorbed on the SAMs*

The procedure of the preparation of samples for MALDI-ToF mass measurements (denaturation, reduction, alkylation, digestion, and spotting of a matrix) is summarized in Figure 1. A mixture of 25 μL of 10 mM ammonium bicarbonate aqueous solution (Wako, Osaka, Japan), 25 μL of 2,2,2-trifluoroethanol (Wako, Osaka, Japan), and 10 μL of 20 mM dithiothreitol aqueous solution (Sigma-Aldrich, St. Louis, MI, USA) was placed on the dried protein layer then was left still in an oven at 60 ◦C under atmospheric pressure for 45 min to denature and reduce the adsorbed proteins. Then, the sample was taken out from the oven, cooled to room temperature, added with 40 μL of 20 mM iodoacetamide (Sigma-Aldrich, St. Louis, MI, USA) aqueous solvent to allow alkylation, and was allowed to stand at room temperature for 1 h. In order to quench the iodoacetamide aqueous solvent, 10 μL of an aqueous dithiothreitol solution used in the reduction was added and then dried at 60 ◦C under vacuum. Next, the proteins underwent digestion with trypsin (Promega, Madison, WI, USA). A 200 μL of trypsin-acetic acid solution was prepared by dissolving 20 μg trypsin in 1 mM acetic acid (Promega, Madison, WI, USA). A 10-μL aliquot of the trypsin-acetic acid solution was mixed with 30 μL of a 10 mM ammonium hydrogen carbonate aqueous solution, and 10 μL of the mixture was dropped onto the sample and kept for 4 h. To stop the digestion, 50 μL of 10% trichloroacetic acid (Wako, Osaka, Japan) was added and then the samples were dried in a desiccator at 60 ◦C. The digestion of proteins in 20% (v/v) FBS was performed with the same procedure.

**Figure 1.** Flow of preparation of samples for MALDI-ToF mass spectroscopic measurements.

The samples were fixed on a slide glass coated with an indium tin oxide (ITO) film with carbon tape. To be able to check if the samples are of the same height level, a commercial peptide solution (Peptide Calibration Standard II: PCS II) (Bruker, Billerica, MA, USA) was placed on all the samples. After that, the matrix solution (α-cyano-4-hydroxycinnamic acid (Sigma-Aldrich, St. Louis, MI, USA)) saturated in a mixture containing 30% of acetonitrile (Wako, Osaka, Japan), 69.9% of water and 0.1% of trifluoracetic acid (TFA) (Wako, Osaka, Japan) is placed on the samples to be measured then dried. The height of the samples was adjusted by placing blank substrates with the same thickness as a spacer at the corners (Figure 2).

**Figure 2.** Samples (protein/SAM/Au/Ge/glass) fixed on a MALDI target plate after all the processes presented in Figure 1.

## *2.5. Standard Samples to Calibrate Peak Intensity*

A mixture of proteins of interest with a known ratio was prepared to calibrate the effect of ionization of peptides, which is dependent on the peptide sequence, on their intensities in mass spectra. In this work, we dissolved BSA, vitronectin, and fibronectin in PBS and prepare the mixed solution at a molar ratio of 1:1.29:0.255 (weight ratio of 1:1:1). This solution was spotted on an ITO substrate and processed as the same procedure as the case of serum proteins on the SAMs.

## *2.6. MALDI-ToF Mass Spectroscopic Measurements*

We used a commercial MALDI-ToF mass spectrometer (ultrafleXtreme, Bruker Daltonics, Billerica, MA, USA) for our measurements. Mass spectra in the range of 500 to 4000 Da were measured in the positive-ion reflector mode with 5000 laser shots with raster. FlexAnalysis software (Bruker Daltonics, Billerica, MA, USA) was used to assign peaks. For each sample, 10 spectra obtained at different positions of the substrates were accumulated after the calibration of m/z values (linear offset in m/z) using the fragment peak from BSA (m/z = 927.49). We prepared two samples for each SAM and measured 15 spectra for each sample giving a total of 30 spectra for each SAM.

#### *2.7. Analysis of the Obtained Mass Spectra*

#### 2.7.1. Finding m/z Values Unique for Proteins of Interest

From the complex spectra we obtained, we attempted to find m/z values unique for each protein. Basing on the top 20 abundant proteins plus the proteins associated with cell adhesion (fibronectin and vitronectin) (Table 2) in FBS, we constructed a list of all possible fragments after digestion using a WEB-based database of ExPASy (https://www.expasy.org/) [30,31]. We omitted m/z values that are close to each other within 0.5 Δm/z. Furthermore, we also removed m/z values that are close (0 ≤ m/z < 3) to that of the isotopes of the peptide fragments of BSA, because BSA is dominant (around 60%) in serum and even the isotopes of its fragments exhibit peaks with considerable intensity. Finally, from the remaining m/z values for each protein, we selected the one with the highest intensity in the spectra and denote as the 'reference m/z' hereafter (Figure 3).

**Figure 3.** Procedure for determination of the reference m/z values.


**Table 2.** Proteins of interest in serum (abundant and those with RGD moieties), their molecular weights, and reference m/z obtained by the procedure presented in Figure 3.

## 2.7.2. Conversion from Peak Intensities to Molar Ratio

The peak intensity in the mass spectrum linearly correlates to the concentration of the peptide in the sample [32]. However, the dependence of the signal intensity on the sequences of peptide must be calibrated. In this work, we employ Equation (2) to evaluate molar ratios,

$$\mathcal{M}\_{a/b} = \frac{I\_a}{I\_b} \cdot \mathcal{K}\_{a/b^\*} \tag{2}$$

where *I*, *M*, and *K* are the intensities of the reference peaks, molar ratio between two proteins, and coefficients to calibrate the intensity and molar ratios, respectively. Using the standard sample with known composition, *K*vitronectin/BSA and *K*fibronectin/BSA were calculated to be 0.302 and 1.51, respectively.

## **3. Results and Discussion**

First, we verified whether the processes of the digestion affect the spectral patterns. Figure 4 compares the spectra obtained from proteins digested in different environments—on substrate and in solution. The overall intensity of the spectrum (a) is about 10 times weaker than that of the spectrum (b), which is rationalized by the total amounts of BSA molecules measured. The peak positions assigned to the fragments of BSA and the relative ratio between the peaks of the fragments are consistent, indicating that the digestion of BSA on C8-SAM provides the same result as in solution. It should be noted that there is disagreement in the peak intensity of the other peaks. A possible reason is a different ratio among BSA, trypsin, and other chemicals in the different environments of digestion.

**Figure 4.** MALDI-ToF mass spectra of digested BSA. (**a**) BSA adsorbed on the C8 SAM then digested (procedure shown in Figure 2) (**b**) BSA was digested in a vial and measured on an C8 SAM.

The results of the adsorption tests with proteins of single composition (Figure 5) showed that EG3-OH exhibited strong resistance to fibrinogen and BSA, whereas the other SAMs adsorbed these proteins. These results are consistent with previous findings reported from our group and others [33–38] showing that SAMs of oligo(ethyleneglycol)-terminated alkanethiols exhibit strong protein-resistance to various proteins at relatively low concentrations of protein solution (typically < 2 mg/mL). In contrast with the adsorption of single protein at low concentration, the results of adsorption tests with the serum diluted to 20% with PBS revealed that all the SAMs used in this work adsorbed serum proteins. Considering that the adsorption amount of about 100 ng/cm<sup>2</sup> corresponds to a monolayer of fibrinogen [35], all the SAMs were certainly covered with a layer of serum proteins. It should also be noted that this concentration of serum employed in this work is often used for cell culturing. There is strong dependence of the amounts of adsorbed proteins on both the concentration and composition of the proteins in the original solution [39]. This is one of the good examples that demonstrate that conventional protein adsorption assays cannot predict the protein adsorption from serum or other body fluids.

**Figure 5.** Amounts of adsorbed proteins and changes in the resonant frequency measured by QCM. (**a**) Adsorption of BSA and fibrinogen from 1 mg/mL solution in PBS (*n* = 4) and (**b**) adsorption of serum proteins from 20% serum diluted with PBS (*n* = 3). Error bars and n denote standard deviation and the number of measurements, respectively.

After the confirmation of the formation of the protein layers on the SAMs by QCM-D, we processed the adsorbed protein as the procedure shown in Figure 1. Figure 6a–e shows mass spectra of digested adsorbed proteins on the SAMs. In contrast with the peaks assigned to BSA, the patterns of other peaks, e.g., in the ranges of m/z = 1300–1400 and 1600–1700 are different depending on the SAMs. This indicates that the composition of proteins other than BSA is different depending on the terminal groups of the molecules constituting the SAMs.

**Figure 6.** MALDI-ToF mass spectra of fragmented proteins on the SAMs; (**a**) C8, (**b**) OH, (**c**) COOH, (**d**) NH2, and (**e**) EG3-OH SAMs. The locations of the reference m/z of BSA, vitronectin and fibronectin are shown (m/z = 927.49, 946.48 and 1323.71 for BSA, vitronectin, and fibronectin, respectively).

The relative peak intensities of the reference m/z for each protein with respect to that for BSA for each SAM are summarized in Table 3. Although we cannot discuss the exact amount of proteins from peak intensities, the change in the composition of proteins after adsorption from serum onto SAMs can be easily seen from the results. The final composition of the adsorbed proteins on materials is a result of competitive adsorption of serum proteins (Vroman effect) [40,41]. In theoretical models proposed so far, small proteins adsorb first because of their high mobility and replaced with larger ones. However, the complexity of the problem includes a wide variety of protein-protein and protein-surface interactions. Moreover, the conformational changes of proteins after adsorption complicates the issue. Therefore, it is very difficult to correlate the composition with the physicochemical properties of the SAMs. It should also be noted that we also compared the compositions of vitronectin and fibronectin with respect to BSA obtained by the conventional methods, which incorporate pipetting processes, and our method (Table S1), clearly indicating the effect of pipetting on the ratio of the peak intensities.

The main finding here is that the compositions of vitronectin and fibronectin on the SAMs, which provide RGD moieties essential for cell adhesion, are higher than in serum (Figure 7). The evaluation of the exact amounts of the serum proteins requires standard samples containing the proteins. Unfortunately, many of the proteins listed in Tables 2 and 3 were not available, and we were not able to evaluate the exact amounts by combining with the results of QCM-D. But fortunately, there have been several works that attempted to predict peak intensities of peptides from their sequences by using techniques of informatics [42–44]. These approaches may realize the exact evaluation of the protein composition in the future.


**Table 3.** Relative peak intensity of the reference m/z for each protein with respect to that of BSA in FBS and on the SAMs. Errors denote a standard deviation of 30 spectra (two substrates for each SAM with 15 spectra each at different positions).

#### **4. Summary and Conclusions**

In this paper, we proposed a new method to evaluate the composition of adsorbed proteins on solid surfaces. In this method, we perform denaturation, alkylation, digestion, and measurements without the collection of proteins. This method enables us to collect all the peptide fragments without the artifact of pipetting. To find reference m/z values used to evaluate the protein composition, we first looked for the m/z value unique for the proteins of interest by listing up m/z values of all possible fragments of the top 20 proteins abundant in serum and the interesting proteins, i.e., those with RGD moieties. Then, we determined a correlation factor to evaluate molar ratios between the proteins from peak intensities by using standard samples with the known composition of proteins.

Our results clearly showed that the composition of the proteins adsorbed from serum onto SAMs critically depends on the terminal groups of the molecules constituting the SAMs as a result of the Vroman effect (competitive adsorption of proteins onto surfaces). Moreover, in this work, we evaluated the ratios of vitronectin and fibronectin to BSA. The compositions of vitronectin and fibronectin govern the adhesion of cells that have integrin-mediated adhesion systems since RGD moieties of the proteins provide binding sites for integrins. We found a clear correlation between cell adhesion and the compositions of vitronectin and fibronectin. This finding will be published elsewhere.

**Supplementary Materials:** The following are available online at http://www.mdpi.com/2079-6412/10/1/12/s1. For quality assessment of identification of proteins. Here we simply show the matching ratio between theoretical m/z values and experimentally observed peak positions, since we removed some of the m/z values and cannot employ ROC plot analysis. In the case of the standard sample, which contains BSA, fibronectin, and vitronectin, more than 95% of the theoretical fragment m/z values were observed in the spectra. In the case of serum proteins, the matching rate is lower (Table S2). If the matching rate is larger than 25%, it can be considered to be marginally matched [45]. Therefore, we consider that the identification is satisfactory in this work.

**Author Contributions:** M.H. and T.M. did all the experiments and analysis. T.N. and E.A.Q.M. contributed to the coding of the software. Y.M. and T.H. designed this project and optimized the experimental conditions. All authors have read and agreed to the published version of the manuscript.

**Funding:** The author (T.H.) acknowledges the financial supports by KAKENHI (19H02565, 17K20095 and 15KK0184) and JST- PRESTO.

**Acknowledgments:** The authors appreciate the help of Ms. Kazue Taki for the administration of this project.

**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* **Chemical and Biological Roles of Zinc in a Porous Titanium Dioxide Layer Formed by Micro-Arc Oxidation**

**Masaya Shimabukuro 1, Yusuke Tsutsumi 2,3, Kosuke Nozaki 1, Peng Chen 2, Risa Yamada 1, Maki Ashida 2, Hisashi Doi 2, Akiko Nagai <sup>4</sup> and Takao Hanawa 2,\***


Received: 12 October 2019; Accepted: 27 October 2019; Published: 29 October 2019

**Abstract:** This study investigated the time transient effect of zinc (Zn) in the porous titanium dioxide formed by micro-arc oxidation (MAO) treatment routinely performed for Zn-containing electrolytes. The aim of our analysis was to understand the changes in both the chemical and biological properties of Zn in physiological saline. The morphology of the Zn-incorporated MAO surface did not change, and a small amount of Zn ions were released at early stages of incubation in saline. We observed a decrease in Zn concentration in the oxide layer because its release and chemical state (Zn2<sup>+</sup> compound to ZnO) changed over time during incubation in saline. In addition, the antibacterial property of the Zn-incorporated MAO surface developed at late periods after the incubation process over a course of 28 days. Furthermore, osteogenic cells were able to proliferate and were calcified on the specimens with Zn. The changes related to Zn in saline had non-toxic effects on the osteogenic cells. In conclusion, the time transient effect of Zn in a porous titanium dioxide layer was beneficial to realize dual functions, namely the antibacterial property and osteogenic cell compatibility. Our study suggests the importance of the chemical state changes of Zn to control its chemical and biological properties.

**Keywords:** zinc; titanium oxide; micro-arc oxidation; antibacterial activity; osteogenic cell compatibility

## **1. Introduction**

Recent studies have reported that biomaterial-associated infections caused by the formation of biofilms on biomaterial surfaces were a major cause of failure in implant surgeries [1–6]. The biofilms are generally formed as a result of bacterial adhesion, growth, colony formation, extracellular polysaccharides, quorum sensing signals, and formation of nutrition channels. Biofilms can weaken the effect of antibiotic agents due to the presence of a wide variety of bacterial species and the barrier effect of the extracellular polysaccharide [7–12]. After the formation of biofilms, it becomes almost impossible to remove the matured biofilms from implanted devices in the human body. The only way of preventing sepsis is by retrieval of the device on which the biofilm was formed from the patient. To avoid this, it is necessary to inhibit biofilm formation during implantation of devices in advance. Prosthetic joint infection is divided into early infection (within three weeks after surgery) and late-onset infection (around three to eight weeks after surgery); this period is considered the incubation stage [13]. Moreover, infections associated with hip implants have been reported in the case of dental treatments [14,15]. Thus, the long-term antibacterial activity of the implant material is strongly desired. An ideal biomaterial surface with antibacterial activity can not only prevent the initial stages of infection such as bacterial adhesion, but also inhibit subsequent bacterial growth at later stages.

Recent studies on antibacterial surfaces have used various surface treatments with silver (Ag) species [16–23]. Ag is a well-known antibacterial agent and its effects on various bacteria have been studied extensively. Many researchers have reported that Ag can strongly influence various kinds of fungal and bacterial strains, including multidrug-resistant bacteria [23–27]. In addition to the efficacy of Ag for bacteria, copper, zinc (Zn), gallium, selenium and silicon have been recently used as antibacterial elements owing to their good antibacterial activity [23,28–35]. Among them, Zn is one of the most important trace elements in living organisms and an effective antibacterial element [30,31,36–42]. Therefore, we expected that the application of Zn for the surface modification of implant devices might offer antibacterial activity as effective as that of Ag.

Titanium (Ti) and its alloys are widely used as major implant materials owing to their excellent mechanical properties and biocompatibility [43]. Recent studies related to the bio-functionalization of a Ti surface has been widely reported and well-summarized elsewhere [44]. Among them, Micro-arc oxidation (MAO) is an electrochemical surface treatment technique performed in a specific electrolyte under high voltage. After MAO, the surface of the substrate metal is covered by a connective-porous oxide layer. The resultant oxide layer formed by MAO treatment contained additional elements that were contained in the electrolyte solution. Therefore, MAO treatment improved the osteogenic cell compatibility of titanium (Ti) when the electrolyte contained calcium and phosphate ions [45–51]. The biocompatibility of MAO coatings has been demonstrated by numerous in vitro and in vivo tests [52–57].

Several studies have focused on the incorporation of Zn onto a Ti surface by MAO treatment [58–64]. Hu et al. [59], Zhang et al. [61], Zhang et al. [62] and Du et al. [63] reported that Zn-incorporated TiO2 coatings showed good antibacterial activity against both *Escherichia coli* (*E. coli*) and *Staphylococcus aureus*. Moreover, Zhao et al. [60] reported that a Zn coating on a Ti surface showed bacteriostatic activity against *Streptococcus mutans*. These experiments provide adequate evidence to demonstrate that MAO-treated Ti in the electrolyte with Zn inhibited both bacterial adhesion as well as growth. In addition, Hu et al. [59], Zhao et al. [60] and Zhang et al. [64] reported that suitable amounts of Zn could promote adhesion, proliferation, and differentiation of osteogenic cells (e.g., rat bone marrow stem cells, MG63 cells and MC3T3-E1 cells). Thus, Zn plays a key role in the development of both antibacterial activity and osteogenic cell compatibility, and the incorporation of Zn by MAO treatment is a promising approach for surface treatment to achieve antibacterial activity in orthopedic and dental implants. Nevertheless, little attention has been given to the chemical and biological changes of Zn in the MAO coatings on their bio-function. Therefore, for the prevention of late-onset infections and the application of Zn to medical and dental implants, it is important to evaluate the long-term behavior of Zn in the body.

The purpose of this study was to investigate the long-term behavior of Zn-incorporated surface oxide produced by MAO for the prevention of late-onset infections. We investigated the time transient effect of Zn on chemical property, antibacterial activity, and osteogenic cell compatibility. The detailed chemical state of Zn in physiological saline was characterized by X-ray photoelectron spectroscopy (XPS). The change in antibacterial activity of the specimens before and after incubation for four weeks in saline was evaluated by the ISO method using *E. coli* as typical gram-negative facultative anaerobic bacteria. In addition, calcification by the cells on the specimens were evaluated at 14 and 28 days after seeding. In this study, we considered different time points and investigated changes in the chemical and biological properties of the Zn-incorporated MAO surface in a simulated body fluid environment.

## **2. Materials and Methods**

## *2.1. Specimen Preparation*

Commercially pure Ti (CP Ti; Rare Metallic, Tokyo, Japan) with grade 2 was used as a substrate material in this study. Two kinds of CP Ti disks with diameters 8 mm and 25 mm were obtained by cutting from a rod of CP Ti. The surfaces of these disks were mechanically polished using #150, #320, #600, and #800 grid SiC abrasive papers, followed by ultra-sonication in acetone and ethanol for 10 min. The disks were then stored in an auto-dry desiccator till further use. Each Ti disk was then fixed onto a polytetrafluoroethylene holder with an O-ring. The area in contact with the electrolyte was 39 mm<sup>2</sup> (7.0 mm in diameter) or 398 mm2 (22.5 mm in diameter). Details of the working electrode were as described earlier [65]. An AISI304 type stainless steel plate was used as a counter electrode. The base composition of the electrolyte for MAO treatment was 100 mM calcium glycerophosphate and 150 mM calcium acetate. Varying concentrations of ZnCl2 (0, 0.5, 1.0, and 2.5 mM) were added to the base electrolyte. After pouring the electrolyte into the electrochemical cell, both electrodes were connected to a DC power supply (PL-650-0.1, Matsusada Precision Inc., Shiga, Japan) and a positive voltage with a constant current density of 251 Am−<sup>2</sup> was applied for 10 min. The resultant voltage during the MAO treatment was 400 V at all conditions. After the MAO treatment, the surfaces were thoroughly washed in ultrapure water in order to remove any electrolyte solution remaining in the porous oxide layer. A major part of the Ti disk was MAO-treated; an annular untreated area was 0.5 mm from the margin. All surface characterization described below was performed in a MAO-treated area.

## *2.2. Incubation in Saline*

The specimens treated in the electrolyte that had various concentrations of Zn were immersed in physiological saline (0.9% NaCl) for 28 days. Saline is the simplest simulated body fluid, which is suitable for long-period measurement. Moreover, chloride ion (Cl−) is unexceptionally contained in all simulated body fluids. The specimens were fixed onto a polyethylene container to allow the release of Zn ions from the surface into physiological saline. Incubation was performed at 37 ◦C in a humidified chamber under constant shaking (100 rpm). Every 7th day, the pooled solution was changed with fresh physiological saline. After immersion, the immersed specimens were used further for surface characterization, evaluation of antibacterial activity, and quantification of Zn ions.

#### *2.3. Surface Characterization and Evaluation of Zinc Ion Release*

Scanning electron microscopy with energy dispersive X-ray spectrometry (SEM/EDS; S-3400NX, Hitachi High-Technologies Corp., Tokyo, Japan) was used to observe the surface morphology and composition before and after incubation. X-ray diffraction (XRD, BRUKER D8 DISCOVER, Bruker AXS KK, Yokohama, Japan) was performed to characterize the crystal structure of the specimens before and after incubation. The diffractometer was used with Cu-Kα radiation (40 keV, 40 mA). XPS was performed using a spectrometer (JPS-9010MC, JEOL, Tokyo, Japan) with Mg Kα X-ray source (energy: 1253.6 eV; acceleration voltage: 10 kV; current: 10 mA). The pressure of the measurement chamber was 1 <sup>×</sup> <sup>10</sup>−<sup>7</sup> Pa. All binding energies mentioned in this paper are relative to the Fermi level. The spectrometer was calibrated against Au 4f7/<sup>2</sup> of pure gold, Ag 3d5/<sup>2</sup> of pure silver, and Cu 2p3/<sup>2</sup> of pure copper. Spectra were obtained using an analysis area of 1 mmφ with a pass energy of 20 eV. The detection angle to the specimen surface was 90◦. The binding energies were calibrated with C 1s photoelectron energy region peak derived from contaminating carbon (285.0 eV). To calculate the integrated intensity of peaks, the background was subtracted from the measured spectrum according to Shirley's method [66]. The chemical state changes of Zn were determined by the modified Auger parameter (α') on the Wagner plot with those of typical Zn compounds. The α' of each specimen was calculated using the following Equation (1).

$$\mathbf{x}' = E\_K(\mathbf{Z}\mathbf{n}\,\mathbf{L}\_3\mathbf{M}\_{45}\mathbf{M}\_{45}) + E\_B(\mathbf{Z}\mathbf{n}\,\mathbf{2}\mathbf{p}\_{3/2})\tag{1}$$

Equation (1) contains both kinetic energy of Zn L3M45M45, *EK*(Zn L3M45M45) and binding energy of Zn 2p3/2, *EB*(Zn 2p3/2). The composition of the specimens were calculated according to a method described previously [67]. The photoionization cross-section of empirical data [68,69] and theoretically calculated data [70] were used for quantification. Inductively coupled plasma-atomic emission spectrometry (ICP-AES, ICPS-7000 ver. 2, Shimadzu Corp., Kyoto, Japan) was used to quantify Zn ion release from the Zn incorporated surface. The amounts of Zn ions released into the pooled solution were quantified each week.

#### *2.4. Evaluation of Antibacterial Activity*

The antibacterial activity was evaluated according to ISO 22196: 2007 method. To examine the antibacterial activity of the specimens, we employed *E. coli* (NBRC3972). The experiment was approved by the Pathogenic Organisms Safety Management Committee of the Tokyo Medical and Dental University (22012-025c). *E. coli* is a representative gram-negative rod-shaped bacterium commonly found in the gut. *E. coli* was cultured in Luria-Bertani (LB) broth (LB-Medium, MP Biomedicals, CA, USA) at 37 ◦C for 24 h. The optical density of the bacterial suspensions were measured at 600 nm using an ultraviolet-visible (UV–vis) spectrometer (V-550, JASCO, Tokyo, Japan) and diluted to obtain concentrations of 4.9 <sup>×</sup> 106 colony-forming units (CFUs) mL<sup>−</sup>1. Prior to the antibacterial activity testing, all specimens were sterilized with 70% ethanol, washed with distilled water, and dried. Drops of bacterial suspension were added on all specimens, which were subsequently covered with a sterilized plastic film and incubated at 37 ◦C for 24 h (*n* = 3). Bacterial cells were then collected from all the incubated specimens; the obtained suspensions were diluted, pipetted onto nutrient agar plates, and incubated overnight at 37 ◦C. The number of viable bacteria was determined by counting the number of colonies formed.

## *2.5. Calcification by Osteogenic Cells*

As described in our previous work [71], MC3T3-E1 cells (RIKEN BioResource Center, Tsukuba, Japan) were maintained in a cell culture medium: alpha modification of Eagle's minimum essential medium (α-MEM; GIBCO, Grand island, CA, USA) supplemented with 10% fetal bovine serum (FBS; GIBCO), 10 U mL−<sup>1</sup> penicillin, 100 mg mL−<sup>1</sup> streptomycin, and 0.25 mg mL−<sup>1</sup> amphotericin B (GIBCO). All specimens were sterilized in 70% ethanol for 20 min and thoroughly rinsed with deionized water before in vitro testing. The cells were seeded on to the sterilized specimens with an approximate initial density of 10,000 cells cm−2. As a control, cells were also seeded on MAO-treated Ti without Zn. The cells were incubated at 37 ◦C in a fully humidified atmosphere under 5% CO2. For induction of osteogenic differentiation, a cell culture medium supplemented with 2 mM β-glycerophosphate (Calbiochem, Darmstadt, Germany) and 50 mg mL−<sup>1</sup> L-ascorbic acid (Wako Pure Chemical Industries, Osaka, Japan) was used when 100% confluence was reached for all the specimens. The medium used for induction was named as the differentiation-inducing medium. The differentiation-inducing medium was changed every three days.

After the seeding for the 1st, 2nd, and 3rd day, the cells attached to each specimen were harvested with trypsin/EDTA, re-suspended in a cell culture medium, and transferred to 96-well microplates. Then, the Cell Counting Kit-8 assay (CCK-8, Dojindo Laboratories, Kumamoto, Japan) was added to each 96-well microplate, which contained cell suspension, and the reaction was continued for 4 h at 37 ◦C. The absorbance of the samples was measured at 450 nm using a microplate reader (ChroMate Microplate Reader, Awareness Technology, Palm, FL, USA). The reference wavelength was set at 630 nm. In this evaluation, the tissue culture polystyrene (TCPS) was used as a control specimen for quantification of the number of attached cells on each specimen. The cells that were attached to the TCPS were harvested by treatment with Trypsin/EDTA followed by re-suspension of cells in the cell culture medium. Cell suspensions were diluted serially. The number of cells in the cell suspension harvested from TCPS were counted by Trypan blue (Trypan Blue Stain 0.4%; Gibco) using

a hemocytometer. Then, standard curves for cell number calibration and for light absorbance were constructed and were used for quantification of the number of attached cells on each specimen.

The calcification of MC3T3-E1 cells on each specimen was evaluated by the calcified deposits by a color change reaction after alizarin red S staining. After the removal of media from all the specimens, cells were rinsed three times with phosphate-buffered saline. Cells were then fixed with 4% formalin for 1 h and rinsed three times with ultra-pure water. Each specimen was stained with 1% alizarin red S solution (adjusted to pH 4.2 with ammonium hydroxide) at room temperature for 30 min. After removing the alizarin red S solution, cells were repeatedly rinsed with ultrapure water. When the specimens were fully dry, the surface of each specimen was studied using an optical microscope (OLYMPUS SZX12, Olympus, Tokyo, Japan).

## *2.6. Statistical Analysis*

Data obtained by surface characterization were calculated from three independent specimens. The results of biological tests are shown from at least two independent tests. Each test was performed using at least three independent specimens. All values are presented as means ± SD, and commercial statistical software KaleidaGraph (Version 4.1.1, Synergy Software, Reading, PA) was used for statistical analysis. One-way analysis of variance (ANOVA) was used following multiple comparisons with Student–Newman–Keuls method to assess the data, and *P* < 0.05 was considered to indicate statistical significance.

## **3. Results**

## *3.1. Surface Characterization and Evaluation of Zinc Ion Release*

Figure 1 shows SEM images of the MAO-treated Ti surface treated in the electrolyte with 2.5 mM Zn before (A) and after (B) incubation in saline. Typical connective porous morphology after MAO treatment was observed. There was no difference in the surface morphology among the specimens before and after incubation in saline for 28 days according to our SEM data under the indicated magnification. In addition, the pore size on the specimens before incubation was 5.3 ± 2.3 μm and on the specimens after incubation during 28 days was 4.8 ± 2.1 μm. There were no significant differences in the pore size among the specimens before and after incubation.

**Figure 1.** Scanning electron microscopy (SEM) images of specimens before (**A**) and after (**B**) incubation in saline during 28 days. These images show the connective porous oxide layer on a Titanium (Ti) surface formed by micro-arc oxidation (MAO) treatment using Zinc-containing electrolyte.

XRD spectra obtained from the MAO-treated Ti surface treated in the electrolyte with 2.5-mM Zn before and after incubation has been shown in Figure 2. The peaks originating from α-Ti and anatase TiO2 were detected from each specimen. Moreover, the peaks originating from Zn were undetected. There was no difference in the crystal structure among the specimens before and after incubation in saline for 28 days.

**Figure 2.** X-ray diffraction (XRD) spectra of the Zn-incorporated specimens before and after incubation in saline during 28 days.

XPS survey spectra of the MAO-treated Ti surface in the electrolyte with 2.5 mM Zn before and after incubation in saline has been shown in Figure 3A. In addition to peaks originating from C, O, P, Ca, and Ti, those originating from Zn and Na were detected. Narrow scan spectra around Zn 2p electron energy region is shown in Figure 3B. The binding energies of Zn 2p3/<sup>2</sup> peaks obtained from all specimens were 1022.2–1022.6 eV, indicating that Zn existed as Zn2+. Phosphorus was found to exist as phosphate species and calcium existed as a divalent ion, because binding energies of the corresponding peaks of P 2p and Ca 2p3/<sup>2</sup> were 134.1 ± 0.2 eV and 347.8 ± 0.2 eV, respectively. The binding energy of Ti 2p3/<sup>2</sup> peak was 459.0 ± 0.1 eV, indicating that Ti existed as TiO2.

**Figure 3.** X-ray photoelectron spectroscopy (XPS) spectra of survey (**A**) and Zn 2p photoelectron energy region (**B**) from specimens immersed in saline for 0–28 days.

Figure 4 depicts the Wagner plot of Zn in the oxide layer incubated for 0–28 days based on Zn 2p3/<sup>2</sup> peaks and Zn L3M45M45 Auger peaks and reference plots from Zn2<sup>+</sup> compounds as reported by previous studies [72–75]. The chemical state of an element was determined by comparison of its binding energy with modified Auger parameter values on the Wagner plot. According to the Wagner plot of Zn, the binding energy decreased, and the Auger parameter increased with incubation time. Moreover, since the value of the Auger parameter finally converged to 2010.2 eV, it became clear that the chemical state of Zn incorporated in titanium dioxide by MAO approached to that of ZnO with increasing incubation times.

**Figure 4.** Wagner plot of Zn in the oxide layer incubated for 0–28 days based on Zn 2p3/<sup>2</sup> photoelectron peaks and Zn L3M45M45 Auger peaks (*n* = 3). Each parameter of Zn2<sup>+</sup> compounds were plotted according to previous studies [72–75].

Figure 5 shows the change in the Zn concentration in the oxide layer at each incubation time, as determined by XPS. The concentration of Zn was relatively small even before the incubation; thus, the amount of Zn incorporated during MAO treatment was small. The amount of Zn dramatically decreased to about 0.8% by the incubation in saline from 0 to 7 days and was maintained for 28 days.

**Figure 5.** Change in the concentration of Zn in the oxide layer with incubation time determined by XPS (*n* = 3).

Figure 6 shows the amount of Zn ions released from the oxide layer into saline, as determined by ICP-AES. The amount of Zn ions released within the first 7 days was the highest and thereafter Zn was not detected. In summary, there was no observed release of Zn ions after 7 days.

**Figure 6.** The amount of Zn ions released into saline every week from Ti MAO-treated in the electrolyte containing 2.5-mM Zn (*n* = 3).

#### *3.2. Evaluation of Antibacterial Activity*

Figure 7 shows the normalized CFU count for *E. coli* or antibacterial effects of MAO-treated Ti in the electrolyte containing various concentrations of Zn. The vertical axis represents the normalized bacterial number, defined as the bacterial number on each of the specimens divided by that of the control, MAO-treated Ti in the electrolyte without Zn. The normalized bacterial number that was below 1 (shown as dashed line in the figure) implied that the tested specimens had antibacterial activity. Before incubation, there was no significant difference in *E. coli* counts among the specimens. On the other hand, after incubation for 28 days, antibacterial effects against *E. coli* were observed. This effect was independent of the concentration of Zn in the electrolyte for MAO treatment, indicating that antibacterial activity of Zn developed after incubation in saline.

**Figure 7.** Antibacterial effects of Ti MAO-treated in the electrolyte containing various concentrations of Zn. Data are shown as the mean ± SD. \*: Significant difference between specimens before and after incubation, \*\*: Significant difference against the specimen without Zn (*n* ≥ 3, *P* < 0.05).

## *3.3. Calcification by Osteogenic Cells*

Figure 8 shows the number of MC3T3-E1 cells on MAO-treated Ti in the electrolyte containing various concentrations of Zn with varying culture times. The number of viable cells on all specimens increased with culture time. The Zn concentrations in the electrolyte did not differ significantly among specimens.

**Figure 8.** Number of MC3T3-E1 cells on Ti MAO-treated in the electrolyte containing various concentrations of Zn with culture time. Data are shown as the mean ±SD. n.s.: No significant difference (*n* ≥ 5, *P* < 0.05).

Figure 9 shows photographs of alizarin red S staining of calcified deposits by MC3T3-E1 cells on each specimen cultured for 14 and 28 days. In case of MAO-treated specimens, we observed that only the central part (7 mm in diameter) was covered with the oxide layer formed by MAO treatment while the margins (0.5 mm) were not treated. Calcified cells were observed in all the specimens. However, untreated Ti showed bare metal-colored parts that indicated the presence of incompletely calcified cells. On the other hand, all MAO-treated specimens showed completely stained surfaces. There was no segregation seen for calcified and non-calcified areas on the MAO-treated surface. In addition, there were no visual differences among the MAO-treated specimens with and without Zn.


**Figure 9.** Color scale stained by alizarin red S of calcified deposits by MC3T3-E1 cells on each specimen cultured for 14 and 28 days. Scale bar represents 2 mm.

## **4. Discussion**

The morphology and crystal structure of the porous oxide layer formed by MAO treatment did not show any changes during incubation in saline (Figures 1 and 2); this property is highly advantageous for implant surfaces.

The peaks originating from Zn were not detected from the specimens by XRD. The amount of Zn incorporated during MAO treatment was 3.5%, and decreased with incubation (Figure 5). Therefore, we believe that XRD peaks originating from Zn were relatively small compared with Ti, which was not detected by XRD.

The chemical state and concentration of Zn in the oxide layer changed during incubation in saline (Figures 3–5). In addition, the Zn ions were released into saline during the 7-day incubation period; thereafter, no Zn ions were eluted (Figure 6). These phenomena indicate that Zn incorporated into specimens would be released only during the initial stages of implantation in vivo. The concentration of Zn in the oxide layer decreased with the release of Zn ions during the first 7 days and was maintained at 0.8 at % in the subsequent period. In other words, despite the presence of Zn in the oxide layer, no Zn ions were released from the surface from 7 to 28 days. This indicated that the released Zn ion was adsorbed on the surface of the oxide layer to form zinc oxide as an insoluble corrosion product of Zn. The insoluble corrosion product of Zn was converted to ZnO to achieve chemical stabilization of Zn in saline. In other words, Zn incorporated into the oxide layer by MAO treatment changed the chemical state by interaction with saline and was finally stabilized by conversion into ZnO. Moreover, ZnO was already generated on the surface of oxide layer in the 14-day incubation (Figure 4). Therefore, it was considered that Zn ions did not elute during the 14–28 days period due to both the conversion into ZnO and stabilization in saline.

The specimens incubated in saline exhibited antibacterial activity against *E. coli*, while the specimens before incubation did not (Figure 7). Ti incorporated with Zn in the surface oxide layer formed by MAO treatment exhibited late-onset antibacterial activity against *E. coli*. The release of Zn ions were observed only during the initial stages of the incubation process (0–7 days). This indicated that Zn ions itself did not have any antibacterial activity. Du et al. [76] reported that the minimum inhibitory concentration (MIC) for Zn ions against *E. coli* was 768 <sup>μ</sup>g·mL<sup>−</sup>1. The maximum concentration of Zn ions released from the specimens in this study was smaller than the previously reported MIC of Zn ions. For this reason, development of antibacterial activity by other factors needs to be considered. Prasanna et al. [77] reported that ZnO could produce reactive oxygen species (ROS) in the dark, owing to singly ionized oxygen vacancy in the crystal lattice of ZnO, and that generation of ROS contributed to its antibacterial properties. This study clearly revealed that Zn was finally converted to ZnO and was stabilized by interactions in saline. Therefore, the development of the late-onset antibacterial activity of Zn was due to the generation of ZnO in saline. In other words, this study proved that the chemical state of Zn incorporated in the oxide layer played a key role in the development of its antibacterial activity. Therefore, we believed that antibacterial activity was developed when ZnO was generated. According to Figure 4, we considered that the specimens incubated in saline during 14–21 days could exhibit the antibacterial activity, while the specimens incubated in saline during 0–7 days may not affect to bacteria. On the other hand, it is believed that bacteria interact with titanium dioxide through surface siderophores, and the oxides provide a template for biofilm formation [78,79]. For this reason, it is expected that bacteria can be captured by the interaction between the porous titanium dioxide and the siderophore on the cell surface, and killed by ZnO. We believe that Zn-incorporation for a Ti surface by MAO treatment may be suitable for realizing both bacterial capture ability and antibacterial activity, namely a 'trap-killing' system [80].

The MAO-treated specimens containing Zn showed no toxic effects on proliferation of osteogenic cells (Figure 8). Moreover, the calcification of cells on MAO-treated specimens with Zn were at the same level as that on the specimen without Zn. Therefore, the existence of a slight amount of Zn did not affect the osteogenic property of cells and proliferation was normal. The porous oxide layer consisted not only of titanium oxide but also incorporated Ca, P, and Zn (Figure 3). Relatively large amounts of Ca and P were contained in the oxide layer of MAO-treated Ti-based alloy at 7.6 and 10.0 at %, respectively [81]. Ca and P are the main mineral components in bone tissues and are present as calcium–phosphate compounds. The presence of Ca and P in the porous oxide layer formed by MAO treatment enhances the activity of osteogenic cells [47]. In our previous study, MAO-treated specimens containing a small amount of Ag showed osteogenic cell compatibility, which accelerated the calcification process by MC3T3-E1 cells without any cytotoxicity [22]. Therefore, we predicted that Zn-containing specimens in this study might have the same calcification ability.

In addition, a specific antibacterial activity of Zn, which was incorporated in the oxide layer was observed. Zn-incorporated oxide layer formed by MAO exhibited late-onset antibacterial activity and osteogenic cell compatibility, namely dual function. The amount of Zn ions released during 7 days incubation (0.02 ppm) was much less than IC50 for MC3T3-E1 cells reported by Yamamoto et al. [82]. Moreover, our results indicated that the small amount of ZnO generation did not affect osteogenic cells. In other words, the time transient effect of a suitable amount of Zn was effective for only bacteria. Therefore, MAO treatment of Ti using Zn made it possible to exhibit both late-onset antibacterial activity and osteogenic cell compatibility on Zn-containing TiO2 layer. In addition, our study proposed the importance of the chemical state of Zn in the oxide layer for the development of antibacterial properties of Zn. Our results have suggested that due to the time transient effect of Zn in TiO2, its chemical and biological properties could be easily changed and this property could be adapted for achieving changes in simulated body fluids (Figure 10). In particular, we suggest the importance of the chemical state changes of Zn in the various environments to control its chemical and biological properties. We hope that in future, that the outcome of the present study would be useful in the design of bio-functional implants related to Zn and contribute to the development of antibacterial biomaterials.

**Figure 10.** The schematic illustration of time transient effect of Zn in the porous titanium dioxide layer on chemical and biological properties.

## **5. Conclusions**

This study investigated the time transient effect of Zn in titanium dioxide with incubation in saline. Although the morphology of the porous oxide layer formed by MAO treatment did not change, the chemical state of Zn in the oxide layer changed to ZnO during incubation in saline. Moreover, the antibacterial activity of the specimen with Zn developed in the late stages after the incubation process in saline for 28 days. We propose that the generation of ZnO in the oxide layer played a key role in development of antibacterial activity. In addition, the Zn-incorporated specimen did not exhibit any cytotoxicity in MC3T3-E1 cells and did not hinder the proliferation and calcification of cells. Therefore, MAO treatment of Ti using Zn is suitable for realizing both late-onset antibacterial activity and osteogenic cell compatibility.

**Author Contributions:** Conceptualization, M.S., Y.T., A.N. and T.H.; formal analysis, M.S., Y.T., K.N. and P.C.; investigation, M.S., K.N., R.Y. and M.A.; methodology, M.S., K.N., P.C. and H.D.; project administration, Y.T., A.N. and T.H.; supervision, A.N. and T.H.; validation, M.S., Y.T., K.N. and R.Y.; writing – original draft, M.S.; writing – review and editing, T.H.

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

**Acknowledgments:** This study was supported by the Research Center for Biomedical Engineering, Tokyo Medical and Dental University, Project "Creation of Life Innovation Materials for Interdisciplinary and International Researcher Development" and project "Cooperative project amount medicine, dentistry, and engineering for medical innovation-Construction of creative scientific research of the viable material via integration of biology and engineering" by the Ministry of Education, Culture, Sports, Science and Engineering, Japan.

**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*

## **Icariin**/**Aspirin Composite Coating on TiO2 Nanotubes Surface Induce Immunomodulatory E**ff**ect of Macrophage and Improve Osteoblast Activity**

## **Aobo Ma** †**, Yapeng You** †**, Bo Chen, Wanmeng Wang, Jialin Liu, Hui Qi, Yunkai Liang, Ying Li \* and Changyi Li \***

School of Dentistry, Tianjin Medical University, Tianjin 300070, China; maaobo@tmu.edu.cn (A.M.); youyapeng@tmu.edu.cn (Y.Y.); chenbo@tmu.edu.cn (B.C.); wangwanmeng@tmu.edu.cn (W.W.); liujialin@tmu.edu.cn (J.L.); qihui@tmu.edu.cn (H.Q.); liangyunkai@tmu.edu.cn (Y.L.)


Received: 24 March 2020; Accepted: 22 April 2020; Published: 24 April 2020

**Abstract:** Surface coating modification of titanium-based alloys is an efficient way to accelerate early osseointegration in dental implant fields. Icariin (ICA) is a traditional Chinese medicine that has bone activating functions, while aspirin (ASP) is a classical non-steroidal anti-inflammatory drug with good antipyretic and analgesic capabilities. Moreover, poly(lactic–co–glycolic acid) (PLGA) has attracted great attention due to its excellent biocompatibility and biodegradability. We superimposed an ASP/PLGA coating onto ICA loaded TiO2 nanotubes structure so as to establish an icariin/aspirin composite coating on TiO2 nanotubes surface. Scanning electron microscopy, X-ray photoelectron spectroscopy, a contact angle test and a drug release test confirmed the successful preparation of the NT–ICA–ASP/PLGA substrate, with a sustained release pattern of both ICA and ASP. Compared to those cultured on the Ti surface, macrophage cells on the NT-ICA-ASP/PLGA substrate displayed decreased M1 proinflammatory and enhanced M2 proregenerative genes and proteins expression, which implied activated immunomodulatory effect. Moreover, when cultured with conditioned medium from macrophages, osteoblast cells on the NT-ICA-ASP/PLGA substrate revealed improved cell proliferation, adhesion and osteogenic genes and proteins expression, compared with those on the Ti surface. The abovementioned results suggest that the established NT-ICA-ASP/PLGA substrate is a promising candidate for functionalized coating material in Ti implant surface modification.

**Keywords:** icariin; aspirin; composite coating; TiO2 nanotubes; immunomodulatory effect; macrophage; osteoblast activity

## **1. Introduction**

Dental implant dentures are the preferred treatment option for patients with dentition defect and edentulous jaws [1]. Titanium and titanium—based alloys have become the main steam material for dental implants due to their excellent mechanical properties and biocompatibility [2]. However, the biological inertness of pure titanium impairs the rapid bonding between dental implant and bone tissue, even causes implant failure [3]. Therefore, surface modification of Ti-based implant materials has become the main way to enhance cell activity, exert antibacterial effect and so as to promote osseointegration [4]. Various strategies have been employed for improving surface characteristics of dental implant, including physical, chemical and biological methods [5].

Among them, the application of protein growth factor on implant surface has been proved to be an effective way for enhancing osteogenesis and could accelerate osseointegration at the implant–tissue

interface, such as bone morphogenetic protein–2 [6–8]. However, the lack of chemical stability of growth factor proteins hinders their wide application in implant surface modification [8,9]. In addition, inflammation reaction generated during the drilling, screwing and inserting processes of implant surgery may hamper bonding strength between implant and surrounding bone tissue [10]. In order to overcome the drawbacks, it becomes imperative to develop a new surface modification method, which could load osteogenic drug with chemical stability and subside inflammation accompanied with implant surgery at the same time.

Epimedium is widely used in traditional Chinese medicine and has the kidney invigorating and bone activating functions [11]. Icariin (ICA), an inexpensive and chemically stable monomer components of Epimedium [12,13], retains the osteogenic function of Epimedium and has antiatherosclerotic, anti-inflammatory and anticancer capabilities [14]. At the same time, icariin can promote the proliferation activity of osteoblasts and inhibit the formation of osteoclasts. Therefore, it is widely used in the treatment of osteoporosis [15]. In addition, icariin is also a promising osteoinductive compound for bone tissue engineering [11,16]. We previously reported that the icariin-functionalized coating on TiO2 nanotubes surface promote osteoblast function and accelerate osseointegration [17]. Moreover, recently the anti-inflammatory [18,19] and immunoregulatory [20] functions of ICA has also been raised. Therefore, ICA becomes an ideal candidate drug for Ti-based implant surface modification to regulate immune response and promote bone formation at the same time.

Aspirin (ASP) is a well-established non-steroidal anti-inflammatory drug (NSAID), with chemical stability and an inexpensive price. In addition to its long-term application as an antipyretic and analgesic for postoperative pain relief [21], ASP is also reported to substantially improve immunomodulatory function, manifested as regulatory T cells upregulation and Th17 cells downregulation in vitro [22]. In addition, systemic use of ASP revealed enhanced bone regeneration in vivo [23]. Moreover, some scholars have recently demonstrated that low doses of ASP can regulate the balance between bone resorption and bone formation in osteoporosis caused by ovariectomy [24–26]. Evidence from both in vitro cell culture tests and in vivo animal studies also revealed that ASP has a protective effect on bones by promoting the proliferation of osteoblast precursor cells and differentiation of osteoblasts [27]. Therefore, in the present work, we further superimposed ASP on the ICA-loaded TiO2 nanotubes surface, in order to simultaneously resolve acute inflammation after implant insertion and improve osteogenesis.

Due to the lack of immobilization sites for drug molecules on the pristine Ti surface, we fabricated TiO2 nanotubes on Ti surface as the reservoir for drug storage and delivery. Prepared by anodization, TiO2 nanotubes are widely used as a surface modification method for dental implants [28]. It is well-accepted that nanostructured morphology can promote osteoblast adhesion, proliferation, differentiation activity and accelerate osseointegration around the implant [29]. However, since the end of nanotubes are in an open state, loading drug on TiO2 nanotubes surface may cause undesired early burst release of the drug [30]. To solve this problem, it is necessary to find a coating material combined with nanotube architecture so as to endow the modified surface with a controlled drug release platform.

In order to effectively overcome the initial fast drug release of the simple TiO2 nanotubes surface, various coatings were applied to acquire sustained drug release capacity on the nanotubes surface [31,32]. Among them, poly(lactic-co-glycolic acid), often shortened to PLGA, is one of the important materials for preparing controlled drug release coatings. A large amount of literature have reported different types of PLGA-based drug delivery systems (DDSs), such as nanoparticles, microspheres, implants, nanogels, nanofibers, rods, films, etc. [33,34]. The PLGA controlled release coating can load a variety of substances from hydrophilic to lipophilic drugs, from micromolecule to macromolecule and from single molecules to multiple molecules [35].

Therefore, in the present study, TiO2 nanotubes structure were first prepared on the Ti surface and then loaded with ICA. Afterwards, the ASP/PLGA coating is employed to cover the pre-existed ICA loaded nanotubes surface. Thus, we constructed the icariin/aspirin composite coating on TiO2 nanotubes surface, abbreviated as NT–ICA–ASP/PLGA substrate. Scanning electron microscope (SEM), X-ray photoelectron spectroscopy (XPS) and contact angle test were utilized for characterization of the physical and chemical features of the established icariin/aspirin composite coating on TiO2 nanotubes surface. While the drug release tests were used to evaluate the release profiles of ICA and ASP, respectively.

Previously, researchers have made great effort on titanium surface modification to improve osteogenic properties in mesenchymal stem cells or osteoblast cells. However, the discrepancies of biological responses between in vitro and in vivo experiments exist when evaluating implant materials so far. It implies that the complexity of the in vivo environment has not been fully understood and some important factors, including immunomodulatory factors, are often ignored [36,37]. Among all the immune cells, macrophages play a central role in inflammation reaction [38,39]. Macrophage cells were also one of the cell types, which firstly contact with implant materials after insertion [40]. Therefore, in the present work, we intend to simultaneously explore the immunomodulatory function and osteogenic effect of the icariin/aspirin composite coating on TiO2 nanotubes surface. At first, we examined the immunomodulatory effect of the established NT–ICA–ASP/PLGA substrate using macrophage cells. Then, we employed the indirect coculture system including macrophage conditioned media (CM) and osteoblast cells to evaluate the osteogenic function of the constructed NT–ICA–ASP/PLGA substrate.

Our hypothesis is that icariin/aspirin composite coating on the TiO2 nanotubes surface could improve osteogenic effect of osteoblast cells through regulating the polarized status of macrophages. By exerting immunomodulatory function of macrophages at the interface between implant and surround bone tissue, an immune microenvironment that favors bone formation is created. The combined effects of icariin and aspirin may synergistically accelerate osseointegration of the dental implant.

## **2. Materials and Methods**

## *2.1. Chemicals and Reagents*

ICA and ASP was provided by Sigma-Aldrich (St. Louis, MO, USA). PLGA (Resomer RG 503, lactide:glycolide 50:50, ester terminated, Mw 24,000–38,000) was provided by Sigma-Aldrich Phosphate buffered saline (PBS) was obtained from Solarbio, Inc. (Beijing, China). Ammonium fluoride (NH4F) and ethylene glycol (EG) were provided by Tianjin FengChuan Chemical Reagent Co., Ltd (Tianjin, China).

## *2.2. Specimen Preparation and Surface Characterization*

## 2.2.1. TiO2 Nanotubes Fabrication

To prepare the TiO2 nanotubes surface, titanium (Ti) discs with diameter of 15 mm and thickness of 1 mm were fabricated from commercially produced titanium plate (Grade 2, ASTM F67 unalloyed Ti; 99.7% Ti; 0.14% O; 0.09% Fe; 0.04% C; 0.02% N; 0.008% H; 0.002% other elements), which were purchased from Baoji Titanium Industry (Baoji, China). Ti discs were polished by 320, 800, 1500 and 2000 grit sand paper in sequence, ultrasonically washed with acetone, ethanol and deionized water in turn for 10 min, and then dried in air. The solvent of the electrolyte is ethylene glycol (EG), the solute is 0.16 mol/L NH4F, with 10% deionized water by volume added. The TiO2 nanotubes discs were fabricated under a voltage of 40 V for 1 h by high-voltage DC power supply (Dongwen High Voltage Power Supply Factory, Tianjin, China). After reaction, discs were ultrasonically cleaned by the EG solution. Later, the discs were rinsed with deionized water for 3 times, and left to dry naturally to obtain the TiO2 nanotubes (NT) surface.

## 2.2.2. Construction of Icariin/Aspirin Composite Coating on the TiO2 Nanotubes Surface

To fabricate the NT–ICA surface, ICA was dissolved in anhydrous methanol, with the concentration adjusted to 1.15 mg/mL, then the solution was mixed by vortex oscillators for 2 min. The NT slices were placed in the wells of a 24–well plate, one well for each slice, then 1 mL prefabricated ICA solution was added to each well, sealed with parafilm membrane and left to stand for 24 h at 4 ◦C. After that, the excess liquid was sucked out and air-dried to obtain the NT-ICA surface.

To construct the Ti-PLGA and NT-ICA-PLGA surfaces, we dissolved 30 mg of PLGA powder in 1 mL acetone solution, mixed thoroughly by vortex. We spread 100 μL mixed solution evenly on the surface of the Ti and the NT–ICA slices, dried naturally and repeated four times to obtain the Ti–PLGA and NT-ICA-PLGA surfaces, respectively.

To establish the NT–ASP/PLGA and NT–ICA–ASP/PLGA substrates, we took 30 mg PLGA and 10 mg ASP powder into a 15 mL centrifuge tube, add 1 mL acetone solution, and mix thoroughly by vortex. Then apply 100 μL of the mixed solution of ASP, PLGA and acetone to the surface of the NT and NT–ICA slices and left to dry naturally. Repeat the coating step four times to obtain the NT–ASP/PLGA and NT–ICA–ASP/PLGA substrates, respectively.

## 2.2.3. Surface Characterization

Observe the macro surface morphology of Ti, Ti–PLGA, NT, NT–ICA, NT–ICA-PLGA, NT–ASP/PLGA and NT-ICA-ASP/PLGA surfaces with stereomicroscope (S9i; Leica Microsystem, Tokyo, Japan). The micro surface topography of various surfaces was observed using a scanning electron microscope (SEM, Zeiss Merlin Compact; ZEISS, Jena, Germany), the working distance was 22.4–27.8 mm. X-ray photoelectron spectroscopy (XPS, AXIS Nova; Kratos Analytical, Manchester, UK) was used for analyzing the elemental composition of each sample's constituents. At room temperature, the wettability of the sample surfaces were evaluated by contact angle measurement using deionized water. The droplet volume used was 2 μL, and a contact angle goniometer (JGW-360A, Chongda Intelligent Technology, Xiamen, China) was used to acquire droplet images. Values of initial contact angle were analyzed using image analysis software (version 1.0, Anglem, Chongda Intelligent Technology, Xiamen, China). Three different positions were measured for each sample, and values were determined as mean ± standard deviation, *n* = 3.

## *2.3. Icariin and Aspirin Drug Release Amount Measurement*

The drug release profiles of two drugs ICA and ASP were detected by the high performance liquid chromatography system (HPLC; 1100 Series; Agilent Technologies, Santa Clara, CA, USA). The MS 105 electronic balance (Mettler-Toledo., Greifensee, Zurich, Switzerland) of 0.01 mg accuracy was used to accurately weigh the ICA and ASP reference substance. Anhydrous methanol was used to prepare the standard solution used in the standard curve by half dilution method. The standard solution concentrations of ICA and ASP are 0, 0.5, 1, 2, 4, 8 and 16 μg/mL and 0, 2.5, 5, 10, 20, 40 and 80 μg/mL, respectively. The samples were immersed in PBS (pH 7.4) at 37 ◦C for drug release experiments. Various samples were immersed in 10 mL PBS and incubated for 24 h, and then the solution was daily collected for analysis by the HPLC system. After that, fresh PBS solution was added to replace the extracted solution. The test periods for ICA and ASP lasted for 30 and 8 days, respectively. The total amount of drug loaded on the sample was calculated according to the accumulated amount of drug released until the end of the test period used in this experiment. The drug cumulative release percentage was calculated by dividing the accumulated amount of released drug at each time point by the total amount of drug loaded. Data are collected from three separate experiments and expressed as mean ± SD (*n* = 3).

#### *2.4. Behaviors of RAW 264.7 Cells Cultured on Di*ff*erent Surfaces*

### 2.4.1. Cell Culture of Macrophage Cells

All samples were sterilized by irradiation with 25 kGy Cobalt 60 for 30 min and placed in 24–well plate. RAW 264.7 macrophage cells (American Type Culture Collection, ATCC, Manassas, VA, USA) were cultured in Dulbecco's modified eagle medium (DMEM, Gibco, Carlsbad, CA, USA) supplemented

with 10% fetal bovine serum (FBS, Gibco, Carlsbad, CA, USA) and 1% (*v*/*v*) streptomycin/penicillin (Gibco, Carlsbad, CA, USA) in a 5% CO2 incubator (Thermo Fisher Scientific, Waltham, MA, USA) at 37 ◦C. Culture medium was daily changed and cells were stimulated with Lipopolysaccharides (LPS; Sigma Aldrich, St Louis, MO, USA) at the concentration of 1 μg/mL for the first three days. All samples were sterilized by irradiation with 25 kGy Cobalt 60 for 30 min.

## 2.4.2. Proinflammatory (M1) and Proregenerative (M2) Marker Gene Expression

The gene expression levels of proinflammatory (M1) and proregenerative (M2) markers in RAW 264.7 macrophage cells cultured on different substrates were analyzed by quantitative real-time polymerase chain reaction (qPCR). RAW 264.7 macrophage cells were seeded on all surfaces at a density of 1 <sup>×</sup> 10<sup>4</sup> cells/well for 3 days. Then, total RNA was extracted by TRIzol (Thermo Fisher Scientific, Waltham, MA, USA). The RNA concentration and purity were measured at 260 nm by Nanodrop spectrophotometer (NanoDrop Technologies, Wilmington, DE, USA). Mouse mRNA encoding genes for tumor necrosis factor-alpha (TNF–a), interleukin-1β (IL–1β), transforming growth factor-beta (TGF–β) and heme oxygenase-1 (HO–1) were selected, glyceraldehyde–3–phosphate dehydrogenase (GAPDH) was amplified in parallel with the target genes and used as an internal control. Reverse transcription and qPCR were performed. Then relative gene expression levels were determined using the relative threshold cycle (CT) method and reported as 2−ΔΔCt. Primers used for both the target genes and the housekeeping gene were shown in Table 1. For every interested gene, the cDNA amplification was performed in triplicate by a single sample. The results were from three independent tests.


**Table 1.** Primers for proinflammatory (M1) and proregenerative (M2) marker genes.

## 2.4.3. Enzyme-Linked Immunosorbent (ELISA) Assay

The proinflammatory (M1) and proregenerative (M2) marker protein expression levels in RAW 264.7 macrophage cells incubated with all kinds of samples were further measured by enzyme-linked immunosorbent (ELISA) assay. RAW 264.7 macrophage cells were seeded on different surfaces at a density of 1 <sup>×</sup> <sup>10</sup><sup>4</sup> cells/well for 3 days. The supernatants were then collected and measured immediately. The concentrations of proteins including tumor necrosis factor-alpha (TNF-a), interleukin–1β (IL–1β), transforming growth factor-beta (TGF–β) and heme oxygenase–1 (HO–1) were determined by ELISA kits (ImmunoWay Biotechnology, Plano, TX, USA). Three different slices were measured for each group. Three independent experiments were repeated.

## *2.5. Behaviors of MC3T3-E1 Cells on Various Surfaces in Conditioned Medium (CM)*

## 2.5.1. Collection and Preparation of CM

RAW 264.7 macrophages were seeded on various surfaces at a density of 1 <sup>×</sup> 10<sup>4</sup> cells/well for 14 days. Cells were stimulated with LPS (1 μg/mL) for the first three days, and culture medium supernatant were collected daily and changed with new medium. The collected supernatant was centrifuged at 1500 rpm for 15 min, and then filtered through a 0.22 μm filter. Then, the filtered supernatant was mixed with DMEM containing 10% FBS and 1% penicillin/streptomycin at a volume ratio of 1:1 to obtain the conditioned medium (CM), and stored at 4 ◦C for later use.

## 2.5.2. Cell Culture of Osteoblast Cells

MC3T3-E1 preosteoblast cells (ATCC, Manassas, VA, USA) were first cultured in the DMEM supplemented with 10% FBS and 1% (*v*/*v*) streptomycin/penicillin in a 5% CO2 incubator at 37 ◦C. After 4 h of inoculation, cells were seeded on different surfaces at a density of 1 <sup>×</sup> <sup>10</sup><sup>4</sup> cells/well and the CM prepared in 2.5.1 was used for cell culture. Various samples were disinfected through irradiation with 25 kGy Cobalt 60 for 30 min.

## 2.5.3. Cell Proliferation

The cell proliferation ability of RAW 264.7 macrophage cells incubated with various substrates were measured using cell counting kit-8 assay (CCK-8, New Cell and Molecular Biotech, Suzhou, Zhejiang, China). Cells were cultured with all kinds of substrates in 24–well plates at a density of <sup>1</sup> <sup>×</sup> 104 cells/well for 7 days. At 1, 3, 5 and 7 d, 1 mL of culture medium containing 100 <sup>μ</sup>L CCK-8 solution was distributed to each well. After incubation for 2 h in the incubator, 100 μL reserved solution was transferred into a new 96-well plate, and then measured by a microplate reader (Cytation 5, Bio-Tek, Winooski, VT, USA) at 450 nm. The values were measured and recorded as mean ± standard deviation, *n* = 3. The cell proliferation tests were repeated three times.

## 2.5.4. Cell Morphology

To examine cell morphology on various surfaces, MC3T3–E1 cells were incubated with different surfaces for 24 h before observation. Then, all samples were rinsed twice with PBS, fixed with 2.5% glutaraldehyde solution (Solarbio, Beijing, China) at 4 ◦C overnight. Then the fixed cells were rinsed by PBS for three times, 10 min each time. Afterwards, samples were further fixed by 1% osmic acid away from light for 1 h and then dehydrated using sequential ethanol solutions (30%, 50%, 75%, 90%, 95% and 100% (*v*/*v*)). Subsequently, various substrates were dried by critical point dryer (EM CPD030, Leica Microsystems, Wetzlar, Germany, the working distance was 8.3–9.4 mm. Lastly, cells grown on various samples were observed by SEM (SU8010, Hitachi, Tokyo, Japan). The typical images from three independent slices were shown for each group.

## 2.5.5. Osteogenic-Related Gene Expression

To detect the osteogenic-related gene expression levels in MC3T3-E1 cells incubated with different substrates, MC3T3–E1 cells were cultured for 14 days and the culture medium was changed daily. At 14 d, cells on the samples were extracted by TRIzol. Expression levels of mouse mRNA encoding genes for alkaline phosphatase (ALP), collagen type 1 alpha 1 (COL1A1), osteopontin (OPN) and osteocalcin (OCN) were detected, with the housekeeping gene *GAPDH* used as the internal control. Reverse transcription and quantitative real-time polymerase chain reaction were performed and data were calculated by the 2−ΔΔCt method. Primers used for the target genes and the housekeeping genes were shown in Table 2. The cDNA amplification of each gene of interest was conducted in triplicate by one sample and data were reported from three separate tests.


**Table 2.** Primers for osteogenesis-related genes.

## 2.5.6. Western Blot Test

The osteogenic-related protein expression levels in MC3T3–E1 cells cultured with various surfaces were examined by Western blot test after 14 days of incubation. MC3T3–E1 cells were cultured for 14 days and medium changed daily. Cells were then harvested on ice, washed twice using PBS, and 200 μL lysis buffer supplemented with protease inhibitor were added to each well. After being incubated on ice for 10 min, cell extracts were subjected to centrifuge with 12,000 rpm at 4 ◦C for 15 min. Protein samples were separated by SDS-PAGE and then transferred to polyvinylidene fluoride (PVDF) membranes. After blocking using skim milk powder (5 g/100mL, Solarbio, Beijing, China) for 1 h, PVDF membranes were hybridized with specific antibodies of ALP (1:1000, ab229126), COL1A1 (1:1000, ab6308), OPN (1:1000, ab63856) and OCN (1:1000, ab93876) from Abcam, Cambridge, UK, respectively overnight at 4 ◦C. Data of expression levels from interested proteins were compared with that of internal control beta-Actin (1:1000, ab8226, Abcam, Cambridge, UK) to normalize loaded protein amount of different samples. The membranes were washed three times, and then incubated with goat anti-mouse secondary antibody at 1:5,000 dilution for 1 h. The protein bands were visualized using chemiluminescence imaging system (ChemiDocTM XRS+ workstation, Hercules, CA, USA). The relative intensity of the protein bands was quantified using image analysis software (Image Pro-Plus 6.0, Media Cybernetics, Silver Spring, MD, USA).

## *2.6. Statistical Analysis*

All experiments were performed three times and the data were expressed as means ± standard deviations (SD). The one-way ANOVA combined with Tukey post hoc test was applied to examine the statistical difference between different samples. All differences considered to be significant when *p* < 0.05.

## **3. Results**

## *3.1. Surface Characterization*

First, the macro surface morphology of various surfaces was observed with stereomicroscope (S9i; Leica Microsystem, Tokyo, Japan) and shown in Figure 1. The Ti slice looks silver and the Ti-PLGA surface observed a transparent film covering the original Ti surface, suggesting that PLGA can form thin film materials on the prepared surface. For the NT surface, a faint gray-green color was seen after anodic oxidation. Whereas, the surface of NT–ICA samples had obvious yellow color, with scattered icariin particles deposited. The NT–ICA–PLGA surface obtained by dripping the PLGA solution on the surface of the NT–ICA sample was still yellow, but the original ICA particles were covered and difficult to identify. Since ASP and PLGA are both white, the NT–ASP/PLGA surface revealed a white-color film on the faint gray-green color NT base material. While for the NT–ICA–ASP/PLGA substrate, a white-color film covering yellow base of NT-ICA was seen.

*Coatings* **2020**, *10*, 427

**Figure 1.** Macro surface morphology of various surfaces was observed under stereomicroscope (S9i; Leica Microsystem, Tokyo, Japan) and representative images were shown (10 ×).

Second, the micro surface morphology of all samples was observed using a scanning electron microscope (SEM, Zeiss Merlin Compact; ZEISS, Jena, Germany). As shown in Figure 2, scratch marks were seen on the surface of Ti after gradient polishing. After loading the PLGA coating, the Ti–PLGA group displayed a smooth and uniform surface. After the anodizing process, nanotubes with uniform and controllable diameter and length can be detected on the surface of NT group. The diameter of nanotubes was 100 ± 5 nm, while the length of which was 3 ± 0.5 μm. On the surface of the NT–ICA group, scattered granular substances of ICA were observed. For the NT–ICA–PLGA surface, ICA was completely enwrapped by the uniform smooth film of PLGA coating, with many needle-like configurations indicating the existence of ICA. While on the NT–ASP/PLGA substrate, short needle-like ASP crystals were observed in the PLGA coating layer. The morphology of NT–ICA–ASP/PLGA surface combined the common features of NT–ASP/PLGA and NT–ICA–PLGA surfaces, and displayed both the configuration of ICA and ASP in the PLGA coating.

**Figure 2.** Surface micromorphology images of different substrates were obtained with scanning electron microscope (SEM, Zeiss Merlin Compact; ZEISS, Jena, Germany). Red and white arrows indicate icariin (ICA) and aspirin (ASP), respectively; 1000×, scale bar = 10 μm. Insets of images of NT and NT-ICA substrates were magnified SEM graphs (30,000×, scale bar = 200 nm).

XPS results were used to determine the elemental composition of different samples. All binding energies were referenced to the C 1s spectrum peak (284.8 eV) as an internal reference after calibrating peak position. As shown in Figure 3 and Table 3, the spectra of the Ti and NT mainly include C1s, O1s and Ti2p3. For the NT-ICA surface, increased peak of C1s may originate from ICA (C33H40O15), while the decreased peak of Ti2p3 may be explained by the partial coverage of the TiO2 nanotubes structure by ICA, compared to the NT surface.

**Figure 3.** X-ray photoelectron spectroscopy (XPS) analysis displaying chemical composition of different surfaces.



In addition, the disappearance of the Ti peak on the Ti–PLGA, NT–ICA–PLGA, NT–ASP/PLGA and NT–ICA–ASP/PLGA surfaces indicate that the PLGA coating totally overspread the original Ti and nanotubes structures. As we known, PLGA ([C5H8O5]n), ICA (C33H40O15) and ASP (C9H8O4) are all composed of C, H and O. Due to the small photoionization cross section of H, the signal of H is too weak for XPS to be detected. Therefore, only C and O could be detected on the Ti–PLGA, NT–ICA–PLGA, NT–ASP/PLGA and NT–ICA–ASP/PLGA surfaces by XPS analysis. Moreover, since both icariin and aspirin have higher C and lower O element content than PLGA, the increased C and decreased O element content observed in NT–ICA–PLGA, NT-ASP/PLGA and NT–ICA–ASP/PLGA groups compared to the Ti–PLGA group may be explained by the successful loading of ICA and ASP.

In addition to surface topography and chemical composition, hydrophilicity is also an essential factor, which affects the biological compatibility of materials. Figure 4 showed that the contact angle of Ti was 89.5 ± 0.6◦, while after the application of PLGA coating, the contact angles decreased to Ti–PLGA (69.5 ± 0.6◦), NT–ICA–PLGA (58.9 ± 0.8◦), NT–ASP/PLGA (69.8 ± 0.5◦) and NT–ICA–ASP/PLGA (64.5 ± 1.6◦), respectively, with statistical significance. This indicates that PLGA coating could improve the wettability of surfaces, compared to pure Ti surface. Moreover, the contact angles of the NT and NT–ICA groups were 19.5 ± 3.0◦ and 21.3 ± 1.0◦, respectively, even lower than those of various PLGA coated substrates, with statistical significance. This phenomenon suggested that the titanium nanotube structure fabricated after anodization could substantially enhance the hydrophilicity of surfaces. In this experiment, we only detected the initial contact angles. In a follow-up experiment, we will further test the frequency acquisition of the droplet images and monitor the time dependence of the water

contact angle so as to obtain more kinetic constant and exponential parameters of material and evaluate hydrophilicity of the material much more accurately.

**Figure 4.** Contact angles of various surfaces. (**A**) Contact angle degree of different surfaces. Different lowercase letters indicate statistically significant differences between different groups (*p* < 0.05). Data are expressed as mean ± SD, *n* = 3 replicates each group. (**B**) Images of contact angles on different surfaces.

## *3.2. In Vitro Drug Release Profile*

The drug release pattern of ICA in the NT-ICA-ASP/PLGA surface is shown in Figure 5A. A sustained release profile of ICA was seen during the 30 days detection period. At the same time, the daily concentration of ICA always maintained within the range of 0.1–1 <sup>×</sup> 10−<sup>5</sup> mol/L per day, which was proved to be effective concentration of icariin to promote MC3T3–E1 cells osteogenesis in our preliminary experiment (data not shown). Taken into consideration of this point, the addition of ICA in the NT–ICA–ASP/PLGA surface is expected to exert continued osteogenic function during the release period.

As shown in Figure 5B, the release profile of ASP also displayed controlled release pattern during the release period, and finally reached a plateau at day 7. In addition, the release concentration of ASP maintained between 0.2 and 4 mM during day 1–5, within the effective concentration of anti-inflammatory and proregenerative function of ASP by our preliminary data (data not shown). Therefore, we hypothesized that the addition of ASP in the coating may promote RAW 264.7 macrophage cells to change from M1 to M2 polarization status.

Meanwhile, the release kinetics of both ICA and ASP in the NT–ICA–ASP/PLGA surface were assessed by fitting the release profiles of the two drugs to the Korsmeyer−Peppas model [41]. The mechanism of ICA and ASP release from the established surface during dissolution investigations in PBS was determined using Equation (1):

$$\frac{M\_{\rm f}}{M\_{\rm sv}} = kt^n \tag{1}$$

where *Mt <sup>M</sup>*<sup>∞</sup> is the fractional drug release at denoted time point *<sup>t</sup>*, *<sup>k</sup>* is a kinetic constant reflecting the structural and geometric properties of the drug/carrier system, and *n* is the release exponent, which relies on the release mechanism. As shown in Figure 6, the first 10 days of ICA release profile was used for simulation and the release exponent *n* was 1 ± 0.0251, which is close to a desirable time-dependent release mechanism, as indicated by Korsmeyer et al. [41]. While for ASP, since there were relatively less time points for simulation, we selected the first 4 days for simulation and the diffusional release exponent *n* value was 0.5438 ± 0.0725, which implied a drug release mechanism approximately controlled by anomalous (non-Fickian) diffusion. As illustrated in Figure 6, the data fit the model well as the correlation coefficient *R*<sup>2</sup> of ICA and ASP release curves are 0.9946 and 0.9485, respectively. In addition, the kinetic constant of ICA and ASP were 5.9699 ± 0.2496 and 47.5629 ± 3.2833, respectively, which indicated different structural and geometric properties of the two kinds of drugs in the drug/carrier system. The drug release percentage scatter diagrams as a function of time and sketch view of the established substrate were shown in Figure 6A–C. The sustained release mechanism of ICA may partly be interpreted by its poor water solubility [42] and the existence of the outside ASP/PLGA coating, which generate a viscous layer that further retarded the drug diffusion to the solvent, similar to a previous report [43]. In the future, drug release amount at more time points, especially for ASP, are needed to be measured to evaluate the kinetic mechanism much accurately.

**Figure 5.** In vitro release profile of ICA (**A**) and ASP (**B**) on the NT-ICA-ASP/PLGA surface. (**a**) Cumulative release amount curve. (**b**) Cumulative release percentage curve. (**c**) Daily release concentration curve. Data are expressed as mean ± SD (*n* = 3).

**Figure 6.** (**A**) ICA release behaviors in the first ten days. (**B**) ASP release behaviors in the first four days. (**C**) Sketch view of the NT–ICA–ASP/PLGA release system.

### *3.3. Polarization Status of RAW 264.7 Cells Cultured on Di*ff*erent Surfaces*

#### 3.3.1. Proinflammatory (M1) and Proregenerative (M2) Marker Genes Expression

Quantitative real-time polymerase chain reaction (qPCR) was performed to detect expression levels of proinflammatory (M1) and proregenerative (M2) gene markers in RAW 264.7 macrophage cells incubated with different substrates. The results were displayed in Figure 7. In the physiological healing of wound, the inflammatory phase last hours to days. Therefore, the initial several days after dental implant insertion was crucial for ideal osseointegration. Since the drug release curve of ASP revealed a relative steady release profile during day 2–4, we selected day 3 to evaluate the polarization of macrophages. As shown in Figure 7, M1 proinflammatory genes IL-1β and TNF-α expression levels decreased in cells cultured on the NT surface, compared with those on the pristine Ti surface, indicating that the 100 nm diameter NT used in this experiment could inhibit inflammatory reaction at the gene level. Whereas the Ti−PLGA surface evoked higher proinflammatory genes expression than Ti surface. It indicated that although PLGA is a non-toxic and harmless material that can be absorbed by the human body, it could promote inflammatory reaction in RAW 264.7 cells during the early period after implantation. However, the decreased M1 proinflammatory gene levels in cells incubated with the NT-ICA-PLGA surface, compared with those with the Ti-PLGA surface implied that ICA may exert anti-inflammatory function. In particular, the addition of ASP, a well-established non-steroidal anti-inflammatory drug, obviously decreased proinflammatory gene expression in macrophage cells cultured on NT−ASP/PLGA and NT−ICA−ASP/PLGA substrates, compared with the Ti surface.

The M2 proregenerative marker genes expression in RAW 264.7 macrophage cells cultured on various surfaces were further evaluated. After 3 days of incubation, the addition of PLGA coating on the Ti−PLGA surface did not change TGF-β and HO-1 gene expression levels much, compared with Ti surface. It indicated that PLGA had no significant effect on induction of macrophages towards M2 polarized status. Previously, ICA has been reported to exert anti-inflammatory effect by elevated cytoprotective gene expression HO−1 [44,45]. Therefore, the observed highest HO−1 gene level on the NT−ICA surface may be attributed to the burst release of ICA occurred during the first 3 days, which evoked the substantially increased HO−1 gene expression. Most importantly, TGF−β and HO−1 gene levels in cells on the NT−ICA, NT−ICA−PLGA, NT−ASP/PLGA and NT−ICA−ASP/PLGA substrates were obviously higher than those on the Ti surface, with statistical significance. These results suggest that both ICA and ASP could induce M2 polarization of macrophage, which endowed the NT−ICA−ASP/PLGA substrate with superior proregenerative immunomodulatory ability.

**Figure 7.** Real-time PCR results of interleukin-1β (IL-1β), tumor necrosis factor-alpha (TNF-a), transforming growth factor-beta (TGF-β) and heme oxygenase-1 (HO-1) gene expression levels in the RAW 264.7 cells on different substrates after 3 days of incubation. Different lowercase letters indicate statistically significant differences between different groups (*p* < 0.05). Data are mean ± SD, *n* = 3 replicates each group.

#### 3.3.2. Enzyme-Linked Immunosorbent (ELISA) Assay

We further examined the proinflammatory (M1) and proregenerative (M2) marker protein expression levels by ELISA assay. As shown in Figure 8, decreased IL−1β and TNF−α and increased TGF−β and HO−1 protein levels were observed in cells cultured on the NT surface, compared with those on the Ti surface, displaying that the nanotube structure could reduce inflammation and induce regeneration. On the contrary, proinflammatory IL−1β and TNF−α protein levels were enhanced, while the proregenerative TGF−β and HO−1 protein levels do not change much in cells on the Ti−PLGA surface, compared with the Ti surface. It showed that the PLGA coating could induce M1 inflammatory reaction, but failed to induce M2 polarization in macrophages. After the addition of ICA and ASP, M1 proinflammatory IL−1β and TNF-α protein levels reduced, whereas TGF−β and HO−1 protein levels elevated in cells on the NT−ICA, NT−ICA−PLGA, NT−ASP/PLGA and NT−ICA−ASP/PLGA substrates, compared to the original Ti surface, with statistical significance. In particular, NT-ASP/PLGA and NT−ICA−ASP/PLGA groups had relatively lower M1 and higher M2 proteins secretion than the other groups.

Combined the results from gene and protein detection in RAW 264.7 cells, similar to the NT−ASP/PLGA surface, the NT−ICA−ASP/PLGA substrate also exhibited a superior ability of inflammation inhibition and regeneration enhancement than other surfaces. It implies that the combination of ASP and ICA on the NT−ICA−ASP/PLGA surface could achieve anti-inflammatory and proregenerative function.

**Figure 8.** ELISA results representing the RAW 264.7 cytokine secretion on different substrates after 3 days of culture. Different lowercase letters indicate statistically significant differences between different groups (*p* < 0.05). Data are mean ± SD, *n* = 3 replicates each group.

*3.4. Behaviors of MC3T3-E1 Cells on Various Surfaces in Conditioned Medium (CM)*

## 3.4.1. Cell Proliferation

The proliferation of MC3T3-E1 cells seeded on all samples was assessed by CCK-8 assay. Figure 9 showed cell numbers cultured on different surfaces after 1, 3, 5 and 7 days of incubation. At day 1 and 3, there was no statistical significance in the cell proliferation capacity of each group although a tendency of a higher cell proliferation rate on the NT−ICA−ASP/PLGA substrate was seen, compared with the Ti group. After 5 and 7 days of incubation, cells cultured on NT−ASP/PLGA and NT−ICA−ASP/PLGA substrates revealed higher cell proliferation ability compared with that on the pristine Ti, *p* < 0.05. This suggests the relative strong effect of aspirin to promote osteoblast proliferation under the existence of CM from macrophage cells. Interestingly, the proliferation rate of cells on the NT−ASP/PLGA group was moderately higher than that on the NT−ICA−ASP/PLGA group, although there was no statistical difference. This phenomenon may be explained by that the NT−ICA−ASP/PLGA substrate had an additional prodifferentiation effect on osteoblasts than the NT−ASP/PLGA surface, due to the existence of ICA. It is well accepted that during cell life activity, higher proliferation activity is often inhibited when cell differentiation is promoted and vice versa [12,46]. Therefore, the slightly inhibited cell proliferation of the NT−ICA−ASP/PLGA group may imply its enhanced cell differentiation due to the additional prodifferentiation effect of released ICA, compared to the NT−ASP/PLGA group.

122

**Figure 9.** The CCK-8 results of MC3T3−E1 cells cultured on different surfaces for 1, 3, 5 and 7 days. Different lowercase letters indicate statistically significant differences between different groups (*p* < 0.05). Data are mean ± SD, *n* = 3 replicates each group.

## 3.4.2. Cell Morphology

Cell morphologies of MC3T3−E1 osteoblast cells on different samples after 24 h of incubation were acquired with scanning electron microscope (SEM, SU8010, Hitachi, Tokyo, Japan) and shown in Figure 10. After 24 h of incubation, cells attached along the direction of scratch marks on the pure Ti surface. While on the NT and NT−ICA substrates, cells revealed well-distributed morphology, with filopodia stretching into the nanotubular structures, which implies good cell adhesion induced by the nanostructure. In addition, cells on the Ti−PLGA, NT−ICA−PLGA, NT−ASP/PLGA and NT−ICA−ASP/PLGA substrates displayed a spread-out morphology, towards different directions. This can be explained by that PLGA improved cell adhesion to some extent, similar with previous report [47]. In particular, for the NT−ICA−ASP/PLGA substrate, cells extend towards all directions, which revealed the satisfied cell adhesion ability.

**Figure 10.** Images of MC3T3−E1 cells showing the representative morphology on different substrates after 24 h of culture were obtained with scanning electron microscope (SEM, SU8010, Hitachi, Tokyo, Japan). The pseudo colored pink cells indicate morphology of MC3T3−E1 cells; 1000×, scale bar = 10 μm.

#### 3.4.3. Osteogenesis-Related Gene Expression

Figure 11 shows osteogenesis-related genes expression of MC3T3−E1 cells cultured on different surfaces after 14 days of culture under the condition of macrophage CM. Gene expression levels of ALP, COL1A1 and OCN all increased on osteoblast cells cultured on the NT−ICA−PLGA, NT−ASP/PLGA and NT−ICA−ASP/PLGA substrates, compared with that on the pristine Ti and Ti-PLGA surfaces, *p* < 0.05. It's worth noting that the NT−ICA−ASP/PLGA substrate showed highest expression levels for all of the osteogenesis-related genes examined than the other surfaces, with statistical significance. These results suggested that the employment of ICA and ASP could both improve osteoblast differentiation, and the superior osteogenic function of the NT−ICA−ASP/PLGA substrate lies in the synergistic function of ICA and ASP.

**Figure 11.** Real-time PCR results representing the gene expression levels of alkaline phosphatase (ALP), collagen type 1 alpha 1 (COL1A1), osteopontin (OPN) and osteocalcin (OCN) in MC3T3-E1 cells after 14 days of culture on different samples. Different lowercase letters indicate statistically significant differences between different groups (*p* < 0.05). Data are mean ± SD, *n* = 3 replicates each group.

## 3.4.4. Western Blot Test

Western blot analysis was used to evaluate the protein expression levels in MC3T3-E1 cells cultured on different samples for 14 days. Figure 12A,B displayed Western blot bands and quantitative analysis, respectively.

Consistent with the gene expression levels, the protein expression levels of ALP, COL1A1 and OCN were all higher in cells incubated with NT−ICA−PLGA, NT−ASP/PLGA and NT−ICA−ASP/PLGA substrates, compared to that on the original Ti and Ti-PLGA surfaces. Especially, NT−ICA−ASP/PLGA substrate revealed enhanced expression levels for all of the osteogenic proteins determined in this experiment than those on the other surfaces, *p* < 0.05. The optimal osteogenesis-related protein expression of the NT−ICA−ASP/PLGA substrate is attributed to the combined effects of both ICA and ASP to facilitate bone formation in osteoblast cells.

**Figure 12.** (**A**) Western blot showing the protein expression levels of ALP, COL1A1, OPN and OCN and beta-Actin in MC3T3-E1 cells after 14 days of culture on different samples. (**B**) The corresponding gray-scale values of proteins levels in the 7 groups. Different lowercase letters indicate statistically significant differences between different groups (*p* < 0.05). Data are mean ± SD, *n* = 3 replicates each group.

#### **4. Discussion**

Numerous strategies have been implemented to integrate bioactive molecules into the implant surface so as to promote osseointegration. Previously, most investigators endeavored to optimize dental implant materials to promote osteogenic differentiation in mesenchymal stem cells or osteoblast cells and have developed satisfied implant materials to a certain degree. However, only promoting bone formation in osteoblast cells is not enough to fulfill implant osseointegration in vivo. The reason is that after implant were inserted in to the bone tissue, in addition to its interaction with osteoblast, immune cells, such as macrophages, also play an important role in osseointegration [48]. In order to uncover the close relationship between skeletal and immune system, we applied the icariin/aspirin composite coating on TiO2 nanotubes surface, named as the NT-ICA-ASP/PLGA substrate, and explore its immunomodulatory effect on macrophage and osteogenic activity in osteoblast cells. We hypothesized that the NT–ICA–ASP/PLGA substrate could induce immunoregulatory function in macrophages and provide an immune microenvironment, which is favorable for osteogenesis in osteoblast cells.

In the current work, the NT-ICA-ASP/PLGA substrate was fabricated. The SEM, XPS detection and contact angle tests confirmed the successful preparation of the icariin/aspirin functionalized coating on TiO2 nanostructured surface. The established NT–ICA–ASP/PLGA substrate achieved controlled release of both ICA and ASP, with satisfied hydrophilicity. The drug release curves revealed that the releasing period of aspirin is about 7 days, while icariin achieved a sustained release period until 30 days.

To maintain the controlled release of the two kinds of drugs, we firstly prepared TiO2 nanotubes on the surface of pure titanium by anodization. It is well established that 70–100 nm diameter nanotubes could induce improved cell adhesion, differentiation toward osteoblasts [49–51]. Previously, we have chosen 80 nm diameter nanotubes for the fabrication of functionalized coating [17]. In the present study, we selected a 100 nm diameter of TiO2 nanotubes for further functionalization study after taking two reasons into considerations. First, as a drug loading reservoir, the larger diameter of the nanotubes, the stronger ability to load drugs [52]; second, 100 nm diameter nanotubes displayed highest ability of protein adsorption, which is critical in mediating cell attachment and proliferation so as to promote bone formation [53,54]. In view of this, we selected 100 nm diameter nanotube as a drug carrier, hoping that it will balance the drug-loading capacity and osteogenic ability at the same time. Another thing should be clarify is that we changed the anodization parameters, compared to our past article [17]. In our previous work, HF was used as an electrolyte, while in the present experiment we employed the reaction system composed of NH4F, ethylene glycol and water. The reason that we adopted this new method lies in that it can increase the viscosity of the electrolyte and reduce the activity of fluoride ions in order to acquire much controllable and orderly aligned nanotubes with uniform diameter [55]. In present work, we analyzed the bioactivity, gene expression and immune regulatory responses of the established icariin/aspirin composite coating on 100 nm diameter TiO2 nanotubes. In the future, we will continue to evaluate the influence of different sizes of TiO2 nanotubes on the biological behaviors of the coated surface.

We further used PLGA to avoid the initial burst release of drug from the TiO2 nanotubes surface. PLGA is commonly applied in bone tissue engineering because of its favorable mechanical properties and drug-loading capacity [56,57] and used in wound dressings due to its excellent bonding ability [58]. PLGA has also been employed as coating materials on the surface of titanium implants [59,60]. Serving as biodegradable material [61], PLGA polymer-coated nanotube displayed longer drugs release period compared to chitosan-coated nanotubes [62]. In addition, compared with protein growth factor, icariin and aspirin were both chemical stable and heat-resistant drugs [63,64]. More importantly, since we adopted the dip-coating method for preparation of the PLGA coating onto the NT–ICA–ASP/PLGA substrate, this mild coating method could further protect icariin and aspirin from denaturation during the coating procedure and achieve sustained and multiple release of two kinds of drugs at the same time. Therefore, in this way, we successfully superimposed the ASP/PLGA coating onto the ICA loaded TiO2 nanotubes structure so as to obtain the established NT–ICA–ASP/PLGA substrate. In the future, we will further detect the interfacial adhesion strength between the coating and titanium substrate.

We have previously reported that ICA and TiO2 nanotubes structures could exert osteogenic effects synergistically [17]. In the present work, apart from its well-established osteogenic effect, we also intend to explore other functions of ICA. Although there is less report about this issue, it has been revealed that icariin could exert anti-inflammatory function [18]. Icariin and its derivate have demonstrated anti-inflammatory effects in macrophage cell lines, human myeloid cells, and a mouse model of inflammation. Additionally the anti-inflammatory effect was concomitant with a down-regulation of IL–10, IL–6 and TNF–α [19]. Other researchers also found that icariin possess significant therapeutic effect on rheumatoid arthritis (RA) clinically [9]. Moreover, the mechanism of

icariin to prevent RA might be related to its immunoregulatory function, which was mediated through the decrease in the cell number of immune cells Th17. Additionally, icariin administration could inhibit IL–17 production [20]. Taking the above-mentioned factors into consideration, we attempt to discover the anti-inflammatory and immunoregulatory functions of ICA in this work. In the present study, during inflammation markers detection, we discovered decreased IL–1β gene expression and inhibited IL–1β and TNF–α protein levels on the NT–ICA and NT–ICA–ASP/PLGA surfaces, compared to Ti surface. Moreover, enhanced gene and protein levels of M2 phenotype markers TGF–β and HO–1 were also observed on ICA–loaded surfaces, compared with Ti and Ti–PLGA substrates. These results confirmed the anti-inflammatory and immunoregulatory functions of ICA.

Aspirin (ASP), a classic non-steroidal anti-inflammatory drug, remains the first-line clinical drug due to its good antipyretic and analgesic capabilities and extremely low drug side effects [65]. Some researchers showed that ASP down-regulate iNOS and TNF-α expression in macrophage cells in vitro and improve bone regeneration in vivo by inhibiting LPS-induced macrophage activation in the early stages of inflammation [66]. In addition, aspirin may be a promising option for preventing and curing osteoclastic bone destruction, including peri-implant osteolysis [67]. Another researcher also proved that on the surface of titanium primed with phase-transited lysozyme, a coating loaded with aspirin could promote osseointegration [68]. In our experiment, compared with Ti surface, NT–ASP/PLGA and NT–ICA–ASP/PLGA surfaces displayed decreased IL-1β gene level and reduced IL–1β and TNF–α protein levels, with upregulated TGF–β and HO–1 gene and protein expression levels, *p* < 0.05. As expected, our results confirmed the anti-inflammatory and immunoregulatory functions of ASP, with the established NT–ICA–ASP/PLGA substrate exhibiting decreased M1 inflammatory and increased M2 proregenerative effects at the same time.

It is well known that macrophages are the main effector cells of the inflammatory response, and have great plasticity. When stimulated by different signals, they can show a M1 or M2 phenotype [38,39]. The timely transformation of M1–type to M2–type macrophages effectively promotes the regression of inflammation and tissue repair [69,70]. Generally speaking, classically activated macrophages (M1) secrete a variety of proinflammatory factors, such as TNF–α, IL–1β, IL–6, etc. While the alternatively activated M2 phenotype macrophages generate anti-inflammatory or immunoregulatory cytokines such as TGF–β and IL–10 [71]. In addition, heme oxygenase HO–1, the derivable isoform of the heme-degrading enzyme HO, plays an important role in inflammation and immunoregulation of homeostasis. Myeloid HO–1 expression modulates macrophage polarization to M2 [72,73]. Furthermore, high HO-1 level increased the expression of osteonectin, OPG and BMP–2, and increase osteoblast function and differentiation [74]. In our experiment, down-regulated M1 proinflammatory and up-regulated M2 proregenerative genes and proteins expression were shown on the novel icariin/aspirin composite coating covered TiO2 nanotubes surface, compared with pure Ti surface. These results indicated that the NT–ICA–ASP/PLGA substrate revealed optimal ability, which could inhibit inflammation and enhance regeneration simultaneously, probably attributed to the synergistic effects of ICA and ASP.

Recently, osteoimmunomodulation (OIM) has been raised to emphasize the importance of immune response during osteogenesis at biomaterial–tissue interface. As the concept of OIM has been gradually accepted, more and more attention has been paid to the positive regulation of immune response at implant–tissue interface so as to accelerate implant osseointegration [75]. After implant insertion, inflammation should be properly regulated to achieve faster bone integration, particularly for patients who are suffering from systemic diseases, which were adverse to bone formation [76,77]. By regulating the early inflammatory response of macrophages, tissue repair can be enhanced, thereby promoting the early rapid formation of bone tissue around the implant.

In order to evaluate OIM, we firstly extracted conditioned media (CM) from macrophage cells incubated with various implant materials. Then the extracted CM was added to osteoblast cells to simulate the immune microenvironment around the implant site. Thus, the indirect coculture system composed of sample slices, macrophage CM and osteoblast cells is ready to evaluate the effects of immune cells on osteoblast differentiation. Our results suggested that in the coculture system, osteoblast cells on the NT–ASP/PLGA and NT–ICA–ASP/PLGA substrates displayed increased cell proliferation, compared with the other groups at day 5 and 7. Moreover, cells showed well–spread morphology on the NT–ICA–ASP/PLGA substrate, suggesting good cell adhesion on the modified substrate. In addition, we also detected the osteogenic gene and protein levels in the coculture system. During the osteoblast cell differentiation process, an early increase in alkaline phosphatase (ALP) was observed. Then followed by an augment in collagen type 1 and OPN, with osteocalcin expressed later [78]. We measured expression levels of the abovementioned four typical marker genes, including *ALP*, *COL*, *OPN* and *OCN*, which expressed at different stages of osteoblast differentiation and osteogenesis. As predicted, the highest gene levels were observed on the NT–ICA–ASP/PLGA substrate, implying the collaborative pro-osteogenic effect of ICA and ASP at the genetic level. Not only that, both ICA and ASP modified surfaces revealed relatively higher protein levels of osteogenic proteins than pure Ti surface. Especially, the NT–ICA–ASP/PLGA substrate demonstrated the highest expression levels for all of the osteogenesis-related proteins. It is well-accepted that enhanced gene and protein expression levels related to osteoblast cell differentiation may accelerate osseointegration at the implant–tissue interface. The advanced expression levels of both osteogenesis-related genes and proteins on the NT-ICA-ASP/PLGA substrate implied its superior ability to induce osteoblast cells differentiation. The optimal effect of the NT-ICA-ASP/PLGA substrate might be ascribed to the synergistic effects of ICA and ASP, which is in line with its excellent cell proliferation and adhesion results.

#### **5. Conclusions**

In view of the concept of OIM, it is especially to be expected to fabricate an implant material, which possesses both early immune regulatory property and long-term osteogenic ability. Both ICA and ASP are candidate drugs for implant surface modification with multiple function including anti-inflammation, immunoregulation and osteogenesis. Taken into consideration of the factors mentioned above, we constructed the NT–ICA–ASP/PLGA substrate. The novel substrate could trigger an effective and timely shift of macrophage from the inflammatory reaction stage to restorative stage. The established NT–ICA–ASP/PLGA surface endowed the modified surface with immunomodulatory function in macrophages and osteogenic effect in osteoblast cells at the same time. This study provides a very promising strategy for fabricating functionalized coating on titanium-based alloys to improve implant osseointegration.

**Author Contributions:** Conceptualization, Y.L. (Ying Li) and C.L.; Data curation, A.M. and Y.Y.; Funding acquisition, Y.L. (Ying Li) and C.L.; Investigation, A.M., Y.Y., B.C., and W.W.; Methodology, A.M., Y.Y., J.L., H.Q., and Y.L. (Yunkai Liang); Project administration, Y.L. (Ying Li) and C.L; Writing—original draft, A.M. and Y.Y.; Writing—review and editing, Y.L. (Ying Li). All authors have read and agreed to the published version of the manuscript.

**Funding:** This research was funded by National Natural Science Foundation of China, Grant No. 81870809, 81500886, 31470920 and Tianjin Natural Science Foundation Grant No. 16JCYBJC28700.

**Acknowledgments:** We sincerely acknowledge Baoe Li from the School of Materials Science and Engineering, Hebei University of Technology, for her precious technical support and discussions. We also thank Huanhuan Zhai from Tianjin Institute of Industrial Biotechnology, Chinese Academy of Sciences, for her important work in preparing SEM images.

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

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*Article*
