**Surface Engineering of Light Alloys**

Editor

**Sara Ferraris**

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*Editor* Sara Ferraris Politecnico di Torino Italy

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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/surf eng light alloy).

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## **Contents**

## **About the Editor**

**Sara Ferraris** received her PhD in Biomedical Engineering in 2010 from Politecnico di Torino. She is Assistant Professor in the Department of Applied Science and Technology at Politecnico di Torino. Her main research interests deal with surface properties of metallic materials for different applications (biomedical, mechanical, industrial), surface modification strategies (e.g., functionalization, coating, chemical or physical treatments), and surface characterizations. She is the author of 102 papers in international journals and holds 5 patents in the materials science field.

## *Editorial* **Special Issue: Surface Engineering of Light Alloys**

#### **Sara Ferraris**

Department of Applied Science and Technology, Politecnico di Torino, C.so Duca degli Abruzzi 24, 10129 Torino, Italy<EMAIL_ADDRESS>
Received: 27 November 2020; Accepted: 30 November 2020; Published: 1 December 2020

Light alloys (mainly aluminum, magnesium and titanium alloys) are of great interest in applications where lightweight has an high impact, such as automotive, aerospace and biomedical fields.

In addition to their bulk properties, such as mechanical properties, thermal properties or density, surface properties are crucial in many applications. In fact the surface is the first layer of the material exposed to the working environment. In biological systems this means that the surface is the first approach of the material to biological entities and it plays a crucial role in situations like tissue integration or bacterial contamination. Moreover the surface is a critical area for corrosion or wear phenomena both in biomedical and automotive or aerospace applications. Finally the surface can be the starting point for joining processes, in this scenario proper surface preparation can significantly affect the joint performance.

In this context surface engineering of light alloys is a versatile instrument to optimize surface properties of these materials for specific applications, without altering their bulk properties.

The present special issue covers all the above described topics with 10 research papers.

Five of them are related to biomedical applications of titanium an magnesium alloys. Surface engineering of titanium alloys moves from the investigation of Ti/Ti-alloys wear behavior in presence of artificial saliva (paper from Alemanno et al. [1]), to bioactive inorganic coatings for the improvement of bone bonding ability of Ti alloys (Yamaguchi et al. paper [2]) and finally to surface functionalization with natural molecules (Cazzola et al. paper [3]) to improve biological properties of Ti alloys. The main criticism of Mg alloys in biomedical applications is related to their poor corrosion resistance and too rapid degradation. Two solutions are explored in the present special issue, organic dopamine functionalization (Ghanbari et al. paper [4]) or inorganic coatings (Mg oxides and carbonates, Xie et al. paper [5]).

The protection of Mg alloys from corrosion and wear is of interest also for industrial applications far from the biomedical ones, Ni-P coatings are investigated for this purpose in the Buchtik paper [6].

Protective coatings (alumina) for industrial applications on Al alloys have been investigated from the wettability and mechanical standpoints in the papers from Nied ´zwied ´z et al. in the present special issue [7,8].

The possibility to activate Al/Al-alloys surface to optimize their joining ability has been analyzed in the paper from Ferraris et al. [9].

Finally the effect of high energy irradiation on the microstructure of Al alloys is the key topic of Wan et al. paper [10].

**Conflicts of Interest:** The author declares no conflict of interest.

#### **References**

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© 2020 by the author. 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* **Tribological Behaviour of Ti or Ti Alloy vs. Zirconia in Presence of Artificial Saliva**

### **Fabio Alemanno 1, Veronica Peretti 2, Angela Tortora 1 and Silvia Spriano 2,\***

Received: 30 July 2020; Accepted: 26 August 2020; Published: 31 August 2020

**Abstract:** Abutment is the transmucosal component in a dental implant system and its eventual appearance has a major impact on aesthetics: use of zirconia abutments can be greatly advantageous in avoiding this problem. Both in the case of one and two-piece zirconia abutments, a critical issue is severe wear between the zirconia and titanium components. High friction at this interface can induce loosening of the abutment connection, production of titanium wear debris, and finally, peri-implant gingivitis, gingival discoloration, or marginal bone adsorption can occur. As in vivo wear measurements are highly complex and time-consuming, wear analysis is usually performed in simulators in the presence of artificial saliva. Different commercial products and recipes for artificial saliva are available and the effects of the different mixtures on the tribological behaviour is not widely explored. The specific purpose of this research was to compare two types of artificial saliva as a lubricant in titanium–zirconia contact by using the ball on disc test as a standard tribological test for materials characterisation. Moreover, a new methodology is suggested by using electrokinetic zeta potential titration and contact angle measurements to investigate the chemical stability at the titanium–lubricant interface. This investigation is of relevance both in the case of using zirconia abutments and artificial saliva against chronic dry mouth. Results suggest that an artificial saliva containing organic corrosion inhibitors is able to be firmly mechanically and chemically adsorb on the surface of the Ti c.p. or Ti6Al4V alloy and form a protective film with high wettability. This type of artificial saliva can significantly reduce the friction coefficient and wear of both the titanium and zirconia surfaces. The use of this type of artificial saliva in standard wear tests has to be carefully considered because the wear resistance of the materials can be overestimated while it can be useful in some specific clinical applications. When saliva is free from organic corrosion inhibitors, wear occurs with a galling mechanism. The occurrence of a super-hydrophilic saliva film that is not firmly adsorbed on the surface is not efficient in order to reduce wear. The results give both suggestions about the experimental conditions for lab testing and in vivo performance of components of dental implants when artificial saliva is used.

**Keywords:** artificial saliva; lubricant; zirconia; titanium alloys; wear

#### **1. Introduction**

The term biotribology comprises tribological phenomena occurring both in natural tissues and organs of living organisms (such as wear of skin or sliding of eyelids over eye) and after implantation of an artificial device (wear of orthopaedic and dental implants) [1]. This research focused on the last case, and deals with the tribology of materials used for dental implants with a focus on the interface between titanium and zirconia.

Wear of natural or artificial teeth mainly results from the closed phase of chewing, with food–tooth contact, or eventually from thegosis (sliding teeth into lateral positions) and bruxism (grinding teeth without the presence of food) when direct tooth–tooth contact occurs. Furthermore, nowadays, chemical effects play an increasingly important role in tooth wear, mainly as a result of the large consumption of acid drinks (erosion). Additionally, tooth wear can also result from tooth cleaning or pipe smoking. If moderate tooth wear has significant clinical consequences (aesthetically and functionally), excessive wear can result in unacceptable damage of the occluding surfaces, alteration of the functional path of masticatory movement, dentine hypersensitivity, and pulpal pathology [2]. Biotribology must include both standard tests for materials characterization and specific tests for the simulation of the in vivo environment and mechanical loading [3]. This research focused on the first topic, but with a focus on the selection of the lubricant, which is, in this case, artificial saliva.

Saliva is the lubricant in the human mouth and consists of approximately 98% water plus a variety of electrolytes and proteins [4]. Salivary proteins can be selectively adsorbed onto solid substrates as well as mucosa membranes exposed to an oral environment, forming a pellicle within a few seconds. Saliva is a source of inorganic ions necessary for remineralization and self-repair capacity because it supplies calcium and phosphate ions. The role of the salivary pellicle is also thought to protect the underlying tooth surface from acid attack because it is a buffer to acids introduced into the mouth [5]. An important function of saliva is to form a boundary lubrication system and to act as a lubricant between hard (tooth) and soft (mucosal) tissues or in dental implants, decreasing wear and reducing friction [5]. It has been widely reported that a salivary film is formed by a layer-by-layer adsorption of salivary proteins and then has a heterogeneous structure consisting of a thin and dense inner layer, able to decrease the friction and wear of teeth, and a thick, highly hydrated and viscoelastic outer layer [6]. The lubricating properties of saliva are presumably related to the strong adhesion strength of the inner salivary film on the tooth surface, which depend on surface roughness, free energy, and the presence of positive or negative surface charges [5]. Harvey et al. observed that the friction coefficient on a hydrophobic substrate was almost an order of magnitude higher than that on a hydrophilic one because of the lower adhesion strength of the salivary film [7]. On the other side, hydrophobic materials in the oral cavity might be more easily cleaned from adsorbed salivary films.

As the transmucosal component in the implant system, the abutment is exposed to the oral cavity and its eventual appearance has a major impact on the aesthetics: using zirconia abutments, aesthetics can be indeed improved. Zirconia abutments are available in two basic designs: one-piece abutment (all parts are made of zirconia in a unit) or two-piece abutment (a secondary metallic component is used as connecting part). Several studies have verified that the fracture strength of two-piece abutments is higher than that of the one-piece design [8,9]. In both cases, the critical issue is severe wear between the zirconia and titanium inducing loosening of the abutment connection and production of titanium wear debris, which can result in peri-implant gingivitis, gingival discoloration, or marginal bone adsorption [10].

A comprehensive description of the tribocorrosion behaviour of different materials coupled in a dental implant can be found in a paper by Sikora et al. [11]. The specific aim of this research was to compare two types of artificial saliva as a lubricant in titanium–zirconia contact and both mechanical (through tribological test ball on disc) and chemical stability (through zeta potential titration curves) of the lubricant film were tested. This investigation is of relevance both for the wear testing of materials and in the case of the clinical use of artificial saliva against chronic dry mouth.

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

The discs used for tribological tests were made of commercially pure titanium (ASTM B348, Gr2, Titanium Consulting and Trading, Florence, Italy) and Ti6Al4V alloy (ASTM B348, Gr5, Titanium Consulting and Trading, Florence, Italy) and obtained from cylindrical bars (10 mm in diameter and 2 mm in thickness). Prior to the tests, the discs were polished using abrasive paper with different grit values (320–2000) and then ultrasonically cleaned in ethanol for 10 min twice and dried with compressed air. The balls used were made of zirconium oxide (ZrO2) partially stabilized with yttrium

(Y-TZP; Ceratec Technical Ceramics BV, Geldermalsen, The Netherlands; composition: 94.7 wt.% ZrO2, 5.3 wt.% Y2O3; 6 mm in diameter).

The aim of the research was to compare artificial saliva of different compositions. To do so, two commercial products were used in the form of liquids: Biotène (Biotène® Moisturizing Spray) and BioXtra (Biopharm). Both fluids contain water, xylitol, sodium benzoate, paraben (propylparaben and methylparaben or sodium propylparaben), and sodium saccharin. Moreover, Biotène contains organic corrosion inhibitors: polyethylene glycol 60 hydrogenated castor oil, cetylpyridinium chloride, and the polymer vp/va copolymer. Finally, Biotène contains glycerin, xanthan gum, limonene (pH = 6.06), while BioXtra contains hydrogenated starch hydrolysate, sorbitol, hydroxyethylcellulose, sodium monofluorophosphate, potassium chloride, sodium chloride, magnesium chloride, dipotassium phosphate, calcium chloride, colostrum whey, lactoperoxidase, citric acid, and potassium sorbate (pH = 7.33). No organic corrosion inhibitor is present in BioXtra.

Tribological tests were performed by using a high frequency reciprocating rig instrument (TR-282, Ducom Instruments Pvt. Ltd., Bengaluru, India) with the following parameters: Load = 2 N; Frequency = 30 Hz; Stroke length = 1 mm; Velocity = 60 mm/s; Temperature = 37 ◦C; Duration = 60 min; Lubrication volume = 2 mL. The lubricant (artificial saliva) was measured with a syringe and poured into the reservoir of the disc holder, completely covering the sample. The test was repeated on three different discs (and balls) for each pair of materials and saliva lubricant: the set of experiments allowed us to independently test the different variables. The aim was to independently compare the effect of the different saliva on the same substrate (Ti c.p. or Ti6Al4V) and of the same type of saliva on different substrates. The results were statistically analysed by using the Student's t-test.

Two-dimensional images of the discs and balls were acquired both before and after each test using optical microscopes (AmScope, ME300 Series and Olympus Vanox-T, Leica Microsystems B.V., Amsterdam, The Netherlands, respectively, for observing the discs and balls) in order to measure the length and width of the wear scars (measurements performed by ImageJ software (version 1.52). The samples after tribological tests were rinsed and wiped with ethanol before imaging.

Adsorption of saliva before zeta potential titration and contact angle measurements were performed by placing discs in a 12-well container and each one was covered with 0.5 mL of saliva. The adsorption phase lasted 1 h at 37 ◦C in the incubator. Subsequently, the samples were left to air dry for two days.

An electrokinetic analyser (SurPASS, Anton Paar, Graz, Austria) was employed for zeta potential titration vs. pH on the discs after saliva absorption. The surface zeta potential was determined in function of pH in a 0.001 M KCl electrolyte solution, varying the solution pH by the addition of 0.05 M HCl or 0.05 M NaOH through the instrument automatic titration unit. The acid and alkaline sides of the curve were obtained in two different steps on two different couples of samples.

Surface wettability by water was determined on the discs after the zeta potential tests through contact angle measurements (sessile drop method, DSA-100, KRÜSS GmbH, Hamburg, Germany). Briefly, a drop (5 μL) of ultrapure water was deposited on the material surface and the angle formed with the surface measured by the instrument software.

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

The coefficient of friction (COF) was measured during the tribological tests of zirconia balls on Ti c.p. or Ti alloy discs under lubrication with different artificial saliva: Biotène (containing organic corrosion inhibitors) and BioXtra (without organic corrosion inhibitors). Figure 1 shows the average friction coefficient as a function of the sliding distance for the four tested material/saliva combinations.

**Figure 1.** Average COF over sliding distance of Ti c.p. (**a**) and Ti6Al4V alloy (**b**) in the presence of Biotène (blue) or Bioxtra (red).

Compared to the BioXtra saliva, Biotène decreased the friction coefficient on both substrates with good lubrication action, as can be seen from the values shown in Table 1: the coefficients of friction of both materials were much lower in the presence of Biotène and they do not differ significantly from each other (*p* > 0.05, Student's t-test). The difference between the average friction coefficient of Ti c.p. in the presence of Biotène or BioXtra as a lubricant was statistically significant (*p* < 0.01); the same observation can be made in the case of Ti6Al4V. If we consider BioXtra as a lubricant and we compare Ti c.p. with Ti6Al4V, the difference of the average friction coefficients is again statistically significant (*p* < 0.01). Considering Biotène as a lubricant, the difference between the average friction coefficient of the two substrates was statistically less evident (*p* < 0.05). Furthermore, Biotène showed, by far, lower variability in the friction coefficient. In fact, there was a clear difference in the curve trend: samples lubricated with BioXtra had an initial region with gradual increase of COF, which means a "running-in period" that was much more extensive than those lubricated with Biotène. In this region, an adjustment of the two contacting surfaces occurred by crushing and smearing of the asperities. Then, a second region can be identified, during which COF remains fairly stable in the presence of Biotène, while the COF of the samples lubricated with BioXtra exhibited strong oscillations. After the running-in period, these oscillations may be attributed to the build-up and accumulation of third-body particles in the contact region.

**Table 1.** Average friction coefficients and standard deviations of all the tested samples.

A picture of the wear scar was acquired from each disc at the end of each test and a representative set of images is reported in Figure 2 for each of the tested materials.

**Figure 2.** Images of the wear scars on Ti c.p and Ti6Al4V after the tests with Biotène and BioXtra.

In the presence of Biotène (Figure 2), the shape of the scars on both materials was very similar and was almost a perfect ellipse. The edges were regular and well defined. The surface on the bottom of the scars was uniform and almost free of debris. Instead, in the presence of BioXtra (Figure 2), the scars on the two materials resulted in a more irregular elliptical shape. The edges were more jagged and the surface on the bottom of the scars was rough with a considerable presence of debris. In both cases, it can be seen that the grooves on the bottom of the scars showed an oriented texture aligned to the sliding direction. This can be explained by third-body abrasion processes caused by the formation of metal debris due to the ultra-hardness of zirconia.

The area of wear scars in the presence of Biotène was smaller than that in the presence of BioXtra (Figure 3); in particular, that on the titanium alloy was the lowest of them all, according to the higher mechanical resistance of the Ti6Al4V alloy with respect to Ti c.p.

**Figure 3.** Average areas of the wear scars on the discs.

An image of the wear scars on the zirconia balls was taken at the end of each test and a representative picture is reported below for each tribo-pair in Figure 4.

**Figure 4.** Images of a zirconia balls used against Ti cp and Ti6Al4V discs after the tests with Biotène and BioXtra.

It can be immediately noticed that the wear tracks on the zirconia balls tested in the presence of Biotène had a similar aspect and the same for those tested with BioXtra. In all cases, the wear scars formed on zirconia balls had an elliptical shape with a darker and more worn-out central area. The occurrence of a strong adhesive interaction between the tribo-pairs lubricated with BioXtra was supported by the higher amount of material transfer on the zirconia balls.

The mean wear scar diameter (MWSD) was calculated by averaging the two main diameters of wear scars on the zirconia ball (Figure 5). The MWSDs registered on Ti c.p. were not statistically different if the scars obtained in the tests with the two different lubricants are compared; no statistical difference also existed between the two substrates tested by using BioXtra as the lubricant. On the other hand, this difference between the MWSDs was statistically significant on Ti6Al4V and if the two substrates tested with Biotène are compared (*p* < 0.01)

**Figure 5.** Average mean wear scar diameters on the zirconia balls.

At the end of the tests, the amount of metallic debris dispersed in the artificial saliva used as a lubricant was qualitatively observed (data not reported); once again, the difference between the two salivas was evident with a greater quantity of residues visible at the end of the tests carried out with BioXtra.

It therefore appears that Biotène decreases the wear between the tested materials; this is due first of all to the lower friction coefficient than that in case of BioXtra. Furthermore, the presence in this artificial saliva of organic corrosion inhibitors (surfactants PEG-60 hydrogenated castor oil, cetylpyridinium chloride, and the polymer vp/va copolymer) can justify the lower tendency to wear. In fact, the mechanism of action of organic corrosion inhibitors is based on the adsorption on the surface to form a protective film that protects it against deteriorating [12]. In the presence of BioXtra, a strong

adhesive interaction has occurred between zirconia and Ti c.p. or the Ti alloy. Hence, the predominant wear mechanism is adhesive galling. It is likely that the surfactants, absent in this artificial saliva and present in Biotène, provide some protection to titanium surfaces from galling.

In this research, in addition to the mechanical tests, the chemical stability of the lubricant film on the titanium surfaces was tested by means of zeta potential titration curves. This is of interest, considering that different pH values can be present in the mouth and that wettability of the lubricant film and chemical interaction at the lubricant–surface interface are critical in determining the biotribological behaviour. The graphs below (Figure 6) show the zeta potential titration curves of the Ti c.p. and Ti alloy discs after adsorption of the two artificial salivas.

**Figure 6.** Zeta potential titration curves of Ti c.p. (**a**) and Ti6Al4V alloy (**b**) as received (green curves), after adsorption of Biotène (blue curves) and of BioXtra (red curves). The arrows mark the pH used at the beginning of the measurements and the acidic (on the left) and alkaline (on the right) ranges of the titration.

Comparing the curves of the samples pre-adsorbed with Biotène with the curves of the as-received substrates (i.e., tested without saliva pre-adsorption), we note that the curves obtained on the samples after adsorption were significantly different from the curves of the as-received materials. This means that the adsorption of some compounds from saliva occurred. The slope of the titration curves after adsorption was significantly lower than that of the as-received materials, which means that hydrophilicity of the surfaces after adsorption is higher than that of the bare metals. There was a plateau at pH higher than 6.5 due to functional groups (such as, for instance, OH groups) exposed on the surface by the adsorbed compounds, meaning that the functional groups on the surfaces have a specific acidic strength and are all completely dissociated when the pH is higher than 6.5, that means that in physiological condition there is a net negative surface charge. According to their dissociation only in the basic range, these groups do not act as a strong acid and they do not significantly affect the isoelectric point on the curve with respect to the Ti as it is (without adsorption of saliva). On the other hand, when strong acidic or alkaline functional groups are present on a surface, they are respectively de/protonated at lower/higher pH values than 7 and both the isoelectric point and the onset of the plateau are shifted to the left/right side.

Observing the curves related to the adsorption of BioXtra, it can be said that the curves obtained after adsorption were significantly different from the curves of the as-received materials, which means that the adsorption of saliva compounds also occurred in this case. The adsorbed layer had zeta potential values not far from 0 across the explored range: this behaviour is typical of a super-hydrophilic surface that does not adsorb either OH− or H3O+ ions instead of water molecules, even if the solution is strongly basic or acidic. On the other hand, the curve was not stable with a high standard deviation at some points and fluctuating trend, according to the presence of an un-stable surface layer. In addition, the step registered at pH 5.5 (that is, the first point of the two titration curves in the basic and acidic range) can be related to some instability, or not perfect reproducibility of the adsorbed layer.

These results can be used for interpreting the results obtained from wear tests. Compared to Bioxtra, Biotène decreased the friction coefficient and wear of both substrates because it allows for the formation of a chemically stable adsorbed layer with a good hydrophilic behaviour, acting as a good lubricant. On the other hand, the layer formed by BioXtra is not chemically stable and it is not able to act as a good lubricant, even if it has a super-hydrophilic behaviour.

This hypothesis was confirmed by contact angle measurements performed with a water drop on the discs after zeta potential tests, as reported in Figure 7. A consistent decrease in the contact angle with respect to the as-received Ti surfaces (without any adsorption of saliva) was measured on the sample pre-adsorbed with Biotène and used for the zeta potential test, which was in agreement with the zeta potential titration curve where a higher wettability was detected after adsorption of Biotène and confirmed that this saliva was still firmly adsorbed after the electrokinetic zeta potential test. On the other hand, the contact angle was almost unchanged with respect to the as-received titanium surfaces (without any adsorption of saliva) on the sample pre-adsorbed with BioXtra and used for the zeta potential test, confirming that the saliva film was almost completely removed after the electrokinetic zeta potential test and was not stably adsorbed on the titanium surfaces.

**Figure 7.** Contact angle measurements of water on Ti c.p. and Ti6Al4V alloy as received and after adsorption of the Biotene and BioXtra artificial salivas.

The obtained results are of interest because they set the relevance of the composition of the selected artificial saliva on the wear resistance of dental materials. Even if this topic has been already partially explored in the literature [13–15], the new evidence derived from this research are here discussed. Friction coefficient and wear were significantly reduced in the presence of organic corrosion inhibitors able to make a chemically and mechanically stable hydrophilic adsorbed layer. First, this means that researchers planning tribological tests have to carefully select the artificial saliva used as a lubricant in order to avoid overestimating the wear resistance of the tested surfaces. Second, it means that patients who need to use artificial saliva and have dental implants have to carefully select the right one. Finally, the use of zeta potential titration measurements is proposed here as a useful experimental tool for obtaining more information for a better understanding of the tribological tests: presence/absence of an adsorbed layer of lubricant on a tribological surface, its chemical and mechanical stability when significant pH changes occur and a flux of liquid is applied on the surface, the presence of functional groups exposed, and hydrophilicity of the surface after adsorption of the lubricant. The chemical instability of the adsorbed layer of saliva at pH 5.5 in the case of Bioxtra may be relevant in physiological conditions considering that acidic pH can be easily reached in the oral environment.

On the basis of this research, other investigations will be performed in order to overcome the limitations of the tests previously described including the investigation of the Ti6Al4V-ELI alloy (Gr23) as a substrate because it is currently widely employed in dental applications, better investigation of the adsorbed film thickness, comparison of artificial saliva with different additives (such as the antibacterial

ones), chemical analysis of the adsorbed compounds, different geometry of the samples, and tribological tests (such as flat-on-flat tests) that better simulate the tribological contact in dental implants.

#### **4. Conclusions**

Results suggest that the use of an artificial saliva containing organic corrosion inhibitors (such as Biotène) can significantly reduce the friction coefficient and wear both of titanium and zirconia surfaces. This is because it is able to be firmly adsorbed on the surface of the Ti c.p. or Ti6Al4V alloy and form a protective film with high wettability. In the absence of this film (such as in the case of BioXtra), much more severe wear occurs with a galling mechanism. In this last case, the occurrence of a super-hydrophilic saliva film that is not firmly adsorbed on the surface (mainly in an acidic environment) is not efficient in order to avoid severe wear.

From the methodological standpoint, the use of electrokinetic zeta potential and contact angle measurements together with standard tribological tests has shown to be useful for a better understanding of the lubricant action of different artificial saliva on different materials.

This research can be the basis for larger investigations of the tribological behaviour of dental materials and implants in the presence of lubricants with different compositions.

**Author Contributions:** Conceptualization, F.A., A.T. and S.S.; methodology, F.A. and A.T.; investigation, F.A., A.T. and V.P.; writing—original draft preparation, S.S.; writing—review and editing, S.S.; supervision, S.S. All authors have read and agreed to the published version of the manuscript.

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

**Acknowledgments:** The authors would like to acknowledge Ducom Instruments Europe B.V. for the friction and wear measurements.

**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/).

## **Influence of Conditions for Production and Thermo-Chemical Treatment of Al2O3 Coatings on Wettability and Energy State of Their Surface**

#### **Mateusz Nied ´zwied ´z, Władysław Skoneczny \* and Marek Bara**

Institute of Materials Engineering, Faculty of Science and Technology, University of Silesia in Katowice, 40-007 Katowice, Poland<EMAIL_ADDRESS>(M.N<EMAIL_ADDRESS>(M.B.)

**\*** Correspondence<EMAIL_ADDRESS>Tel.: +48-32-368563

Received: 26 June 2020; Accepted: 13 July 2020; Published: 15 July 2020

**Abstract:** This article presents the influence of the anodizing parameters and thermo-chemical treatment of Al2O3 coatings made on aluminum alloy EN AW-5251 on the surface free energy. The oxide coating was produced by DC (Direct Current) anodizing in a ternary electrolyte. The thermo-chemical treatment of the oxide coatings was carried out using distilled water, sodium dichromate and sodium sulphate. Micrographs of the surface of the Al2O3 coatings were characterized using a scanning microscope (SEM). The chemical composition of the oxide coatings was identified using EDS (Energy Dispersive X-ray Spectroscopy) microanalysis. Surface free energy (SFE) calculations were performed by the Owens–Wendt method, based on wetting angle measurements made using the sessile drop technique. The highest value of surface free energy for the only anodized coatings was 46.57 mJ/m2, and the lowest was 37.66 mJ/m2. The contact angle measurement with glycerine was 98.06◦ ± 2.62◦, suggesting a hydrophobic surface. The thermo-chemical treatment of the oxide coatings for most samples contributed to a significant increase in SFE, while reducing the contact angle with water. The highest value of surface free energy for the coatings after thermo-chemical treatment was 77.94 mJ/m2, while the lowest was 34.98 mJ/m2. Taking into account the contact angle measurement with glycerine, it was possible to obtain hydrophobic layers with the highest angle of 109.82◦ ± 4.79◦ for the sample after thermal treatment in sodium sulphate.

**Keywords:** aluminum oxide layers; surface morphology; thermo-chemical treatment

#### **1. Introduction**

Nowadays, when a common criterion for material selection is its weight to strength ratio, and when great importance is attached to environmentally friendly solutions, aluminum has become very widely used in material engineering [1]. The most important advantages of aluminum are undoubtedly its low weight in relation to strength, high electrical and thermal conductivity, ease of forming in all machining processes such as rolling and extrusion, considerable corrosion resistance and almost 100% recyclability without losing its properties [2–4]. In its pure form, aluminum, in addition to the number of advantages, also has significant disadvantages, which are its low melting point and low absolute strength [5]. In order to minimize these disadvantages, aluminum is combined with elements such as copper, silicon or magnesium to form durable alloys [6]. The industries that most often use aluminum and its alloys are the automotive, aviation and shipbuilding industries [7–10].

In order to better protect aluminum alloys against mechanical damage and corrosion, the process of so-called anodizing (oxidation) is used [11,12]. Anodizing aluminum and its alloys is an electrochemical process consisting of creating an oxide coating on the surface of an alloy, usually with a thickness of a few to several dozen micrometers [13]. The anodized coating is many times harder than pure aluminum or aluminum alloys (1700 to 2000 HV Al2O3 hardness at about 30 HV for pure aluminum and 200 HV for PA9 alloy) [14,15]. During the electrochemical process, the formation of the oxide coating occurs at the expense of substrate loss. This results in very good adhesion of the coating to the aluminum substrate. The most commonly used type of anodizing is hard DC (Direct Current) anodizing. During the process, as the coating thickness increases, the anodizing voltage increases [16]. The beginning of the DC anodizing process is characterized by a rapid increase in voltage, necessary to break through the compact and thin (0.01–0.1 μm thick) natural barrier layer, also called the barrier layer [17]. In the next stage of the process, the voltage decreases. When the minimum value is reached, the barrier layer is rebuilt at the oxide–electrolyte interface. The voltage increases again over time due to the formation of pores and their further deepening as the thickness of the Al2O3 coating increases, up to a maximum thickness of approximately 150 μm [18,19]. The oxide coating produced during anodizing ideally reaches the form of regular pores arranged in the shape of a hexagonal system, along with a thin barrier film in direct contact with the aluminum substrate. The oxide pore sizes mainly depend on the anodizing method used, but also on the process parameters. A high temperature and long anodizing time increase the solubility of the oxide, while creating a porous coating with poorer protective properties [20,21].

For better corrosion protection and sealing of the absorbent oxide coating while maintaining its advantages, thermo-chemical treatment is used [22]. As a result of processing in the Al2O3 coating, aluminum oxide is transformed into hydrated forms such as γ-Al (OH)3 hydrargillite and γ-AlOOH boehmite. The transformation occurs due to swelling of the walls of the Al2O3 cells owing to their hydration, as a result of which the pores are closed, and a smooth surface is created [23]. During thermo-chemical treatment, the process of producing a pseudo-boehmite sublayer on the surface may also occur; this happens most often when the process is short (several minutes), but also at high temperatures [24].

Researchers have investigated the thermo-chemical treatment of oxide coatings, but the research lacks reference to the impact of this process on the wettability of the Al2O3 layers [25,26]. The porosity of oxide coatings has a significant impact on their energy state, and thus the properties enabling the repulsion (hydrophobicity) and attraction (hydrophilicity) of water molecules [27]. That is why studies on the wettability of oxide coatings produced as a result of anodizing and thermo-chemical treatment are so important. Nowadays, the energy state is a very important feature of materials. Materials with hydrophobic properties, i.e., those exhibiting low wettability, repelling water particles, are widely used in surface engineering as self-cleaning or anti-icing materials [28,29]. Materials with high wettability are also in high demand, or, in other words, attracting water particles (hydrophilic); they are used in photovoltaic panels, biomaterials and wherever it is important to reduce friction [30–32]. A surface is considered hydrophobic when the water contact angle is greater than 90◦. The conventional name for surfaces reaching contact angles over 150◦ is ultrahydrophobicity. The opposite of hydrophobicity is hydrophilicity and occurs when the contact angle on the surface is less than 90◦ [33].

Many authors have carried out research on the modification of Al2O3 layers by thermo-chemical treatment. In [34], 6061 aluminum alloy was anodized at a constant voltage in an electrolyte at the temperature of 273 K. A 15% H2SO4 solution was used as the electrolyte, the anodizing time was 60 min, and a voltage of 15–30 V was used. Then, each of the produced layers were modified by immersion in stearic acid at the temperature of 353 K for 45 min, successively in ethanol at the temperature of 343 K. The whole process was completed by drying in an oven at the temperature of 253 K. The authors of the work managed to obtain hydrophobic surfaces; the contact angles were 138◦–152◦ depending on the anodizing voltage used. In subsequent studies, researchers conducted studies using two-stage anodizing on 1050 aluminium alloy at the temperature of 273 K. As the electrolyte, 10.5% and 1% solutions of H3PO4 were used; the process was carried out using the constant voltage method at voltages in the range of 160–195 V for 60 min in Stage 1 and 15–300 min in Stage 2. After Stage 1 of anodizing, the layers were immersed in a solution of 6% H3PO4, 1.8% CrO3 and deionized water at the temperature of 333 K. Wet chemical etching was carried out in a 5% solution of H3PO4 at the temperature of 333 K for 60–150 min. In the next step, the layers were sonicated in acetone for 10 min,

then dried in an oven at the temperature of 343 K for 360 min and modified by immersion in a 5% C12H24O2 solution for 90 min. The last stage of layer modification was spraying with silane and dried at the temperature of 413 K for 30 min. Thanks to the use of lauric acid, the contact angles were higher by 5◦–30◦, while thanks to silane spraying, a hydrophobic coating with a contact angle of 146◦ was obtained [35]. In the next cited article, researchers also dealt with constant voltage anodizing using a C2H2O4 electrolyte. The anodizing time was 420 min at the voltage of 40 V. The layers were modified by immersion in an 0.3% polytetrafluoroethylene solution and then cured at room temperature. In the next stage, the samples were immersed in a solution of 1.8% H2CrO4 and 6% H3PO4 at the temperature of 338 K for 300 min. Contact angles of 120◦ before immersion and 160◦ after immersion in a polytetrafluoroethylene solution were obtained [36]. Another article that can be cited is research based on two-stage constant voltage anodizing. Anodizing was carried out at the temperature of 275 K in a 1% H3PO4 solution. A voltage of 194 V was applied for 60 min in the 1st stage and 300 min in the 2nd stage. After completing Stage 1, the layers were immersed in a 6% H3PO4 solution and a 2% H2CrO4 solution at the temperature of 323 K for 40 min. After anodizing was completed, the layers were immersion modified in deionized H2O at the temperature of 373 K for 1 min. This process mainly involved preparing the substrate for subsequent treatments. The samples were then dried at the temperature of 323 K and finally treated with hexamethyldisilane vapor (HMDS) for 240 and 540 min. As a result of the performed modifications, contact angles of 139◦–153◦ were obtained (depending on the number and duration of HMDS cycles) [37].

In all the cited publications, the researchers focused mainly on the modification of oxide coatings produced as a result of constant voltage anodizing. In the next publication, the coatings were prepared using the DC method for 120 min at a current density of 1–2 A/dm2. The process was carried out in a 5% H3PO4 solution at a temperature of 283–293 K. Some of the samples were plasma treated and all of them were immersed in trichloro octadecylsilane-hexane for 120 min as well as being washed with hexane and oven dried at the temperature of 333 K. Wetting angles of 152◦ (without plasma treatment) and 157◦ (after plasma treatment) were obtained [38].

After analyzing the literature, it was found that most of the world research on the wettability of Al2O3 coatings is mainly based on the modification of the coating by applying additional layers to their surfaces. There is a lack of research in this area regarding the impact of the anodizing parameters before and after applying thermo-chemical treatment and the impact of the compounds used for this treatment on the contact angles, and thus the energy state of the surface of the coatings, which makes our research innovative. The studies presented below focused mainly on the impact of the anodizing parameters (current density, electrolyte temperature, anodizing time) and the impact of the thermo-chemical treatment carried out with various compounds (water, sodium dichromate, sodium sulphate) on the contact angles of oxide coatings. Measurement of the contact angles allowed the energy state of the layers to be determined by calculating the surface free energy.

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

#### *2.1. Research Material*

The research material was Al2O3 coatings made on aluminum alloy EN AW-5251 by means of hard DC anodizing. The aluminum alloy used has a high magnesium content, which contributes to such properties as good ductility, high corrosion resistance and high weldability. This alloy is most often used in the construction of aircraft and marine structures.

The samples for producing the oxide coatings had a surface area of 0.1 dm2. Before proceeding with the anodizing process, each sample was etched in a 5% KOH solution for 20 min, followed by a whitening process in a 10% HNO3 solution for 5 min. The temperature of the used solutions was 296 K. At the end of each process, the samples were rinsed in distilled water. The coatings were produced in the DC anodizing process using a GPR-25H30D power supply. The process was carried out in variable production conditions (current density, anodizing time), at a variable electrolyte temperature

(anodizing without thermo-chemical treatment) and a constant electrolyte temperature for the samples subjected to thermo-chemical treatment after the anodizing process. A ternary electrolyte consisting of an 18% sulphuric acid solution (33 mL/L), oxalic acid (30 g/L), and phthalic acid (76 g/L) was used for the anodizing process. During anodizing, the electrolyte was mixed mechanically at 100 rpm, the direction of rotation was changed every 10 min.

The studies were carried out according to a statistically determined poly-selective experiment plan. Hartley's plan was applied for three input factors based on a hypercube, for which coefficient α = 1. Table 1 presents the list of factors on a natural and standard scale for coatings produced in the anodizing process without thermo-chemical treatment.

**Table 1.** List of factors on natural and standard scale for coatings produced in anodizing process without thermo-chemical treatment.

Table 2 presents the list of factors on a natural and standard scale for coatings produced in the process of anodizing and thermo-chemical treatment. The thermo-chemical treatment was carried out after thorough rinsing of the samples in distilled water. The anodizing process was carried out at the constant electrolyte temperature of 298 K. Thermo-chemical treatment was carried out in solutions of distilled water, sodium dichromate and sodium sulphate. For each compound, the treatments lasted 60 min, the temperature of the solutions was 371 K.

**Table 2.** List of factors on natural and standard scale for coatings produced in anodizing process and thermo-chemical treatment.

#### *2.2. Research Methodology*

The thickness of the oxide coatings was measured by the contact method using a Dualscope MP40 instrument (Helmut Fischer GmbH + Co.KG, Sindelfingen, Germany), which uses the eddy current method for measurement. Ten measurements were taken along the entire length of the samples with three repetitions and the mean values were calculated along with deviations.

Using a Hitachi S-4700 scanning microscope (Hitachi, Tokyo, Japan), a micrograph of the morphology of the Al2O3 coatings was made. A magnification of 50,000× was used to observe the surface. Before further studies were carried out, the surfaces were sprayed with carbon to remove electrons struck during interaction of the microscope beam. Analysis of the chemical composition of the oxide coatings was carried out using a Noran Vantage EDS (Energy Dispersive X-ray Spectroscopy) system attached to the Hitachi S-4700 scanning microscope.

Contact angle measurements were performed using four liquids, two polar (water and glycerine) and two non-polar (α-bromonaphthalene and diiodomethane). Ten drops of each liquid were applied to each of the investigated coatings using an 0.5 μL micropipette over the entire length of the sample. Photographs of each drop were taken with a camera, which were then exported to a computer. The software used allowed the marking of three extreme points on the drop photograph, which enabled automatic measurement of the contact angle. The smallest and largest angles were rejected, and the other eight angle values were used to calculate the average contact angle of a particular coating. Surface free energy was calculated by the Owens–Wendt method using one polar (water) and one non-polar liquid (α-bromonaphthalene). Calculations using the Owens–Wendt method were carried out based on the equation:

$$
\gamma\_s = \gamma\_s^{\text{cl}} + \gamma\_s \mathbb{P} \tag{1}
$$

where γs is the surface free energy of the solid, γs d is the surface free energy (SFE) dispersion component of the examined materials, and γs p is the SFE polar component of the investigated materials.

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

Table 3 presents the average results of thickness measurements of the oxide coatings produced in the anodizing process (without thermo-chemical treatment).

**Table 3.** List of Al2O3 coating thicknesses produced in anodizing process (without thermo-chemical treatment).

The measurements showed significant changes in the thickness of the oxide coatings resulting from the use of different anodizing parameters. The thickness of the coatings depends on the process time, current density and electrolyte temperature. An increase in anodizing time at a constant current density and electrolyte temperature causes a significant increase in the coating thickness (samples 01I, 01J, 01K). As the current density increased, while the electrolyte time and temperature remained constant, a significant increase in coating thickness was also observed (samples 01E, 01F, 01K). In both cases, the increase in coating thickness is due to the increasing value of the electric charge. An increase in electrolyte temperature, while maintaining a constant anodizing time and current density, contributes to a decrease in coating thickness, which is caused by an increase in Al2O3 secondary solubility along with an increase in electrolyte temperature (samples 01G, 01H, 01K).

Table 4 presents the average results of thickness measurements of oxide coatings produced in the anodizing process and thermo-chemical treatment.

**Table 4.** Summary of Al2O3 coating thickness producedin anodizing process and thermo-chemical treatment.

Comparing the measurements of the thickness of the oxide coatings before and after the thermo-chemical treatment produced under the same anodizing conditions, it can be stated that the treatment causes a slight increase in the thickness of the oxide coating (comparison of samples 01I, 01J, 01K with samples 02 I, 02J, 0.2 K). Confirmation of this can also be seen by comparing samples 01E, 01F, 01K with samples 02E, 02F, 02K. The increase in the layer thickness after thermo-chemical treatment is insignificant and varies within 1.5 μm. As in the case of unmodified coatings, after the thermo-chemical treatment, the dependence of the increase in Al2O3 thickness on the increase in current density and the time of the anodizing process can also be seen. Table 5 shows the contact angle values measured using distilled water and α-bromonaphthalene. Table 6 shows the contact angle values measured using glycerine and diiodomethane. The measurements were made on coatings produced during anodizing without thermo-chemical treatment. Depending on the anodizing parameters, different contact angle values were obtained. The highest contact angle for polar liquids (distilled water, glycerine) was obtained for the 01D sample produced at the current density of 4 A/dm2 for 90 min at the electrolyte temperature of 303 K. The lowest value was obtained for the 01J sample produced at the current density of 3 A/dm2 for 90 min at the electrolyte temperature of 298 K. Considering the values of the contact angles of the coatings measured using glycerine, sample 01D is a hydrophobic surface with a contact angle of 98.06◦ ± 2.62◦.

Table 7 contains the contact angle values measured using distilled water and α-bromonaphthalene. Table 8 shows the contact angle values measured using glycerine and diiodomethane. The Al2O3 coatings underwent thermo-chemical treatment after previous anodizing. The highest value of contact angle measured with distilled water was obtained for sample 02E produced at the current density of 2 A/dm2 for 60 min, then subjected to thermo-chemical treatment in sodium sulphate. The smallest contact angle value was obtained for the 02D sample produced at the current density of 4 A/dm2 for 90 min, subjected to thermo-chemical treatment in sodium dichromate. The produced coating was highly hydrophilic—the measured angle was only 8.62◦ ± 2.02◦. Considering the second polar liquid (glycerine), the smallest contact angle value was also measured for the 02D sample; however, the largest value was as much as 109.82◦ ± 4.79◦—the highly hydrophobic surface was measured on the 02K sample produced at the current density of 3 A/dm2 for 60 min and then subjected to thermo-chemical treatment in sodium sulphate. Sample 02E also showed hydrophobic properties—its contact angle was 92.78 ± 6.62◦. Thermo-chemical treatment of the Al2O3 coatings contributed to a significant change in

the wettability of the measured surfaces. In the case of the sample anodized at the current density of 3 A/dm2 for 60 min, the thermo-chemical treatment in sodium sulphate contributed to an increase in the contact angle from 73.86◦ to 81.14◦ for water and from 71.95◦ to 109.82◦ for glycerine, creating a highly hydrophobic surface. Furthermore, the thermo-chemical treatment of samples 01E, 01F, 01K in sodium sulphate contributed to a significant reduction in wettability (a higher contact angle) for glycerine. A reduction in the electrolyte temperature during anodizing from 303 to 298 K and the thermo-chemical treatment process in sodium dichromate of samples 01C, 01D, 01H contributed to a very large increase in the wettability of the coatings, creating surfaces with strong hydrophilic properties, both using water and glycerine. An increase in the electrolyte temperature during anodizing from 293 to 298 K for samples 01A, 01B, 01G and the thermo-chemical treatment process using water resulted in a significant reduction in the contact angle.

Table 9 presents the surface free energy values calculated by the Owens–Wendt method. Wetting angles measured using distilled water and α-bromonaphthalene using the sessile drop technique were used for the calculations. The angles were measured on Al2O3 coating without thermo-chemical treatment.

**Table 5.** Wetting angles of oxide coatings produced in the anodizing process (without thermo-chemical treatment) for distilled water and α-bromonaphthalene.

**Table 6.** Wetting angles of oxide coatings produced in the anodizing process (without thermo-chemical treatment) for glycerine and diiodomethane.

**Table 7.** Wetting angles of oxide coatings produced in anodizing process and thermo-chemical treatment for distilled water and α-bromonaphthalene.

**Table 8.** Wetting angles of oxide coatings produced in the anodizing process and thermo-chemical treatment for glycerine and diiodomethane.

**Table 9.** Surface free energy values of Al2O3 coatings without thermo-chemical treatment.

The highest value of surface free energy was determined for the 01J sample, i.e., for the coating characterized by the smallest contact angle for water, while the lowest SFE value was determined for the surface characterized by one of the highest contact angles for water (84.06◦)—sample 01G. After the calculations, it can be stated that on oxide coatings without thermo-chemical treatment, the contact angles show inverse proportionality to the value of surface free energy. Three-dimensional graphs were made to visualize the effect of the anodizing parameters on SFE (Figure 1).

(**c**)

**Figure 1.** Surface free energy (SFE) dependence on: (**a**) process time and current density, (**b**) electrolyte temperature and current density, (**c**) electrolyte temperature and process time.

By analyzing the impact of the process time and current density on SFE, it can be seen that the highest values of surface free energy of over 44 [mJ/m2] were exhibited by a coating produced within 100 min and at the current density of 3 A/dm2. Values over 40 [mJ/m2] were also observed for coatings produced within 30 min for all the current density values. The graph showing the effect of electrolyte temperature and current density on SFE shows a clear increase in the surface free energy value in the middle of the electrolyte temperature axis for a current density of about 2 A/dm2. The lowest values < 39 [mJ/m2] were attained for the current density of 3 A/dm2 and electrolyte temperature around 293 K. The next graph presents the dependence of SFE on the electrolyte temperature and process time. The highest values > 44 [mJ/m2] were found at the beginning and end of the process time axis at values in the middle of the electrolyte temperature axis (298 K).

Table 10 shows the SFE calculated for coatings after thermo-chemical treatment. The calculations were carried out using the Owen–Wendt method based on the contact angles of distilled water and α-bromonaphthalene.

**Table 10.** Values of surface free energy of Al2O3 coatings after thermo-chemical treatment.

Sample 02D showed the highest surface free energy; this is a coating which, after anodizing, was subjected to thermo-chemical treatment in sodium dichromate. The coating also has the lowest contact angle for water (8.62◦) and for glycerine (27.67◦). The lowest SFE value was calculated for sample 02E, which has the highest contact angle for water (89.61◦) and one of the highest for glycerine (92.78◦). As with the Al2O3 coatings without modification, the contact angles of the coatings after thermo-chemical treatment show an inverse proportionality to the value of surface free energy. Figure 2 presents 3D graphs of the dependence of the anodizing parameters and compounds used for thermo-chemical treatment on the SFE value.

The surface free energy of coatings after thermo-chemical treatment varies depending on the process time and current density. The highest SFE values > 75 mJ/m2 were found at the intersection of the end of the process time axis (about 90 min) and the end of the current density axis (about 4 A/dm2). A clear increase can also be seen at the intersection of a current density of about 2 A/dm2 and a process time of about 30 min. By analyzing the impact of the current density and the compound used for thermochemical treatment, it can be stated that the greatest impact on the increase in the SFE value is the use of sodium dichromate in the coating modification process. When using sodium dichromate, a clear increase in surface free energy is visible in the entire current density range, with a peak value > 76 mJ/m2 for a current density of about 4 A/dm2. A clear decrease in the SFE value can be seen for sodium sulphate at the current density of 2 A/dm2. The graph showing the dependence of SFE on the time of the process and the compound used for thermochemical treatment shows a clear maximum SFE value > 76 mJ/m2 for sodium dichromate, decreasing with the time of the process. Surface free energy values > 56 mJ/m2 are also visible in the case of anodized coatings that underwent thermo-chemical treatment in water for about 90 min. A clear decrease in the SFE value < 44 mJ/m2 is visible in the middle of the graph for anodized coatings modified in the sodium sulphate solution for about 60 min.

Figure 3 presents surface morphology micrographs taken with a scanning microscope of selected oxide coatings with the largest differences in the contact angle and SFE. A 50,000× magnification was used for all the morphology micrographs.

The micrographs show surface porosity, which is characteristic of Al2O3 coatings. There are significant differences in the size of the nanopores, which depend on both the anodizing time and the current density used in the anodizing process.

Figure 4 presents micrographs of the surface morphology of layers anodized and thermo-chemically treated in sodium dichromate and sodium sulphate.

The thermo-chemical treatment of the 02D sample carried out in sodium dichromate contributed to complete sealing of the porous Al2O3 coating and the formation of a high sodium content sublayer. On the surfaces of the 02E and 02K samples, which after the anodizing process were subjected to thermo-chemical treatment in sodium sulphate, complete coverage of the oxide coating surface with a kind of crosslinked structure is noticeable.

**Figure 2.** SFE dependence on: (**a**) process time and current density, (**b**) chemical compounds used in coating modification and current density, (**c**) chemical compounds used in coating modification and process time.

**Figure 3.** Surface morphology micrographs of oxide coatings without thermo-chemical treatment: (**a**) Sample 01D, (**b**) Sample 01H, (**c**) Sample 01J.

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

(**c**)

**Figure 4.** Surface morphology micrographs of oxide layers after thermo-chemical treatment: (**a**) Sample 02D, (**b**) Sample 02E, (**c**) Sample 02K.

Table 11 summarizes the chemical composition of the oxide coatings produced in the DC anodizing process without the thermo-chemical process.

Samples with significant differences in the values of surface free energy and contact angle were selected for analysis. The atomic content of both aluminum and oxygen in the selected coatings are very similar. An increase in aluminum content was observed for the samples with a smaller Al2O3 coating thickness. The increased aluminum content is due to the shorter distance between the coating surface and the substrate. The values presented in the table are very similar to stoichiometric calculations of the alumina chemical composition.

Figure 5 presents SEM micrographs for samples 02D, 02E, 02K together with the marked field used for chemical composition analysis. Figure 6 shows the EDS spectrum for the 02D, 02E, 02K samples with chemical composition analysis.

(**c**)

**Figure 5.** SEM micrographs with chemical analysis field marked: (**a**) Sample 02D, (**b**) Sample 02E, (**c**) Sample 02K.

The coating composition of sample 02D is characterized by a high content of oxygen, at the same time reducing the content of aluminum. Thermo-chemical treatment in sodium dichromate contributed to the building of over 18% sodium compounds on the surface. On the surfaces of the 02E and 02K samples modified in sodium sulphate, a reduction in oxygen content for about 2% sulphur was shown.

**Figure 6.** *Cont.*

**Figure 6.** EDS (Energy Dispersive X-ray Spectroscopy) spectrum along with chemical analysis for samples: (**a**) Sample 02D, (**b**) Sample 02E, (**c**) Sample 02K.

#### **4. Conclusions**

Based on the conducted studies, it can be concluded that both the anodizing parameters and thermo-chemical treatment have a significant impact on the wettability of Al2O3 coating surfaces, and thus their energy state. By controlling the anodizing parameters (current density, electrolyte temperature, process time), a hydrophobic surface with a contact angle of 98.06◦ ± 2.62◦ was obtained for the 01D sample using glycerine as the measuring liquid. The 01D sample oxide coating also achieved the highest contact angle with distilled water. The coating was produced at the current density of 4 A/dm2 for 90 min at the electrolyte temperature of 303 K. The lowest contact angle value was obtained by the coating produced at the current density of 3 A/dm2 for 90 min at the electrolyte temperature of 298 K—sample 01J.

The second part of the research focused on the modification of the Al2O3 coatings by chemical compounds through thermo-chemical treatment. A very clear effect of thermo-chemical treatment on the values of contact angles was noticed. The highest contact angle value measured using distilled water was obtained by sample 02E modified in sodium sulphate. The smallest contact angle value was obtained by a sample modified in sodium dichromate. The coating modified in sodium dichromate showed strongly hydrophilic values; the measured angle was only 8.62◦ ± 2.02◦. For the 02K sample modified in sodium sulphate, the contact angle measurements with glycerine revealed an angle of 109.82◦ ± 4.79◦ (a strongly hydrophobic surface). Moreover, sample 02E exhibited hydrophobic properties; the contact angle for this coating was 92.78◦ ± 6.62◦. The thermo-chemical treatment in sodium sulphate of the 02K sample increased the contact angle from 73.86◦ to 81.14◦ for water and from 71.95◦ to 109.82◦ for glycerine, creating a highly hydrophobic surface. The modification also contributed to a significant reduction in wettability for glycerine. The reduction in electrolyte temperature in the anodizing process of the coatings modified in sodium dichromate contributed to a very large increase in the wettability of the coatings, creating surfaces with strong hydrophilic properties, both when measuring with distilled water and glycerine. The increase in electrolyte temperature during anodizing from 293 to 298 K and the thermo-chemical treatment process using water resulted in a significant reduction in the contact angle measured for distilled water. The studies have also shown that the contact angle is inversely related to surface free energy.

**Author Contributions:** M.B. performed the analysis of the results; M.N. carried out the tests and wrote the manuscript; W.S. contributed to the concept and modified the manuscript. All authors have read and agreed to the published version of the manuscript.

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

**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/).

## **The Influence of Anodic Alumina Coating Nanostructure Produced on EN AW-5251 Alloy on Type of Tribological Wear Process**

#### **Mateusz Nied ´zwied ´z, Władysław Skoneczny \* and Marek Bara**

Institute of Materials Engineering, Faculty of Science and Technology, University of Silesia in Katowice, 40-007 Katowice, Poland<EMAIL_ADDRESS>(M.N<EMAIL_ADDRESS>(M.B.) **\*** Correspondence<EMAIL_ADDRESS>Tel.: +48-32-368563

Received: 11 December 2019; Accepted: 21 January 2020; Published: 24 January 2020

**Abstract:** The article presents the influence of the anodic alumina coating nanostructure produced on aluminum alloy EN AW-5251 on the type of tribological wear process of the coating. Oxide coatings were produced electrochemically in a ternary electrolyte by the DC method. Analysis of the nanostructure of the coating was performed using ImageJ 1.50i software on micrographs taken with a scanning electron microscope (SEM). Scratch tests of the coatings were carried out using a Micron-Gamma microhardness tester. The scratch marks were subjected to surface geometric structure studies with a Form TalySurf 2 50i contact profiler. Based on the studies, it was found that changes in the manufacturing process conditions (current density, electrolyte temperature) affect changes in the coating thickness and changes in the anodic alumina coating nanostructure (quantity and diameter of nanofibers), which in turn has a significant impact on the type of tribological wear. An increase in the density of the anodizing current from 1 to 4 A/dm2 causes an increase in the diameter of the nanofibers from 75.99 ± 7.7 to 124.59 ± 6.53 nm while reducing amount of fibers from 6.6 ± 0.61 to 3.8 ± 0.48 on length 1 × 103 nm. This affects on a change in the type of tribological wear from grooving to micro-cutting.

**Keywords:** aluminum oxide layers; nanostructure; tribological wear

#### **1. Introduction**

Aluminum alloys are currently widely used in industry due to the low weight of the metal in relation to high mechanical strength and good thermal conductivity [1,2]. Due to the atmosphere, aluminum is automatically covered by a passive oxide coating a few nanometers thick, isolating the metal from contact with the environment [3]. However, in acidic or alkaline environments, the spontaneous oxide coating does not protect the metal well enough and aluminum corrosion occurs fairly quickly, leading to mass loss [4]. Therefore, ensuring proper protection of the aluminum surface involves the formation of an oxide coating of an appropriate thickness [5]. One method that allows adequate protection of aluminum alloys is electrochemical oxidation of the alloy surface. This process is called aluminum anodizing, and due to its properties that improve the hardness of metal, it is used on a large scale today [6]. Anodic alumina coatings in the electrochemical process can be formed using DC [7], constant voltage [8], pulse methods [9], and using AC [10], as well as AC imposed on DC [11]. Electrochemically produced oxide coatings have an amorphous structure [12]. The anodic alumina coating adopts a porous structure with vertically aligned cylindrical pores, depending mainly on the applied current density, electrolyte temperature, and aluminum alloy [13,14]. An important feature of anodic alumina coatings is their resistance to abrasive wear. This results in a wide application of anodized aluminum in friction pairs of engineering kinematic systems [15]. For this reason, tests are often carried out to assess the impact of the anodizing parameters on the mechanical and tribological

properties of oxide coatings [16]. One of the tests determining the mechanical properties of the coatings is the scratch test. Such tests allow one to not only determine the coefficient of friction value but also to determine the value of critical load causing damage to the coating [17,18]. In the first cited article, scientists investigated scratches of anodized aluminum in 10% oxalic acid at room temperature at 10–40 V. Subsequently, the crystal structure, chemical composition, surface morphology, and surface topography were examined. Studies have shown that higher anodizing potential leads to the formation of porous, thicker, and harder anodic alumina coatings. Anodized aluminum showed the formation of horizontal and vertical parallel cracks when attempting to scratch at a load of about 1 N. In the second article cited, researchers investigated the mechanical properties of AZ91 magnesium alloy coated with a double-layer Al2O3/Al. Mechanical properties were examined, among others, by a scratch test. It has been found that the mechanical properties of the coated alloy are significantly better, and the coating exhibits excellent adhesion to the substrate. Nanostructure studies are an important element of research related to the wear of oxide coatings and their adhesion to the substrate [19]. In the cited article, the authors started to produce Al2O3 layers by hard anodizing and thermo-chemical treatment. Then, they studied the microstructure, morphology, and tribological wear. Microstructure tests showed fiber sealing after physicochemical treatment, and it also increased the microhardness and polymer consumption during the tribological test. Our work focuses on the scratching of aluminum surfaces and anodic oxide coatings produced using various manufacturing parameters. In addition to the scratch resistance test, nanostructure analysis was performed to show the relationship between changes in the nanostructure and their impact on the type of tribological wear of the coatings which has not yet been studied. Scientists studied the nanostructure and friction of oxide layers in sliding running using finite elements [20]. They focused mainly on the oxidation of aluminum alloy for 60 min in an electrolyte consisting of sulfur and oxalic adipic acid at 298 K and a current density of 3 A/dm2. The tribological properties of the oxide layer during reciprocating friction were tested using four different polymers (TG15, TGK20/5, TMP12, PEEK/BG). The highest wear and the highest friction coefficient were found for TMP12 material, while the lowest wear and friction coefficient for PEEK/BG. However, all cited articles lack reference to nanostructure for the type of tribological process, which makes our research innovative.

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

#### *2.1. Research Material*

The research material was anodic alumina coatings produced by an electrochemical method on the surface of plates made of EN AW-5251 aluminum alloy. This aluminum alloy is characterized by high mechanical strength and corrosion resistance. The chemical composition of the EN AW-5251 alloy (Table 1) allows one to easily modify its surface by anodizing.

**Table 1.** Chemical composition of the EN AW-5251 alloy.

Samples with an area of 0.1 dm2 were subjected to etching using a 5% KOH solution (20 min) and a 10% HNO3 solution (5 min). The temperature of the solutions was 296 K. The etching treatments ended with rinsing in distilled water. The sample surfaces (except for the reference sample) were modified by anodizing to form an anodic alumina coating on their surface using an electrochemical method. The surface of the samples was anodized by the DC method, using a GPR-25H30D power supply. During anodizing, a variable current density and electrolyte temperature were used to determine how the process parameters affect the type of tribological wear. A ternary electrolyte constituting an aqueous solution of 18% sulfuric acid (33 mL/L), oxalic acid (30 g/L), and phthalic acid (76 g/L) was used for the anodizing process. Anodic hard coatings produced in electrolytes constituting of an aqueous solution sulfuric or oxalic acid require the use of low process temperatures (273–278 K). The addition of phthalic acid ensures that hard layers are obtained and enables the process to be carried out at room temperature. The anodizing process conditions are shown in Table 2.

After anodizing, the samples were rinsed in distilled water.

#### *2.2. Research Methodology*

The thickness of the anodic alumina coatings was measured by the contact method using a Fischer Dualscope MP40 instrument (Helmut Fischer GmbH+Co.KG, Sindelfingen, Germany) using the eddy-current method. Ten measurements were made (repeated 3 times) over the entire length of the sample, then the average values were calculated. Nanostructure tests of the oxide coatings were carried out on metallographic specimens using a Hitachi S-4700 scanning microscope (Hitachi, Tokyo, Japan) at 30,000× magnification. Anodic alumina coatings are non-conductive material, so they charge electrically when scanning through an electron beam, which prevents correct observation of the preparations. For better observation of the nanostructure, the samples were sputtered with carbon, which enables the discharge of electrons rebound during research. The micrographs were used for computer analysis of the image carried out using ImageJ 1.50i software. Using the image analysis procedures (smooth, bandpass filter, threshold), the number of fibers per nm and the diameter of the fibers were calculated. The scratch test was performed using a Micron-Gamma device (MicronSystema, Kiev, Ukraine). A Rockwell diamond indenter was used with a tip radius of 0.2 mm. A load of 4 ± 0.01 N was used during the test. The scratch marks were subjected to surface geometric structure (SGS) analysis. For this purpose, a Form TalySurf Series 2.50i profiler was used. The systematic scanning method was used on the transverse profile. The SGS parameters allowed the authors to determine the type of tribological wear.

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

Measurements of the thickness of the oxide coatings showed significant differences in the thickness depending on the anodizing conditions. Averaged values of measurements of the anodic alumina coating thickness together with deviations are presented in Table 3.

**Table 3.** Influence of anodizing conditions on thickness of oxide layers formed on EN AW-5251 aluminum alloy surface.

The current density as well as the electrolyte temperature during the anodizing process affects the thickness of the anodic alumina coatings. Considering the current dependence, one can notice a significant increase in coating thickness at a constant electrolyte temperature and anodizing time together with an increase in current density (Samples A and D). An increase in the thickness of the coating along with an increase in current density (at a constant process time) is associated with an increase in electric charge. The anodic oxidation process, like any electrochemical process, proceeds according to Faraday's law, according to which a specific amount of electricity transforms aluminum into a specific amount of its oxide. Taking into consideration the temperature dependence at low current density values (Samples A and B), as the electrolyte temperature increases, the coating thickness increases. The increase in electrolyte temperature accelerates the migration of ions, which in turn increases the growth rate of the coating. However, at higher current density values (Samples C and D), the coating thickness decreases as the electrolyte temperature increases. Depending on the type of electrolyte in which the electrolysis process is carried out, a certain amount of produced Al2O3 is always dissolved by the electrolyte. In the thicker coatings, due to the resistance of the electrolyte column lying in the pores located between the oxide fibers, Joule's heat is released. The reduction in the Al2O3 layer thickness as the electrolyte temperature rises, at higher current densities, can therefore be attributed to the increasing secondary solubility of aluminum oxide associated with the higher electrolyte temperature.

Micrographs of the oxide coating nanostructures taken on metallographic specimens at 30,000× magnification (Figure 1) showed structure oriented along the direction of growth of the coating under the influence of an electric field.

**Figure 1.** SEM micrograph (cross-section) of nanostructure of the coatings: (**a**) Sample A, (**b**) Sample B, (**c**) Sample C, (**d**) Sample D (sample designations as in Table 2).

In order to assess the impact of the coating production conditions on their nanostructure, the thickness and number of coating fibers were measured. For this purpose, each of the micrographs was subjected to image analysis, which enabled sharpening and visualization of the nanofibers. The first procedure used was the smooth function (first column), thanks to which the amount of noise in the images was significantly reduced. Then, the bandpass filter function (second column) was used, whose task was to filter the image in a banded way. The last function used was the threshold function with the dark back option selected (third column), thanks to which binarization of the images was achieved—only black and white colors appear (the nanofibers are shown in black). The micrographs of the coatings after applying the appropriate procedures are shown in Figure 2.

**Figure 2.** SEM micrograph (cross-section) of nanostructure of the coatings: (**a**) Sample A, (**b**) Sample B, (**c**) Sample C, (**d**) Sample D (sample designations as in Table 2), 1—image with smooth function, 2—image with bandpass filter function, 3—image with threshold function with dark back option selected.

The obtained results of the values of the aluminum oxide fiber diameters and the number of fibers on length 1 × 103 nm are presented in Table 4.

**Table 4.** Average values of diameters and number of coating fibers.

The relationship between the coating thickness and the diameter of the aluminum oxide fibers is given by Equation (1).

$$\mathbf{y} = 7.83\mathbf{x} \text{ - 3.8, }\tag{1}$$

*y*—coating thickness,

*x*—average fibers diameter.

On the basis of the image analysis data obtained from ImageJ 1.50i, histograms of the nanofibers with appropriate diameter groups in individual oxide coatings were made (Figure 3).

**Figure 3.** Histograms of nanofibers: (**a**) Sample A, (**b**) Sample B, (**c**) Sample C, (**d**) Sample D (sample designations in Table 2).

Computer analysis of the coating micrographs showed several significant relationships between the production parameters and the number of nanofibers and their diameters. For a constant current density of 1 A/dm2, an increase in electrolyte temperature causes a decrease in the number of nanofibers per 1 × 103 nm while increasing their diameter (Samples A and B). In turn, at the current density of 4 A/dm2 (Samples C and D), the increase in electrolyte temperature causes an increase in the number of nanofibers while reducing the diameter of the fibers. The above relationships most likely result from the increasing secondary solubility of aluminum oxide fibers with increasing temperature at higher current density values. At a constant electrolyte temperature, an increase in current density (Samples A and D) causes a significant reduction in the number of nanofibers while increasing their diameter. Another important relationship is the linear relationship between the average diameter of the Al2O3 nanofibers coatings and the thickness of these coatings—an increase in the thickness of the oxide coating increases the diameter of the nanofibers. The coatings produced under extreme conditions (Samples A and D) with the largest and smallest thicknesses are also characterized by the largest and smallest number of nanofibers.

Figure 4 shows the transverse profiles of the surface of the oxide coatings and aluminum alloy made with the profiler after the scratch test. The area above the zero point is marked in green (material upsetting on both sides of the scratch), and in red the area after the scratch was made.

**Figure 4.** Transverse profiles of oxide coating and aluminum alloy surfaces after scratch test: (**a**–**d**): Sample designations in Table 2, (**e**): Aluminum alloy EN AW-5251.

The images showing the transverse profiles of the samples show areas both below and above the zero point, whose surface areas change depending on the coating production parameters. The shape of these areas also undergoes change. The aluminum alloy exhibited the most rounded area profiles. The increase in current conditions of the anodizing process leads to "sharpening" of the area profiles, which is caused by the increasing thickness of the oxide coatings and rising hardness. Table 5 shows the values of the ratio of the cross-sectional surface area of material upsetting around Scratch *f1* and the recess of Scratch *f2*. Based on these values, the processes of anodic alumina coating wear were determined.

**Table 5.** Values of ration of cross-sectional surface area of material upsetting around Scratch *f1* and recess of Scratch *f2.*

Analysis of parameters *f1* and *f2* allowed the authors to determine the surface wear process of the tested samples. It was noticed that increasing the thickness of the anodic alumina coating (starting with pure aluminum) reduces the *f1*/*f2* dependence. As a result of the analysis, it was found that both in the case of pure aluminum and Sample A (smallest coating thickness) there is only plastic deformation of the surface layer, i.e., grooving. The tested surface is indented by the presence of unevenness or abrasive grain, which causes the material to move outside the surface. Grooving occurs if the value of the ratio *f1*/*f2* > 1. In the remaining tested samples, there is scratching and microcutting, which in turn occurs in the case of 0 ≤ *f1*/*f2* ≤ 1. Roughness measurements carried out before the scratch test showed slight differences between layer surfaces. The Ra parameter value was 0.35 ± 0.035 μm. Slight differences in roughness do not affect the type of layer wear, however.

#### **4. Conclusions**

Based on the conducted research, it was found that the anodizing process parameters of the aluminum alloy EN AW-5251 surface affect the nanostructure of the anodic alumina coatings, whose construction, in turn, affects the type of tribological wear. An increase in current density while maintaining a constant electrolyte temperature results in the creation of an anodic alumina coating with a significantly increased diameter of coating nanofibers, while reducing their quantity. A linear relationship between the increase in the thickness of the oxide coating and the increase in diameter of the nanofibers was also demonstrated. A reduction in the *f1*/*f2* ratio value was also observed as the diameter of the nanofibers increased. From among the tested samples, only the surface of the aluminum alloy and anodic alumina coatings of the smallest thickness and the diameter of the nanofibers underwent grooving. The increase in production parameters, and hence the thickness of the coatings causes a change in the wear process of the coating from grooving to scratching and microcutting. Given the structure of the layers (fiber size, coating thickness depending on the density of the electric charge), the above studies are in line with previous scientific studies. Due to the fact that there is no literature research on the impact of nanostructure on the speed and mechanism of wear, we are not able to refer to literature.

**Author Contributions:** M.B. performed the analysis of the results; M.N. carried out the tests and wrote the manuscript; W.S. contributed to the concept and modified the manuscript. All authors have read and agreed to the published version of the manuscript.

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

**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*

## **Investigation of the Oxidation Mechanism of Dopamine Functionalization in an AZ31 Magnesium Alloy for Biomedical Applications**

## **Arezoo Ghanbari 1, Fernando Warchomicka 2,\*, Christof Sommitsch 2 and Ali Zamanian 1,\***

Received: 9 July 2019; Accepted: 12 September 2019; Published: 16 September 2019

**Abstract:** Implant design and functionalization are under significant investigation for their ability to enhance bone-implant grafting and, thus, to provide mechanical stability for the device during the healing process. In this area, biomimetic functionalizing polymers like dopamine have been proven to be able to improve the biocompatibility of the material. In this work, the dip coating of dopamine on the surface of the magnesium alloy AZ31 is investigated to determine the effects of oxygen on the functionalization of the material. Two different conditions are applied during the dip coating process: (1) The absence of oxygen in the solution and (2) continuous oxygenation of the solution. Energy dispersive spectroscopy (EDS) and Fourier transform infrared spectroscopy (FTIR) are used to analyze the composition of the formed layers, and the deposition rate on the substrate is determined by molecular dynamic simulation. Electrochemical analysis and cell cultivation are performed to determine the corrosion resistance and cell's behavior, respectively. The high oxygen concentration in the dopamine solution promotes a homogeneous and smooth coating with a drastic increase of the deposition rate. Also, the addition of oxygen into the dip coating process increases the corrosion resistance of the material.

**Keywords:** biodegradable magnesium; dopamine; Impedance behavior; molecular dynamic simulation

#### **1. Introduction**

Magnesium implants have intrigued the medical industry as a way to substitute temporary implants with more permanent counterparts. In the application of temporary fasteners, including screws or pins, magnesium implants are superior to permanent titanium or stainless-steel implants, since their use eliminates the need for second surgery to remove the implant after bone healing. Moreover, magnesium alloys provide mechanical properties with close compatibility with bone, thereby decreasing the risk of stress shielding [1,2]. As with other bio metals, magnesium is investigated to promote the cell's response on the surface. The healing rate of bone is critical [3], and, for successful bone regeneration, cell viability is highly significant [4]. Cell adhesion is unlikely to occur directly on the implant. Thus, an appropriate surface treatment is needed [5]. For this treatment, several biomimetic coatings are investigated: Aminopropyltriehtoxysilane, vitamin C (AV), carbonyldiimidazole (CDI), stearic acid (SA), aminated hydroxyethyl cellulose, gelatin, and peptides [5–9]. Depending on the composition, biomimetic coatings can have the capability to retard the corrosion affinity of magnesium to body fluids [5,8,9]. Therefore, a biomimetic coating should be adhesive and compact enough to

protect the substrate against corrosive body fluid. In this area of study, polydopamine is introduced as a biomimetic coating to enhance the adhesion of other biomimetic coatings [10]. For instance, the use of a coating made by gelatin (a mixture of peptides and proteins) and modified with polydopamine can improve the binding affinity between the coating and the metal (e.g., on a titanium surface [11]) or promote corrosion resistance [12]. In other cases, the use of dopamine helps to immobilize functional molecules as a robust anchor in biofunctional surfaces (e.g., nano magnetic particles [13]) or increase the cytocompatibility and osteogenic potential of the scaffold [14], among other examples.

Recent works on magnesium have shown the advantage of using a functionalized surface with dopamine to enhance corrosion resistance [15–17], as well as the cell proliferation on the surface of the implant [17,18]. Inspired by the composition of adhesive proteins in mussels, dopamine can be self-polymerized to form a surface-adherent, thin, and robust film on almost all kinds of materials, including organic and inorganic ones [19]. In this area, the dip coating method [15–17,20] has been used as a convenient way to anchor the dopamine and create a film on magnesium alloys. In situ spontaneous deposition takes place through an aerobic auto-oxidation mechanism in a mildly basic aqueous buffered media and, consequently, a bio-inspired film is generated on the substrate. Although the film is known as polydopamine, there is controversy over whether the film consists of supramolecular of monomer aggregates or has a polymeric structure [21]. The oxidation mechanism during the dopamine coating and the evolution of the molecules have been deeply investigated [18,20,22–25]. Many oxidants were evaluated as a buffer for the formation of the coating, including compounds like copper sulfate, sodium periodate, and ammonium peroxodisulfate [25,26], and oxygen as a biocompatible oxidant [20,23,25,26]. It is reported that polydopamine or dopamine- melanin film deposition seems to depend not only on the monomer's availability but also on the availability of O2 [16,25].

Whether polydopamine is incorporated in a magnesium implant as an adhesive intermediate biomimetic layer or an individual biomimetic layer, the corrosion rate of this layer is fundamental. The present work focuses the investigation on variation of the corrosion resistance of the AZ31 alloy through functionalization of the surface via dopamine, with the presence and absence of oxygen. Oxygen is selected as a condition variable to investigate the effects of an oxidant on the coating's quality.

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

#### *2.1. Substrate*

The magnesium alloy AZ31 was used in this work. The chemical composition of AZ31, in wt.%, is as follows: 3 Al, 0.4 Mn, 0.8 Zn, max. 0.005 Fe, 0.004 Cu, 0.003 Ni, 0.07 Si, max. 0.01 Cr, and balance Mg. The as-received AZ31 was in a cold rolled and partially annealed condition.

#### *2.2. Coating Process*

Sheet samples of 1 × 1 cm2 were polished on both sides with SiC abrasive papers (up to 4000 grit) to obtain a uniform roughness. Samples were washed in acetone for 10 min in an ultrasonic bath to remove any surface contaminations. At the end of the surface preparation, samples were washed in deionized water and immediately dried in the air.

For the dip coating, dopamine hydrochloride ((HO)2C6H3CH2CH2NH2·HCL), hydrochloric acid (HCl), and tris(hydroxymethyl)aminomethane, also known as Tris, ((HOCH2)3CNH2) were purchased from Merck (Darmstadt, Germany) as raw materials. The coating solution was prepared using 2 mg/mL dopamine hydrochloride, 100 mL deionized water, and 48 mM Tris as a buffer. In order to understand the oxidant effect on the coating, two different conditions were applied during the dip coating process: (1) the absence of oxygen in the solution and (2) the solution continuously being oxygenized with a flow rate of 0.4 L/min. In both conditions, the pH of the solution was adjusted to 8.5 and was kept constant by adding adequate amounts of hydrochloric acid with a concentration of 1 M. The pH value was monitored with a digital pH meter (Metrohm, Herisau, Switzerland).

The substrate was immersed in the solution for two hours in darkness [17,27]. After immersion, the samples were removed quickly to inhibit further oxidation and cleaned in deionized water to remove any unattached dopamine on the surface. Coated samples were finally dried in a nitrogen flow to avoid any undesirable surface oxidation. In this work, samples functionalized via dopamine with the absence of oxygen are designated as "dopa", while samples immersed in a solution rich in oxygen are named "dopa-O", and the pristine substrate is "AZ31".

#### *2.3. Characterization*

The morphological feature observations, accompanied with a chemical analysis of the pristine and functionalized AZ31 alloys, were carried out using a field emission scanning electron microscope (FESEM, Mira3 Tescan, Brno, Czech Republic) equipped with an energy dispersive spectroscope detector (EDS: EDAX- AMETEK, Tilburg, The Netherlands). FESEM images were taken in secondary electron mode to monitor the topography of the layer. Chemical analysis was carried at the surface and at the cross section to estimate the homogeneity of the deposited polydopamine. Moreover, Fourier transform infrared spectroscopy (FTIR, Thermo; Waltham, MA, USA) in transmittance mode was used to verify the presence of polymerized dopamine on the surface of the "dopa" and "dopa-O" samples.

#### *2.4. Electrochemical Analysis*

The degradation behavior of the samples was investigated via electrochemical impedance spectroscopy (EIS; AutoLab Potentiostat, Metrohm, Herisau, Switzerland). Samples were stabilized in phosphate-buffered saline (PBS) solution prior to the tests, and all impedance measurements were performed in a PBS solution at a temperature of 37 ◦C. A three-electrode electrochemical cell setup, with an Ag/AgCl electrode as a reference electrode and a platinum sheet as a counter electrode, was utilized in this work. Moreover, samples were used as working electrodes. EIS scans were acquired from 100 kHz to 0.01 Hz, and the impedance data were analyzed with the NOVA software (version 1.11).

#### *2.5. Cell Cultivation*

A human osteosarcoma cell line (G292), taken from National Cell Bank of Iran, was selected to evaluate the cellular behavior of samples. Osteoblast-like cells were cultured in a humidified atmosphere of 5% CO2 at a temperature of 37 ◦C using Dulbecco's modified Eagle's medium (DMEM, BIO-IDEA, Iran) supplemented with 10% fetal bovine serum (FBS) (Gibco, Thermo, Waltham, MA, USA) and 1% antimicrobial penicillin/streptomycin solution 100X (Biosera, Kansas City, MO, USA). The cells were subcultured in a cell culture medium, which was refreshed every 3 days before utilization. Two series of samples were placed in 6-well plates and sterilized by ultraviolet light in 70% (*V*/*V*) ethanol–water solution and kept in sterilized media before analysis. Afterwards, samples were seeded with the same number of G292 cells (105 cells) in a culture medium, and the cells were cultivated for 72 h in an incubator. Then, the samples were gently rinsed with phosphate-buffered saline solution (PBS) to remove the unattached cells.

Microscopic observations were carried out via both FESEM and fluorescent microscope to observe the cell morphology and adhesion to the samples. For the FESEM observation, the cells were fixed with a 2.5 v % glutaraldehyde solution for 2 h in a dark environment. Subsequently, the samples were gently washed again with PBS and then immersed in a bath of 50%, 60%, 70%, 80%, 90%, and 100% (*V*/*V*) ethanol / water solutions for 30 min each time to dehydrate the cells attached to their surfaces. The dehydration conducted in absolute ethanol was performed twice, and each time the samples were immersed in 100% ethanol for 1 h. The morphology of the G292 cells attached to the samples was observed after drying and sputtering with gold.

The cell behavior was also examined by a fluorescence microscope (Olympus, Japan) using a cell staining method. A total of 0.1% acridine orange (AO, sigma, Saint Louis, MO, USA) was utilized to visualize the remaining cells on the substrates [28], staining the cells for 1–2 min.

#### *2.6. Molecular Dynamics Simulation*

The effect of the oxidant on polydopamine coating was studied via molecular dynamics simulations. Two different conditions were designed, simulating the "dopa" (condition one) and "dopa-O" (condition 2) deposition conditions. In both conditions, the magnesium matrix was simulated via EAM (embedded atom model) potential. Moreover, each condition consisted of a mixture of dopamine monomers, indolequinone monomers, and dimmers. Concerning the oxidation mechanism of dopamine, the different ratios of dopamine to indolequinone were considered for each deposition condition: 90/10 for "dopa" and 10/90 for "dopa-O". Molecular dynamics simulations were performed via a large-scale atomic/molecular massively parallel simulator (LAMMPS). The atomic computations were carried out in periodic boundaries and in three dimensions. The generic force field, Dreiding, was employed to describe the inter-molecular interactions of the polymer. In this case, the valence interactions were considered to be bond stretch, bond-angle bend, and dihedral angle torsion terms [29]. The non-bonded or Van der Waals interactions of polymer/polymer and polymer/Mg are given by the Lennard–Jones [30] potential. A total of 10,000,000 time-steps with an increment of 0.01 fs were simulated for each condition. The temperatures of the conditions were controlled using a Langevin thermostat accompanied by a micro-canonical ensemble NVE. Visualizations were created using the Ovito software.

#### **3. Results**

#### *3.1. Characterization of the Coating*

The surface observation and the elemental map analysis of "AZ31", "dopa", and "dopa-O" conditions are displayed in Figure 1. In the as-received "AZ31"condition, intermetallic phases rich in aluminum are observed on the polished surface, which is reported as a typical secondary phase Mg17Al12 in AZ31 alloy [31,32]. Figure 1b, c shows the functionalized surface of the "dopa" and "dopa-O". Both samples are homogeneously covered with a coating, which exhibits the presence of nitrogen and oxygen on the Mg matrix, in concordance with observations in the literature, together with carbon and hydrogen elements in the polydopamine layer [24]. At the surface of the coating, cracks are formed during the drying stage [33].

A cross section of the coated samples is analyzed in Figure 2 to determine the thickness of the layer, and the variation of the elements depends on the dip condition. The thickness of the coatings was 42 and 85 μm for "dopa" and "dopa-O", respectively, in agreement with both SEM observations and element profile scans. The element scan line shows a strong influence of the type of coating on the amount of nitrogen and oxygen elements in the layers. The "dopa-O" has higher intensity than the "dopa" condition. Furthermore, the presence of Mg in the layer of "dopa-O" may be related to the oxidation of the substrate due to the flow of oxygen during the coating process. The element profile in Figure 2c shows that although the amount of element is lower in "dopa", it is more homogenous in composition than that the "dopa-O". The first 20 μm (Region I) keeps a constant composition for both conditions, while the reduction of N through the coating in "dopa-O" (region II) indicates a lack of polydopamine in this area, which has been replaced by magnesium oxide.

**Figure 1.** The SEM micrographs and relative elemental mapping of (**a**) the as-received AZ31 alloy; **(b**) "dopa" and (**c**) "dopa-O". Samples functionalized via dopamine with the absence of oxygen are designated as "dopa", while samples immersed in a solution rich in oxygen are named "dopa-O".

**Figure 2.** SEM micrographs and relative elemental mapping of the cross section (**a**) "dopa" and (**b**) "dopa-O"; (**c**) elemental line scan of "dopa" and "dopa-O" coatings, marked in the micrographs in (**a**) and (**b**), respectively.

FTIR characterizations detect the functional groups in the polydopamine layer. The functional groups of the formed polymeric layer on the surface of "dopa" and "dopa-O" are depicted in Figure 3. A peak at 702 cm–1 can be observed in the "dopa" curve, which implies the ring deformation vibration

of 1,2,4-trisubstituted benzene. In both conditions, the peaks at 800 cm–1 (and also the band broadening between 800 and 860 cm–1) can be attributed to the C–H wagging vibration of the ring hydrogens in the 1,2,4-trisubstituted and 1,2,4,5-tetrasubstituted benzenes. The band accompanied with peaks in the range of 950–1420 cm–1 is ascribed to ring deformation vibrations and C–N stretching vibrations [34]. The adsorption bands located in the range of 1450–1625 cm–1 originate from the C=C stretching vibration of benzene rings and the N–H bending [27]. In addition, a peak at 1637 for both conditions is attributed to the C=C stretching vibrations, while the "dopa-O" condition shows a peak at 1729 cm–1 related to the C=O stretching vibrations [34]. A broad absorbance in the spectrum between 3100 and 3600 cm–1 is ascribed to N–H/O–H stretching vibrations [33], and in this region, NH and NH2 stretching vibrations happen in the range of 3100–3360 cm–1 [34].

**Figure 3.** The Fourier transform infrared spectroscopy (FTIR) spectrum of "dopa" and "dopa-O".

#### *3.2. Electrochemical Properties*

The effect of oxygen on the corrosion resistance of the coated samples was analyzed by impedance tests in a PBS solution at 37 ◦C. The variation of the real impedance component (*Z'*) versus the imaginary impedance component (*Z"*), the Nyquist plot, is shown in Figure 4a. The Bode plots (given in Figure 4c,d) improve the accuracy of the degradation behavior. Herein, the consistency of the obtained EIS data are evaluated with the Kramers–Kronig (K–K) transformation. The impedance spectra of all samples are characterized in high and medium frequency regions, with one semicircle loop in the Nyquist plot. Moreover, the one phase extremum in the Bode plot (Figure 4c) is an indicator of one semicircle in the Nyquist plot (Figure 4a). Figure 4d illustrates the decrease in impedance when the frequency decreases. This process is influenced by the adsorbed intermediate species, which form layers of corrosion products [35].

To compare the impedance data of the samples in quantity, the impedance plots are fitted with an equivalent electrical circuit, which is depicted in Figure 4b. The semicircle in the high and medium frequency region is ascribed as the constant phase element (*CPE*), simulating the charge transfer resistance (*Rc*) in parallel with the electric double layer. Moreover, the inductive loop is simulated by the inductance (*L*) and the relative inductive resistance (*Rl*). The total circuit is in series with the solution resistance (*Rs*).

The impedance data extracted from the equivalent electrical circuit are given in Table 1. From the data, the quantity of the circuit elements is increased in the following order: "dopa-O" > AZ31 > "dopa". Due to the insufficient quality of the polymer coating, the impedance is decreased in the "dopa". However, the "dopa-O" sample shows the highest constant phase element, inductance, and real resistances. Apparently, the presence of oxygen gas flowing during the dopamine dip coating process improves the corrosion resistance of the functionalized AZ31 alloy.

**Table 1.** The extracted impedance data from equivalent electrical circuit. CPE, constant phase element.

**Figure 4.** Electrochemical data for samples; "AZ31", "dopa", and "dopa-O" (**a**) Nyquist plot; (**b**) Equivalent electrical circuit; (**c**) Bode-phase angle plot, and (**d**) Bode-impedance plot.

#### *3.3. Cell Cultivation*

The SEM micrographs of the G292 cells attached to the samples, after 72 h cultivation, are illustrated in Figure 5. Artifacts observed on the surface might be related to the sputtering with gold. The cells keep a rounded and spherical shape and show a slightly tendency to be spread out on the surface of the "AZ31" alloy and "dopa" (Figure 5a,b). However, the cells also show a tendency to spread out on the surface of the "dopa-O" condition, as illustrated in Figure 5c.

**Figure 5.** SEM micrographs of the adherent G292 cells, (**a**) "AZ31"; (**b**) "dopa"; and (**c**) "dopa-O" after 72 h of cultivation time.

Figure 6 shows the G292 cells attached on the surface of samples by fluorescent microscopy. The adhesion of the cells can be evidenced in all the cases and a change of the cell's behavior in the "dopa-O" condition is perceptible. However, this information is not enough to draw the effect of the coating on the cell's morphology due to the cytoskeleton/membrane cannot be observed properly. Further investigations focused on cell viability and proliferation are needed to better quantify the effect of the oxygen on the cells' behavior.

**Figure 6.** Fluorescent microscopy of adherent cells stained with acridine orange: (**a**) "AZ31"; (**b**) "dopa"; and (**c**) "dopa-O" after 24 h of cultivation time; (Magnification 100×).

#### *3.4. Molecular Dynamics Simulation*

Molecular dynamics simulations help to determine the effect of oxidants on the polydopamine of the AZ31 alloy substrate. The displacement of the polymer mass towards the substrate surface was simulated. Thus, the center-of-mass (com) command was computed in LAMMPS for Z direction. Additionally, two different conditions were implemented by LAMMPS, considering that the dopamine molecule transforms into indolequinone in the presence of oxygen [25], as described in Section 2.6. The dimmers were designed for both conditions due to the further interactions between the monomer molecules [36]. In fact, polydopamine is not a covalent bond polymer but instead an aggregate of monomers held together by noncovalent forces [24]. The designed molecular structures of dopamine and indolequinone monomers [25], along with the dimmers [24], are illustrated in Figure 7a–c. In addition, the visualized conditions (one and two) are illustrated in Figure 7d,e. The displacements of the center-of-mass (com) in the polymer were calculated for each time-step in the Z direction, shown in Figure 7f. The slope of the graphs represents the difference of the displacement per difference of time. Hence, the slope indicates the deposition rate (angstrom/femtosecond). In this work, the slope of the "dopa-O" condition is higher, showing that the presence of oxygen accelerates the deposition of the polymer on the substrate. It seems that the tendency of indolequinone to bond to the substrate is higher than that of dopamine during the coating process.

**Figure 7.** Molecular dynamic simulation: (**a**) Dopamine monomer and (**b**) indolequinone monomer; (**c**) dimers; (**d**) condition one; (**e**) condition two; (colors definition: green: Mg; gray: C; white: H; red: O and blue: N); (**f**) center of mass versus time-step.

#### **4. Discussion**

#### *4.1. Characterization of the Coating*

The oxidant can intensify polydopamine formation [18,25]. According to the EDS maps given in Figure 1, the surface of the "dopa" and "dopa-O" is covered homogeneously with oxygen and nitrogen. The thickness measurements given in Figure 2 indicate that the thickness of the coating layer in "dopa-O" is, relatively, twice that of "dopa", in agreement with the deposition rate estimated by molecular dynamic simulations (Figure 7f). From respective cross section EDS mapping, the difference in nitrogen and oxygen content is obvious. This difference can be further perceived from the element profile scans (Figure 2c). The coatings (in depth) show two main regions, where the depth of region I is relatively the same for both conditions (~20 μm). Region II shows more depth for the "dopa-O", and the presence of Mg in both layers might be related to the magnesium oxide formed during the immersion in the dip process. This effect is more noticeable in the "dopa-O" condition, where the

continuous flowing of gas can promote the formation and growth of the oxide at the surface. On the other hand, a difference between the two conditions with respect to the ratio of oxygen to nitrogen in weight percent (O/N) is pronounced. Region I indicates a ratio of around 3 for "dopa-O" and 1 for "dopa". The presence of polydopamine should show a ratio of O/N in dopamine and an indolequinone ratio of 2. For this reason, the presence of these molecules on the surface of "dopa" is more reliable in this work. Analysis of the polydopamine layer via FTIR spectrum (Figure 3) indicated that the NH and NH2 stretching vibrations take place in the wavenumber range, 3100–3360 cm–1 [34], and a higher absorbance of "dopa-O" is associated with the amount of nitrogen (Figure 2c). Through the oxidation mechanism of polydopamine, the dopamine molecules transform to indolequinone molecules, whose structure is benzene-like with different positions for the substitutes. Here, the three-substituted structure of dopamine changes to the four-substituted structure of indolequinone [25]. The peak at 702 cm–1 in "dopa" is due to the ring deformation vibration of 1,2,4-trisubstituted benzene, which is similar to the dopamine structure [34]. Thus, a high amount of dopamine that is not fully oxidized can still be expected in the coating of the "dopa". If oxidation and cyclization do not take place completely, dopamine molecules co-exist in the solution [37].

On the other hand, the peak at 800 cm−1 and the band broadening between 800 and 860 cm–1 are attributed to the C–H wagging vibration of ring hydrogens in 1,2,4-trisubstituted and 1,2,4,5 tetrasubstituted benzenes [34]. This peak and band are slightly higher in the "dopa-O" condition, which has a similar structure to 1,2,4,5-tettrasubstituted benzenes. This means that the presence of oxygen activates the oxidation mechanism, and more three-substituted benzenes (dopamine structures) transform into four-substituted benzenes (indolequinone structures). Moreover, the appearance of a peak at 1729 cm–1 for "dopa-O" is due to the carbonyl (C=O) stretching vibrations of the indolequinone [24,25,34] structure, showing that adding oxygen increases the ratio of the dopamine that transforms to indolequinone. This effect promotes the presence of polydopamine instead of dopamine molecules in the "dopa-O" condition.

#### *4.2. Electrochemical Properties*

A comparison of the impedance responses to the corrosive PBS solution (Figure 4a) shows that the largest semicircle is in the "dopa-O" condition, thereby validating the observations of Kim et al. in various commercial microporous membranes [38]. The amount of the constant phase element, representing the resistance to the adsorption/diffusion processes at the electrode surface [39], increased from 50.83 μMho/cm2 in "dopa" to 105.83 μMho/cm2 in "dopa-O". Thus, the impedance of the capacity behavior (CPE) is relatively doubled (see Table 1). This trend is also observed for real resistors. The charge transfer resistance increased from 31.44 Ω cm2 in "dopa" to 142.32 Ω cm2 in "dopa-O". Thus, the transferring of charges responsible for anodic and cathodic reactions is hindered in "dopa-O", so the formed polymeric layer on "dopa-O" acts as a stronger inhibitor of charge transfer and charge storage (CPE) than "dopa". Moreover, there is a considerable increase in inductance element values after surface functionalization in the presence of oxygen, representing greater prevention against specious movement, including Mg2+ and Mg(OH)2. On the other hand, the equivalent circuit component values of "dopa" are lower than those of "AZ31". It seems that the polymeric layer formed on the surface of "dopa" cannot act as an effective barrier to magnesium dissolution in the corrosive PBS solution.

The polymeric layer formed in the "dopa-O" deposition condition acts as a superior barrier layer to PBS corrosive media. As seen in the EDS line (Figure 2c), region I in "dopa-O" contains a higher O/N ratio. Further, according to the FTIR view, the functional layer groups indicate a greater oxidation state. Since oxidation is the polymerization mechanism, it can be deduced that the polymerization of dopamine takes place more frequently on "dopa-O". Through dopamine polymerization or polydopamine formation, aggregates are formed on the surface by linking the particles together [33]. Therefore, this layer can act as a better barrier than "dopa" against corrosive liquid penetration, with less oxidized dopamine in the layer. On the other hand, the presence of MgO in "dopa-O" might influence the corrosion resistance of the material.

#### **5. Conclusions**

The use of oxygen as an oxidant agent for the dopamine functionalizing process was considered in this work, using the magnesium alloy AZ31. The introduction of oxygen into the dip coating draws the following conclusions:

**Author Contributions:** Conceptualization, F.W., A.G. and A.Z.; methodology, A.Z.; software, A.G.; formal analysis, A.G. and F.W.; investigation, F.W. and A.G.; writing—original draft preparation, A.G., F.W. and A.Z.; writing—review and editing, A.G., F.W., C.S. and A.Z.; supervision, A.Z. and C.S.

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

**Acknowledgments:** Open Access Funding by the Graz University of Technology. Authors would like to thank R. Buzolin for FESEM observations, A. Bordbar-Khiabani for samples preparation, T. Ramezani for helping with cell cultivation and R. Sabetvand for supporting the molecular dynamics simulations.

**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/).

## **Tri-Functional Calcium-Deficient Calcium Titanate Coating on Titanium Metal by Chemical and Heat Treatment**

#### **Seiji Yamaguchi \*, Phuc Thi Minh Le, Morihiro Ito, Seine A. Shintani and Hiroaki Takadama**

Department of Biomedical Sciences, College of Life and Health Sciences, Chubu University, Aichi 487-8501, Japan **\*** Correspondence<EMAIL_ADDRESS>Tel.: +81-568-51-6420

Received: 6 August 2019; Accepted: 27 August 2019; Published: 3 September 2019

**Abstract:** The main problem of orthopedic and dental titanium (Ti) implants has been poor bone-bonding to the metal. Various coatings to improve the bone-bonding, including the hydroxyapatite and titania, have been developed, and some of them have been to successfully applied clinical use. On the other hand, there are still challenges to provide antibacterial activity and promotion of bone growth on Ti. It was shown that a calcium-deficient calcium titanate coating on Ti and its alloys exhibits high bone-bonding owing to its apatite formation. In this study, Sr and Ag ions, known for their promotion of bone growth and antibacterial activity, were introduced into the calcium-deficient calcium titanate by a three-step aqueous solution treatment combined with heat. The treated metal formed apatite within 3 days in a simulated body fluid and exhibited antibacterial activity to *Escherichia coli* without showing any cytotoxicity in MC3T3-E1 preosteoblast cells. Furthermore, the metal slowly released 1.29 ppm of Sr ions. The Ti with calcium-deficient calcium titanate doped with Sr and Ag will be useful for orthopedic and dental implants, since it should bond to bone because of its apatite formation, promote bone growth due to Sr ion release, and prevent infection owing to its antibacterial activity.

**Keywords:** antibacterial activity; bone growth; apatite formation; titanium; silver; strontium; calcium titanate; ion release; cytotoxicity; controlled release

#### **1. Introduction**

Titanium metal (Ti) and its alloy are widely used for orthopedic and dental implants since they fulfill certain clinical needs from the point of view of mechanical properties, durability, and biocompatibility. Osseointegration occurs when the surfaces of the metals were roughened at micrometer scale such as 0.5–2.0, or 3.6–5.6 μm in calculated average roughness *R*a, and 43–50 μm in maximum height *R*z by plasma spraying, grid blasting and/or acid etching [1,2]. The nanometer-scale roughness produced by anodic oxidation also has been found to increase cell adhesion, proliferation and alkaline phosphatase activity [3,4]. Although these roughened Ti surfaces are able to directly contact with living bone, they still do not bond to it adequately. Mineralization process is a promising method to achieve strong and stable bone-bonding. It has been reported that bioactive glass/ceramics such as bioglass, hydroxyapatite, and glass ceramics A-W directly bonded to bone through the bone-like apatite layer formed on their surfaces [5].

Various types of surface coating of bioactive glass/ceramics by plasma spray, sputtering, sol-gel, and alternative soaking have been attempted [6–8]. Among them, a plasma spray coating of hydroxyapatite has been widely used to confer bone-bonding to total hip joint, dental implant, and so on. However, this does not form a stable bioactive surface layer, since the surfaces of the hydroxyapatite particles exposed to the plasma are partially melted, so the resultant calcium phosphate coating is liable to be decomposed in the living body in process of time [6].

Various types of surface modifications including alkali/acid solution and heat treatment, hydrothermal treatment, and ion implantation have been developed to confer apatite-forming capability on the metals so that the activated metals form bone-like apatite spontaneously on their surfaces by using the calcium and phosphate present in body fluids, and thereby bond to bone through the apatite [9–12]. Among the variety of the modification techniques, the alkali/acid solution and heat treatment has the needed characteristics for producing a uniform activated surface layer, even on the inner wall of a porous body, without requiring any especial apparatus [13,14]. It has been demonstrated that a bioactive sodium titanate layer is produced on Ti when the metal is soaked in 5 M NaOH solution at 60 ◦C for 24 h and subsequently heat-treated at 600 ◦C for 1 h [15,16]. The surface of treated metal forms a bone-like apatite spontaneously in the living body and bonds bone through this layer [16]. Total artificial hip joints (THAs) with the bioactive sodium titanate layer on their porous Ti layer have been under clinical use since 2007. A recent ten-year follow-up revealed the beneficial effects of the NaOH-heat-treated THAs to be a high survival rate (98%), no radiographic signs of loosening, and both early and stable bone-bonding [17]. However, two joints were retrieved owing to deep infection and periprosthetic femoral fracture, since the NaOH-heat-treated THAs neither promoted bone growth nor prevented infection [17]. Subsequently, the NaOH-heat treatment was modified to NaOH-CaCl2-heat-water treatment to produce a calcium-deficient calcium titanate layer on Ti and Ti alloys, which resulted in more stable apatite formation and bone-bonding [18–20]. On the other hand, there are still the challenges of providing antibacterial activity and promoting bone growth on Ti. It has been reported that typically 1%–2% of patients with total hip arthroplasties suffer deep infections [21]. Dental peri-implant disease and infection have become a main focus of oral implantology [22].

Strontium (Sr) and silver (Ag) ion are candidates for the promotion of new bone formation and prevention of infection since the former exerts a therapeutic effect on osteoporosis bone due to increased new bone formation and deceased bone resorption, while the later prevents infection because of its strong effect against a broad range of microorganisms [23–26]. Pre-clinical study reports have shown that the Sr ions released from dosed strontium ranelate improve mineral density at various skeletal sites such as the total hip and lumbar spine, resulting in the improvement of osteoporosis [23,24]. The mechanism of the antimicrobial action of Ag ions is understood as resulting from an interaction with the thiol (sulfhydryl) groups in enzymes and proteins [25], and is effective even in the living body [26]. Studies have reported the separate incorporation of Sr or Ag into the surface of Ti and Ti alloys [27–30], but there are few reports of these ions being incorporated simultaneously. There are even reported attempts to incorporate these ions into the Ti surface in an effort to confer a capacity for apatite formation. It was reported that Ag-doped calcium phosphate coatings was produced on Ti by a combination of anodic oxidation, electrophoretic deposition, and magnetron-sputtering [31,32]. Although the coated metal exhibited strong antibacterial activity against *Escherichia coli* (*E. coli*), it a little decreased cell viability [32]. A novel method is desired to confer Ti the capabilities of excellent antibacterial activity without any decrease in cell viability, direct bone-bonding, and promotion of new bone formation at the same time.

In this study, Sr and Ag ions were introduced into the calcium-deficient calcium titanate produced on Ti under controlled conditions so that the treated Ti slowly released Sr and Ag ions in order to exhibit the functions of promoting new bone formation while preventing infection without decreasing apatite formation. The potential of the treated metal for clinical applications is discussed in terms of Sr and Ag ion release, antibacterial activity, cytocompatibility, and apatite formation.

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

#### *2.1. Surface Treatment*

Commercially pure Ti sections (Ti > 99.5%; Nilaco Co., Tokyo, Japan) 10 mm × 10 mm × 1 mm in size was grinded with #400 diamond plates and then cleaned in an ultrasonic bath by using acetone, 2-propanol and ultrapure water for 30 min, and dried at 40 ◦C overnight. They were immersed in

5 M NaOH (Reagent grade; Kanto Chemical Co., Inc., Tokyo, Japan) solution at 60 ◦C with shaking at 120 strokes/min for a period of 24 h followed by gentle rinse under ultrapure water flow for 30 s. The treated metals were soaked in a mixed solution consist of 50 mM CaCl2 (Reagent grade; Kanto Chemical Co., Inc., Tokyo, Japan) and 50 mM SrCl2 (Reagent grade; Kanto Chemical Co., Inc., Tokyo, Japan) at 40 ◦C, shaken at 120 strokes/min for 24 h, then washed and dried in a similar manner (designated as "Ca + Sr"). They were subsequently heated at 600 ◦C with programming rate of 5 ◦C/min and holding time of 1 h, then naturally cooled in an electric furnace. After the heat treatment, they were immersed in a mixed solution of 1 M Sr(NO3)2 (Reagent grade; Kanto Chemical Co., Inc., Tokyo, Japan) and *X* mM AgNO3 (Reagent grade; Kanto Chemical Co., Inc., Tokyo, Japan) with an adjusted pH from 3 to 8 by adding HNO3 or NH3(aq) at 80 ◦C, where *X* is a range from 1 to 100 mM and designated as "Sr + *X* mM Ag, shaken , washed, and dried in the in a similar manner. The nominal and measured pH of the 1 M Sr(NO3)2 + 1 mM AgNO3 are summarized in Table 1. Some of the Ti samples subjected to the same NaOH-Ca + Sr-heat treatment were subsequently soaked in 1 M SrCl2 solution without Ag for comparison.

**Table 1.** Measured pH of 1 M SrNO3 + 1 mM AgNO3 solution corresponding to nominal pH.

#### *2.2. Surface Analysis*

#### 2.2.1. Scanning Electron Microscopy and Energy Dispersive X-ray Analysis

The samples treated with the aqueous solution and heat were examined by field emission scanning electron microscopy (FE-SEM: S-4300, Hitachi Co., Tokyo, Japan) equipped with an energy dispersive X-ray spectrometer (EDX: EMAX-7000, Horiba Ltd., Kyoto, Japan). In FE-SEM observation, the samples were subjected to a thin-film coating of Pt–Pd, and 15 kV accelerate voltage was selected. In EDX anlaysis, the elements of Ca, Ag, O, and Ti were quantified using 9 kV. The measurement was repeated on five different points, and their averaged value was used.

#### 2.2.2. Thin-Film X-ray Diffraction and Fourier Transform Confocal Laser Raman Spectrometry

The surface structure of the samples subjected to the aqueous solution and heat treatment were analyzed by a thin-film X-ray diffractometer (TF-XRD: model RNT-2500, Rigaku Co., Tokyo, Japan) and Fourier transform confocal laser Raman spectrometer (FT-Raman: LabRAM HR800, Horiba Jobin Yvon, Longjumeau, France). In TF-XRD, the measurement was conducted at a power of 50 kV and 200 mA. A CuK was used as X-ray source and the incident beam angle was set to 1◦ against the sample surface. In FT-Raman, measurement was conducted with 514.5 nm Ar laser at 16 mW of power excitation.

#### 2.2.3. Scratch Resistance

The scratch resistance of the surface layer to the metal substrate was examined by a thin-film scratch tester (CSR-2000, Rhesca Co., Ltd., Tokyo, Japan) according to JIS R-3255. A stylus with diameter of 5 μm and spring constant of 200 g/mm was pressed into the treated metal surface under the conditions of scratch speed of 10 μm/s, loading rates of 100 mN/min, and amplitude of 100 μm. Five measurements were performed on each sample, and their averaged values were used for analysis.

#### 2.2.4. X-ray Photoelectron Spectroscopy