**About the Editor**

**Andrzej Zieli ´nski** is a Full Professor at the Gdansk University of Technology, and former Head ´ of the Materials Science and Engineering Department and Division of Biomaterials. He is an author of about 300 scientific papers and 9 books/chapters, with close to 400 citations, and a Hirsch index of 14. His main interests include the deposition of ceramic, polymer, and composite biocoatings on titanium surfaces; oxidation of 3D laser-assisted printing of titanium; and degradation of materials. He is the supervisor of 18 Ph.D. students.

## *Editorial* **Special Issue: Recent Developments of Electrodeposition Coating**

#### **Andrzej Zieli ´nski**

Department of Biomaterials Technology, Institute of Machines Technology and Materials, Faculty of Mechanical Engineering and Shipbuilding, Gda ´nsk University of Technology, Narutowicza 11/12 str., 80-233 Gda ´nsk, Poland; andrzej.zielinski@pg.edu.pl

Academic Editor: Charafeddine Jama Received: 24 January 2021; Accepted: 26 January 2021; Published: 28 January 2021

Coatings are one of the forms of surface modifications of several parts produced in many branches of industry and daily life. Coatings may be applied for the protection of carbon steels, aluminum alloys, and even wood and concrete against environmental influence. The hard coatings decrease the wear resistance. However, even though painting is the most popular method of deposition of coatings on buildings, bridges, ships, etc., several functional coatings are being intensively developed. Such examples include the gas-barrier thin films protecting food [1], the coatings against excessive wear of different parts [2,3], textile coatings, e.g., by nanosilver [4], super-hydrophobic coatings [5], and ceramic–metal coatings for protection against erosion [6]. The coatings are often applied in medicine to make the healing time of load-bearing implants with bone faster, to enhance the antibacterial properties, to make steel implants bioinert and easy to remove, and many others. The examples are numerous [7–12].

This Special Issue is aimed at reviewing the newest achievements, particularly in biocoatings. They may be obtained through many techniques, such as direct electrocathodic deposition, pulse electrocathodic deposition, electrophoretic deposition, micro-arc oxidation, chemical and plasma vapor deposition, magnetron sputtering, pulsed laser deposition, electropolymerization, and the sol–gel method, further described in [13–22]. The biocoatings may be made of ceramics, polymeric, metals, or be composite coatings, all so far proposed in the literature.

The electrodeposition coating is among the most plausible techniques, because it makes it possible to design and obtain coatings with different microstructure, thickness, adhesion, and mechanical, physical, chemical, and biological properties. In this Special Issue, composed of nine papers, such examples are shown. One of them shows the deposition of metallic coatings, four papers consider oxidation processes of titanium alloy, three papers are devoted to composite coatings, and the last is review paper on the materials and methods.

Vainoris et al. [23] focused on metallic copper coatings deposited on a flat surface and 3D foams of Cu substrate. The copper deposition occurred much faster on copper foams than on a flat surface, making the metal foams highly suitable for electrowinning. The mechanism of copper deposition was determined, and the capacities of the double electric layer (DL) were calculated. In particular, the DL capacity was much higher and the charge transfer resistance slightly lower for the Cu foam electrodes. As a consequence of this research, the metal foam electrodes were recommended for use in several electrochemical processes.

Ossowska and Zieli ´nski [24] investigated the behavior of new and already used dental implants and the role of oxide layers. In particular, the possible mechanisms of oxide degradation and its influence on titanium corrosion at inflammation states were considered. The extremely low dissolution of rutile, slightly increasing along with pH, was measured. The diffusion of titanium ions through the oxide layer was shown as negligible. The single important mechanism of corrosion was demonstrated as initiated by the oxide layer damage at the defects caused by either the manufacturing process or implantation surgery. Therefore, a stepwise appearance and development of cracks through the oxide layers could be observed and enhance titanium corrosion.

Ossowska et al. [25] focused on the development of sandwich oxide coatings on a titanium base. Two-stage oxidation resulted in the inner solid layer and the outer nanotubular layer of oxides. Such structure of the coating significantly improved mechanical (hardness) and chemical (corrosion resistance) properties. This new technique may be used to substantially improve the surface of titanium load-bearing implants.

Ja˙zd˙zewska and Bartma ´nski [26] aimed at increasing the corrosion resistance and improving the biocompatibility by oxidation of a model screw dental implant made of the Ti–13Nb–13Zr alloy. The obtained nanotubular layers were of thickness 30–80 nm. The important difference in roughness was noticed between the top of the helix and its bottom. Uneven oxidation of screw model implants resulted in higher corrosion current and less noble corrosion, also known as pitting.

Dziaduszewska et al. [27] studied the micro-arc oxidation in some Ca- and P-containing electrolytes of the selective laser-melted Ti–13Nb–13Zr alloy to obtain ceramic–ceramic composite coatings. The study showed the voltage as the most significant process parameter influencing the coating characteristic. They obtained the coatings with a high Ca:P ratio, hydrophilicity, early-stage bioactivity, Young's modulus, and hardness close to those of bone, and appropriate adhesion of the coating to the titanium surface preventing delamination. Such coatings are especially suitable for dental implants.

Majkowska et al. [28] investigated deposition by the electrophoretic method of ceramic–ceramic coatings composed of hydroxyapatite and carbon nanotubes achieved as bilayers (subsequent deposition) and hybrid coatings (simultaneous deposition). It was shown that the pure multi-wall carbon nanotubes (CNTs) layer showed the best mechanical and biological properties. Both bilayers and hybrid coatings demonstrated insufficient properties attributed to the presence of soft, porous hydroxyapatite and the agglomeration of CNTs.

Pawłowski et al. [29] studied the ceramic–polymer coatings obtained by electrophoretic deposition and composed of chitosan and Eudragit compounds. The best process parameters were estimated. The Young's modulus of coatings was close to that of human cortical bone. The doping of Eudragit significantly reduced the degradation of coatings in artificial saliva at neutral pH, while maintaining high sensitivity to pH changes. The composite coatings showed a slightly lower corrosion resistance compared to the chitosan coating, and comparable hydrophilicity.

Zhang et al. [30] studied metallic (Fe)–ceramic coatings obtained by micro-arc oxidation on Mg alloys. The deposition of such coatings substantially increased the degradation resistance and in vitro cytocompatibility. The developed coatings exhibited potential in clinical applications.

In their paper, Zieli ´nski and Bartma ´nski [31] reviewed the state of the art in electrodeposition coatings. The developments of metallic, ceramic, polymer, and composite electrodeposited coatings were investigated. The direct cathodic electrodeposition, pulse cathodic deposition, electrophoretic deposition, micro-arc oxidation in electrolytes rich in P and Ca ions, electro-spark, and electro-discharge methods were characterized. The most popular were the direct and pulse cathodic electrodeposition, and electrophoretic deposition. The justification of the development of different coatings was an expected increase in bioactivity, mechanical strength, adhesion of coatings, and antibacterial properties.

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

#### **References**


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## *Article* **Metal Foam Electrode as a Cathode for Copper Electrowinning**

#### **Modestas Vainoris 1, Henrikas Cesiulis 2,\* and Natalia Tsyntsaru <sup>1</sup>**


Received: 31 July 2020; Accepted: 22 August 2020; Published: 25 August 2020

**Abstract:** The geometry of porous materials is complex, and the determination of the true surface area is important because it affects current density, how certain reactions will progress, their rates, etc. In this work, we have investigated the dependence of the electrochemical deposition of copper coatings on the geometry of the copper substrate (flat plates or 3D foams). Chronoamperometric measurements show that copper deposition occurs 3 times faster on copper foams than on a flat electrode with the same geometric area in the same potential range, making metal foams great electrodes for electrowinning. Using electrochemical impedance spectroscopy (EIS), the mechanism of copper deposition was determined at various concentrations and potentials, and the capacities of the double electric layer (DL) for both types of electrodes were calculated. The DL capacity on the foam electrodes is up to 14 times higher than that on the plates. From EIS data, it was determined that the charge transfer resistance on the Cu foam electrode is 1.5–1.7 times lower than that on the Cu plate electrode. Therefore, metal foam electrodes are great candidates to be used for processes that are controlled by activation polarization or by the adsorption of intermediate compounds (heterogeneous catalysis) and processes occurring on the entire surface of the electrode.

**Keywords:** metal foam; surface area; electrowinning; Cu electrodeposition; EIS; double electric layer capacitance

#### **1. Introduction**

The ever-increasing need for electronics, especially, handheld and portable electronics, and the need to reduce their size and increase their efficiency, generates a lot of various electronics waste all over the globe [1–3]. There are many ways to reclaim used metals in electronic waste; however, electrowinning is a very efficient and quite selective process allowing the recovery of high amounts of various pure metals [4–6]. Metallic foams and porous electrodes have an outstanding potential to be used as a cathode to collect deposited metals because of the functionality of their combined material properties resulting from their specific morphology. There already is great interest in the synthesis of various porous materials such as metal foams, nanowires, porous coatings, thin porous films, etc. [7–15]. Depending on the materials, type of pores (open or closed cells), the porosity and size of pores, such materials have broad application capabilities, from simple ones such as heat transfer or electrodes to more complicated cases of various redox reactions, catalysis, sensing, supercapacitors, or even gas storage because of the high surface area and low density available [9,16–29].

Any solid metal surface that acts as a substrate for electrochemical reactions possesses a certain roughness that can affect in different ways the values of the limiting diffusion current and the exchange current density. On the other hand, if the surface coarseness is relatively small, the limiting diffusion current density does not depend on the surface roughness, and it can be only correlated to the apparent surface of the electrodes. If the surface roughness of electrodes increases, the effective values of the exchange current density are also increased for the process under consideration, which is standardized to the apparent electrode surface area. At the same time, the limiting diffusion current density depends on the surface coarseness due to the decrease of the effective value of the diffusion layer thickness. If the level of the electrode surface coarseness remains low, the change of the limiting diffusion current density can be neglected [30]. In addition, it has been shown that when the metal deposition is controlled by diffusion (particularly silver), the surface with the highest surface roughness had a lower number of active sites but higher deposition efficiency and a higher efficiency of charge transfer [31]. The dependence between surface roughness and deposition efficiency is non-linear; the surface roughness needs to be quite high to affect the deposition efficiency [31,32]. It was proven that when the deposition reaction is controlled by the diffusion, the geometry of the electrode has no significant influence on the reaction [33]. Using very porous or surfaces with high roughness, one can eliminate activation and diffusion overpotentials, making the reaction process controlled by Ohmic effects and thus making the reaction much faster [30,31]. All these effects make porous metal electrodes with pore diameters higher than 50 μm high-performance cathodes for deposition reactions under diffusion control.

The estimation of the active surface area of highly porous conducting materials is also very important. Thus, various in situ or ex situ techniques can be used for these purposes. In situ techniques are preferred, since drying the sample can cause changes in the surface area and/or oxidation of the surface, changing its characteristics. Depending on the material and its porosity, one can use techniques for double electric layer estimations (cyclic voltammetry, initial charge-up dependencies, electrochemical impedance spectroscopy (EIS)), or adapt various adsorption/redox reactions that occur on the surface (underpotential depositions, adsorption measurements, reduction of various dyes, etc.) [29,34–39]. The classical techniques for surface area estimation—liquid permeability, gas adsorption (Brunauer–Emmett–Teller technique)—in some cases can also be used [34,35,39]. However, these techniques require higher amounts of materials and can have quite large error margins, depending on the geometry of the pores and the sample itself. For porous materials that are quite level, and with ordered pores, more sophisticated techniques could be used for porosity estimation such as atomic force microscopy (AFM) or spectroscopic ellipsometry; the latter requires a rather complex modeling [40–42].

EIS is a very powerful and versatile in situ technique that allows not only estimating the true surface area of conducting materials but also investigating the surface and the processes happening at the surface [9,20,23,27,29]. Using the EIS technique, one can investigate both Faradaic and non-Faradaic processes on the surface simultaneously [27,29,34,35,43–47]. Even the size and distribution of the pores can be characterized by employing the EIS technique [45,46]. However, the surface area determined by EIS or any other electrochemical method is not the true surface area, but rather the electrochemically active surface area, which can be much more useful when trying to determine the activity of porous materials for a hydrogen evolution reaction (HER) or other electrochemical reaction [39–47].

In this work, we investigated the deposition of copper on the plate (2D) and foam (3D) copper substrates using voltammetry and EIS. The comparison of 2D and 3D electrodes has been carried out to determine differences in double electric layer formation, charge transfer, diffusion, and deposition rates. These results are important for trying to enhance the potential application of foam electrodes in industry, and particularly for the electrowinning of copper from electronics waste.

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

#### *2.1. Materials and Sample Preparation*

All of the chemicals used for analysis were of analytical grade (Carl Roth, Karlsruhe, Germany). Solutions have been prepared using deionized water (DI). Solution compositions used for electrochemical experiments are shown in Table 1. The pH of solutions was adjusted using sulfuric acid and controlled by a benchtop pH-meter ProLine Plus (Prosence B.V., Oosterhout, The Netherlands). Cu plates and Cu foam electrodes served as working electrodes. The Cu foam sheets used to fabricate electrodes were purchased from Alfa Aesar. To characterize commercially available copper foams, we have done some experiments trying to determine the basic characteristics of this foam. Foam density has been determined as gravimetrically being equal to 0.748 g/cm3, making the porosity of the foam to be around 90.5%. The copper foam has a 3D interconnected porous structure, which can be observed in SEM images (Figure 1). The pore size varies from 1 to 0.1 mm. The surface of the foam is very uneven, making the true surface area of the already porous copper foam even larger.

**Figure 1.** SEM images at low (**a**) and high (**b**) magnification of 3D copper foam.

Working electrodes (copper plates and copper foams) have been washed and degreased using acetone, ethanol, and water in succession and in combination with ultrasonic bath. Both flat and porous samples were 1 cm × 1 cm in geometrical size with both sides conducting. To ensure that the working surface was that of the desired size, other parts of the samples were isolated using insulating plastic spray (PRF 202, Taerosol Oy, Kangasala, Finland). Just before measurements, the native copper oxide layer has been removed by dipping copper samples into 2 M H2SO4 solution for 2 s and afterward rinsing with DI water.

**Table 1.** Composition of solutions used for electrochemical measurements.


#### *2.2. Instrumentation and Methodology*

Morphology: The morphology of copper foams has been investigated using a scanning electron microscope (SEM, Hitachi's Tabletop Microscope TM-3000, Tokyo, Japan).

Electrochemical Measurements: Electrochemical measurements (voltammetry, EIS, chronoamperometry, etc.) have been performed using programmable potentiostat/galvanostat AUTOLAB PGSTAT 302N (Metrohm, Utrecht, The Netherlands). The software used for controlling the hardware was Nova 1.11.2.

Conditions of Electrochemical Measurements: A three-electrode system was used for all the electrochemical experiments, where Cu plates or Cu foams were used as working electrodes, circular platinized titanium mesh (Alfa Aesar, Ward Hill, MA, USA) was used as a counter electrode, and Ag/AgCl filled with saturated KCl solution (Sigma-Aldrich, St. Louis, MO, USA) was used as a reference electrode. The distance between the counter and working electrode was fixed at 2.5 cm. All electrochemical experiments have been performed at room temperature. Voltammetry measurements were done using the potential sweep voltammetry technique on Cu plates and Cu foams as working electrodes, starting at open circuit potential and going up to −1.2 V versus Ag/AgCl at a 2 mV/s scan rate. Voltammetry measurements have been performed using all the solutions shown in Table 1. Chronoamperometry experiments were performed at 4 distinct potentials (−0.1, −0.2, −0.4 and −0.6 V versus Ag/AgCl) using different substrates as working electrodes (Cu plates or foams) in 0.1 M CuSO4 and 0.4 M Na2SO4 solution. The same amount of electric charge was used to deposit coatings, i.e., 30 C. The current efficiency was calculated using chronoamperometry data and change in substrate mass after deposition.

Electrochemical Impedance Spectroscopy (EIS): Electrochemical impedance spectroscopy (EIS) measurements have been done using a standard three-electrode system, carried out in a frequency range of 10 kHz to 0.1 Hz, using perturbation amplitude of 10 mV. Obtained data were fitted to the equivalent electric circuit model (EEC) using ZView 2.8d software.

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

#### *3.1. Copper Foam Characterization*

In order to determine how the behavior of copper foams differs from flat surfaces in solutions, voltammetry experiments with different copper sulfate concentrations were carried out; the compositions of the solutions are shown in Table 1. The concentration of the sulfate anion was kept at 0.5 M to maintain the same buffering power in all of the solutions. The obtained polarization curves for the plate and foam electrode are shown in Figure 2, where the ordinate axis is displayed in a logarithmic scale because of a big difference in the current values between tested concentrations. To estimate the influence of porosity on the copper deposition, the geometrical sample size was the same for both Cu plates and Cu foams (1 cm × 1 cm). As can be seen from Figure 2, Cu deposition starts somewhere around −0.075 V versus Ag/AgCl and did not depend on the substrate used. After the peak representing the Cu2<sup>+</sup> reduction to Cu0, the current on both surfaces and all the concentrations turns into an almost constant one. The reason for this could be the mass transport limitations because the leveling off of the current depends on the concentration of Cu(II) in the solution. This is also supported by the slight increase of the current with the rise of polarization at higher concentrations (50 mM to 0.2 M), showing that with higher potential, the positive ions are attracted from further away, and the deposition rate increases.

In addition, voltammetry tests also showed that independently of the substrate used, the hydrogen evolution reaction (HER) started in the range of −1.0 to −1.1 V versus Ag/AgCl in the solutions containing 10 and 50 mM of CuSO4. This fact could be attributed to the governing role of pH change in the pre-electrode layer during electrodeposition, and this change seems to be similar for both solutions. However, in the solution containing 0.2 M CuSO4, the HER started around −0.75 V versus Ag/AgCl on both surfaces. It can be linked to the higher rate of copper electrodeposition, and in turn, the pH decrease near the working electrode. Thus, the major difference between the two surfaces can be noted from voltammetry experiments: there was an approximately 3 times higher current on the foam substrate at all potentials in comparison to the flat surface. This difference can be explained by the better hydrodynamic conditions of copper foams substrate: the porous surface allows for faster mass transport and exchange.

For further investigation, the solution containing a similar amount of Cu(II) as in solutions used for the metals recovery from the electronic waste was chosen. Regarding the influence of the surface type on the Cu electrochemical deposition, chronoamperometric measurements have been done in 0.1 M CuSO4 and 0.4 M Na2SO4 solution at four fixed potentials: −0.1, −0.2, −0.4, and −0.6 V versus, Ag/AgCl, and at a fixed amount of charge passed through the cell (30 C). The results have been summarized and are shown in Table 2.

**Figure 2.** Cathodic voltammograms on Cu plate (**a**) and foam (**b**) obtained in the electrolytes with various concentrations of CuSO4 (the compositions of solutions are shown in Table 1), potential scan rate 2 mV/s.

**Table 2.** Cu deposition rates on 2D and 3D electrodes in the solution containing 0.1 M CuSO4 and 0.4 M Na2SO4.


Chronoamperometric measurements (Table 2) clearly show an approximately 3 times faster copper deposition rate on the foam at all tested potentials. In this case, there was no hydrogen evolution, and the deposition efficiency was almost 100% on both substrates. A considerably higher deposition rate on the cooper foam substrate supports the idea that the deposition is controlled by diffusion to the electrode having a higher specific surface area. In addition, a higher metal deposition rate on the foam electrodes makes them an attractive substrate for the electrowinning of metals compared to other

materials having a similar geometric area. The morphology of deposits is influenced by the potential and type of substrate, as it is shown in the SEM images in Figure 3.

**Figure 3.** SEM images of potentiostatically electrodeposited Cu coatings at different cathodic potentials on flat and foam copper substrates after 30 C passed charge. The bath was 0.1 M CuSO4 and 0.4 M Na2SO4.

The copper deposits have globules shapes on the flat electrodes, and the morphology did not differ at these two potentials. This is related to the very similar electrochemical deposition rates at these potentials, and as it can be seen from the voltammetry data (Figure 2) and efficiency of deposition, there were no side reactions, and the current was similar at these two potentials. Another case is the deposition on the porous substrate. At −0.2 V versus Ag/AgCl, copper forms cauliflower-like crystalline agglomerates with well-defined edges. At higher potential, the copper forms smoother surfaces that are still cauliflower-like structures. The coverage of both surface geometries was good even without external agitation, even at low potentials.

#### *3.2. Surface Area and Di*ff*usion Rate Estimations*

To characterize copper foams and estimate the active surface areas for the charge and mass transfer processes that occur during the electrochemical deposition of copper, we utilized the EIS technique. EIS measurements have been done for all the solutions listed in Table 1. EIS measurements were performed at cathodic potentials of −0.125, −0.15, −0.175, and −0.2 V versus Ag/AgCl on flat and porous copper substrates. These potentials were chosen based on chronoamperometric data. At such low potentials, the change of surface morphology during deposition is still minimal and can be ignored in this case. Typical EIS scans on the copper plate at various potentials are shown in Figure 4. From the EIS data plots, we can see that at investigated potentials, the data plot can be divided into two zones: the high-frequency semicircle and the low-frequency (starting around 75–100 Hz) 45◦ angle line. The high-frequency semicircle can be attributed to charge up of the double layer and charge transfer to the copper ions, whilst the low-frequency line is attributed to the formation of the concentration gradient of the copper ions. To better evaluate ongoing processes, EIS data were fitted to the equivalent electric

circuit (EEC) that is shown as an inset in Figure 4 of the Nyquist plot (a). The elements of applied EEC have the following physical meaning: R0 is resistance at the electrode/electrolyte interface, CPE(DL) is a double-layer capacitance modeled via the constant phase element (CPE), R(CT) is a charge transfer resistance, CPE(W) stands for the capacitance caused by the concentration gradient, and R(Diff) is a resistance caused by the concentration gradient. The element CPE(W) is attributed to the diffusion because of the signature 45◦ angle seen in the Nyquist plots at low frequencies (Figure 4), and the value *n* in this CPE element was very close to 0.5 in all the experiments. This constant phase element acting only in the low-frequency region represents diffusion, and it can be used as a Warburg element when *n* = 0.5 [48,49]. The values of the constant phase element CPE(DL) have been recalculated into true capacitance using Hsu and Mansfeld's equation [50]. All values of components of the fitted EEC are indicated in Table 3.

**Figure 4.** *Cont*.

**Figure 4.** Nyquist (**a**,**d**) and Bode plots (**b**–**c**,**e**–**f**) on Cu plate (**a**–**c**) and foam (**d**–**f**) registered at various potentials (indicated on graphs) in 0.1 M CuSO4 + 0.4 M Na2SO4 solution at 20 ◦C. Points—experimental data, solid lines—results of fitting to equivalent electric circuit (EEC) shown in the inset (**a**).

As it is seen, the proposed EEC describes well experimental EIS data on both substrates in a whole investigated potential range. The values of the capacitance of the double electric layer on both substrates might be used to estimate differences in real areas between the plate and foam electrodes, i.e., to estimate the roughness factor as a ratio of C(DL) on foam and plates that have the same geometric area (1 cm × 1 cm). Notably, the double-layer capacitance (C(DL)) extracted from the EIS data is 50 μF (see Figure 5), and it is in good agreement with the theoretical values assigned to 1 cm2 of copper [49]. The capacitance of the double layer of a commercial foam, that has the same geometric area as a plate, is 7 to 14 times higher in comparison with a plate electrode. The thickness of the double electric layer is very small and is in tens of nanometers; therefore, this layer replicates the surface morphology on the nano-level, and the ratio with the value obtained on the plate electrode can represent the roughness

factor, and it matches the ratio of C(DL) of both surfaces –(C(DLfoam); C(DL plate) is 7–14:1). However, the increase of double-layer capacitances with the increase of applied cathodic potential on both flat and porous surfaces is different. On the porous electrode, the C(DL) increase is much higher when compared to the change in capacitances of the flat electrode. This increase is related to the much higher surface area, and the distribution of current on the surface of the foam. With higher potential, the current distributes more evenly on the whole foam surface, and the edge effect is less apparent, which also influences the surface area estimations [51,52].

When looking at the effect that the concentration of copper ions has on the EIS parameters (Table 3), we can divide the results into three sections: high concentration (0.2 M), mid-level concentrations (0.1 and 0.05 M), and low concentrations (0.01 M). The double electric layer (DL) capacitance values do not differ that much with the change of the concentration on both surface geometries. However, when looking at charge transfer resistance, the differences between concentrations are significant. At low concentrations, the charge transfer resistance is very high; this is caused by the lack of copper ions. In contrast, this resistance at mid-level concentrations is around 6–9 Ω, which depends on the surface geometry as well as applied potential (Figure 5). At high concentrations (0.2 M and higher), the charge transfer resistance values decrease approximately 3 times on both surfaces, because of an abundance of conducting particles. Nevertheless, this charge transfer resistance is lower at all investigated potentials and all concentrations on the foam electrode, showing that the reduction reaction occurs faster on the copper foams.

When taking a look at the charge transfer resistance dependence on potential (Figure 5) with both types of electrodes, it is clear that the 3D electrode displays approximately 1.5–1.7 times lower charge transfer resistance than the 2D electrode, agreeing with the results of voltammetry (see Figure 2). The differences in the charge transfer resistance on plate and foam electrodes are lower than the differences in the capacitances of DL, because the reaction layer is thicker than the DL, and in some areas of the foam electrode, it overlaps. As it can be seen from Figure 5, the difference between 2D and 3D electrodes in charge transfer resistance is higher at low potentials; thus, the charge transfer reaction on the foam occurs easier, and it partially explains the higher Cu deposition rate (see Table 2). However, lowering the charge transfer resistance, or in turn, the increase of the rate of the charge transfer reaction by approximately 2 times, does not result in increases in the Cu deposition rate by approximately 3 times.

**Figure 5.** Dependence of double-layer capacitance (ordinate at the right) and charge transfer resistance (ordinate at the left) on potential applied for Cu plate and foam electrodes in 0.1 M CuSO4 + 0.4 M Na2SO4 solution.

To further characterize the difference in copper deposition reactions on flat and porous copper surfaces, the components of EEC related to diffusion have been investigated in detail (Figure 6). The foam has lower charge transfer resistance, meaning faster reactions and better hydrodynamic qualities, allowing for faster diffusion and in turn the much faster deposition, even with a larger surface and in turn, lower current density.

**Figure 6.** Dependence of diffusion-related elements of EEC on the potential applied. Measurements performed using a copper plate and copper foam as working electrodes in 0.1 M CuSO4 + 0.4 M Na2SO4 solution.

**Table 3.** Values of electrochemical impedance spectroscopy (EIS) parameters obtained by fitting data obtained on copper foam and copper plates at −0.175 V versus Ag/AgCl at different copper concentrations. EC used for modeling shown in Figure 3 inset. CPE(DL): a double-layer (DL) capacitance modeled via the constant phase element (CPE), CPE(W): the capacitance caused by the concentration gradient, R(CT): charge transfer resistance, R(Diff): resistance caused by the concentration gradient.


The parameter related to diffusion CPE(W) at low concentrations is almost equal on both surface geometries, showing that the diffusion effect is similar, but the resistance at low concentration is about 2.5 times higher. It means that the diffusion layer is much thicker on the copper foams surface because of the porosity effect. Therefore, it causes a higher rate of copper electrodeposition. The overall trend in mid-level and high concentrations is that with the increase of Cu2<sup>+</sup> concentration, the CPE(W) value increases, and the R(Diff) decreases. As it is seen from Table 3, the difference between R(Diff) values at 0.2 and 0.05 M concentrations on the flat surface is only around 14%, whereas on the foam electrode, the values of R(Diff) are lower, but all values are sensitive to the concentration of Cu(II) in the solution. The highest value of R(Diff) is obtained on the foam electrode at a relatively low concentration of Cu(II), i.e., 0.01 M, which is probably due to the faster depletion of copper ion concentration in the 3D diffusion layer and the necessity of a longer time to supply Cu(II) ions into the pores. Since the deposition rate on the foam electrode at a higher concentration of Cu(II) is 3 times faster than on the

flat electrode, this is mirrored by the behavior of CPE(W), showing that the diffusion occurs 3 times faster on the foam. The efficiency of charge transfer on the porous surfaces is higher as well, which is in good agreement with other studies of metal depositions on porous surfaces [31].

To even better understand the diffusion peculiarities on 2D and 3D electrodes, the diffusion impedance using extracted values from total impedance data (presented in Table 4) was calculated. As it is shown in Figure 4, the copper deposition occurs under diffusion control at low frequencies (below 100 Hz) on both foam and plate electrodes, and diffusion is modeled by a parallel connection of CPE(W) and R(Diff) elements (see Figure 4). In this case, diffusion impedance, *Zdi*ff, as a function of frequency is calculated by the equation:

$$Z\_{diff}(\omega) = \frac{R\_{Diff}}{1 + (j\omega)^a Q R\_{Diff}} \tag{1}$$

where *Q* and α are parameters of CPE(W), R(Diff) is resistance caused by diffusion, and ω is the phase angle (ω = 2π*f*). However, when α = 1 − *Q* is pure capacitance, in our case, α = 0.5, and the CPE represents diffusion [53].

The calculated diffusion impedance data are presented in Figure 7. As it is seen, the diffusion impedance on the plate Cu electrode is 2–4 times higher than that on the foam Cu electrodes, which is dependent on the frequency and potential applied.

**Figure 7.** Bode plots of extracted diffusion impedance at various potentials on flat Cu substrate (**a**); and Cu foam substrate (**b**).

These results once again confirm the chronopotentiometric data obtained on both 2D and 3D Cu electrodes. For chronopotentiometry experiments, current values have been chosen higher than the limiting current values seen in Figure 8**.** In this case, the transition time at which the concentration of metal ions on the electrode becomes equal to zero is visual on the chronopotentiograms, and the effective diffusion coefficient can be calculated by the Sand equation:

$$i\sqrt{\tau} = \frac{nFAC\_0\sqrt{\pi D\_{eff}}}{2} \tag{2}$$

where τ is a transition time (s), *i* is a current (A), *C*<sup>0</sup> is the concentration of Cu(II) ions (mol/cm3), *D*eff is the effective diffusion coefficient (cm2·s−1), *F* is Faraday's constant, *n* is the number of electrons participating in the electrochemical reaction; and *A* is a geometrical surface area.

**Figure 8.** Chronopotentiograms on flat (continuous lines) and porous (dashed lines) electrodes at various current densities in 50 mM CuSO4 and 0.45 M Na2SO4 solution. All the densities have been calculated for the geometrical area of the substrate of 1 cm2.

In our case *i* √ τ ∼ *const*, so the maximal deposition rate is controlled by the mass transfer. The values of the effective diffusion coefficient of Cu2<sup>+</sup> ions on both plate and foam Cu electrodes were calculated by Equation (2), and the data are shown in Table 4. The effective diffusion coefficient on the plate electrode is almost three times lower than on the foam electrode, and it is in good agreement with EIS data.


**Table 4.** Effects of electrode geometry on effective Cu(II) ions diffusion coefficient.

So, copper foams are great substrates for reactions that are either limited by the mass transfer (electrochemical depositions, etc.) or the ones that are restricted by adsorption or activation (HER and similar), making them great candidates to reduce the size of electrodes, but not to lose out on the efficiency and activity of electrodes.

#### **4. Conclusions**

A comprehensive investigation of the electrochemical deposition of copper onto 2D (plate) and 3D (foam) Cu substrates has been done. Using various electrochemical methods, it was determined that the rate-determining step in a copper deposition is diffusion. The main processes occurring on the electrode are the charge-up of double electric layer, charge transfer, and diffusion. The specific electrochemically active area of Cu foam was estimated from EIS data, and based on the values of the double electric layer, it was determined to be 7–14 times higher than that for the plate electrode. Based on the EIS data, it was determined that the charge transfer resistance on the Cu foam electrode is 1.5–1.7 times lower than that on the Cu plate electrode, which results in an increase in a charge transfer rate of approximately 2 times. Based on the analysis of the diffusion impedance and chronopotentiometry data, it was found that Cu2<sup>+</sup> mass transfer and the copper deposition rate is up to 3 times faster on the foam surface in comparison with a flat surface having the same geometric area in the same potential range. In addition, effective diffusion coefficients have been calculated from chronopotentiometry data using Sand's equation. These findings make Cu foam an attractive material for metal electrowinning processes as well as for processes controlled by adsorption (e.g., hydrogen evolution reaction).

**Author Contributions:** Investigation, M.V. and N.T.; methodology, M.V. and H.C.; supervision, H.C. and N.T.; visualization, M.V.; writing—original draft preparation, M.V.; writing—review and editing, M.V., H.C. and N.T. All authors have read and agreed to the published version of the manuscript.

**Funding:** This work was funded by the Lithuanian Business Support Agency (LVPA); project J05-LVPA-K-01-0022. **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* **The Mechanisms of Degradation of Titanium Dental Implants**

#### **Agnieszka Ossowska \* and Andrzej Zieli ´nski**

Department of Materials Engineering and Bonding, Gda ´nsk University of Technology, 80233 Gda ´nsk, Poland; azielins@pg.edu.pl

**\*** Correspondence: agnieszka.ossowska@pg.edu.pl; Tel.: +48-58-347-19-63

Received: 29 July 2020; Accepted: 26 August 2020; Published: 28 August 2020

**Abstract:** Titanium dental implants show very good properties, unfortunately there are still issues regarding material wear due to corrosion, implant loosening, as well as biological factors—allergic reactions and inflammation leading to rejection of the implanted material. In order to avoid performing reimplantation operations, changes in the chemical composition and/or modifications of the surface layer of the materials are used. This research is aimed at explaining the possible mechanisms of titanium dissolution and the role of oxide coating, and its damage, in the enhancement of the corrosion process. The studies of new and used implants were made by scanning electron microscopy and computer tomography. The long-term chemical dissolution of rutile was studied in Ringer's solution and artificial saliva at various pH levels and room temperature. Inductively coupled plasma mass spectrometry (ICP-MS) conjugated plasma ion spectrometry was used to determine the number of dissolved titanium ions in the solutions. The obtained results demonstrated the extremely low dissolution rate of rutile, slightly increasing along with pH. The diffusion calculations showed that the diffusion of titanium through the oxide layer at human body temperature is negligible. The obtained results indicate that the surface damage followed by titanium dissolution is initiated at the defects caused by either the manufacturing process or implantation surgery. At a low thickness of titanium oxide coating, there is a stepwise appearance and development of cracks that forms corrosion tunnels within the oxide coating.

**Keywords:** dental implants; corrosion; ringer's solution; artificial saliva; titanium oxide layers; inductively coupled plasma mass spectrometry (ICP-MS)

#### **1. Introduction**

Titanium alloys possess good strength properties and high resistance to the most aggressive environments such as hydrochloric acid or sulfuric acid [1–3]. The compact, stable oxide layer [4] is responsible for corrosion resistance and biotolerance, effectively stopping the anodic pickling of the substrate [5–7]. Another important function is to chemically stabilize the implant in the living organism [4,8–10]. The more compact and bonded the passive layer is to the substrate, the better the corrosion resistance. In the case of thick oxide layers, an improvement in tribological properties may also be observed [7]. According to Hanawa et al. [11], the top sublayer of the titanium oxide layer inhibits metal ion release [12,13] and its transformation in vitro. Additionally, the oxide layer promotes osseointegration and bone adhesion [14–16].

It is known that each implant inserted into the body is treated as a foreign body and can cause allergic reactions, inflammations, even the rejection of the implant. The human body is a very specific environment as the body fluids—extracellular fluids and blood—contain aqueous solutions of certain organic substances, dissolved oxygen, various inorganic anions (Cl<sup>−</sup>, HPO4 <sup>2</sup>−, HCO3−), and cations (Na+, K<sup>+</sup>, Ca2<sup>+</sup>, Mg2<sup>+</sup>), which together represent a highly aggressive environment [17]. The presence of amino acids and proteins accelerates the corrosion processes [18]. Besides, in the case of dental implants, the composition of the saliva is highly complex, containing both inorganic salts and organic components. This composition depends on many factors such as food, age, and diseases and the pH of saliva can vary around dental implants. The ingestion of acidic beverages can decrease the buccal pH, and the infections can also acidify the pH of saliva, contributing to the corrosion of dental implants. On the other hand, titanium and its alloys are sensitive to tribocorrosion [19]. The oxidation is performed mainly to prevent corrosion of titanium and its alloys in severe conditions. For dental implants, such oral environments include varying pH, acid attack and the presence of chemical compounds such as cetylpyridinium chloride, sodium fluoride and hydrogen peroxide [20].

The action of media containing fluoride ions causes degradation of the continuity of the oxide film followed by damage to the titanium as a result of the ingestion of fluoride ions into the oxide layer, thereby reducing its protective properties [21,22].

Corrosion processes influence changes in the structure of the implanted material, weakening its integrity, which can result in material discontinuities and cracks. Cells in direct contact with the exposed surface of the material are stimulated for the intensive secretion of inflammatory mediators, mainly neutrophils and macrophages [23]. In vitro studies [24] show that corrosion products are harmful to cell differentiation and proliferation processes.

The dissolution of the titanium oxide layer is due to the process of ion diffusion into the layer. In titanium, the oxygen atoms migrate via the interstitial diffusion mechanism, occupying the free, octahedral interstitial positions in the titanium hexagonal lattice. Studies conducted by Wu and Trinkle [25] showed that for oxygen atoms not only interstitial but also axial positions are available, i.e., all arrangements of oxygen atoms in the titanium matrix are possible.

The oxidation of titanium is faster when the material is subjected to high temperatures and the influence of an oxygenated environment. The overall oxidation reaction includes the formation of oxide followed by the diffusion of oxygen into the bulk of the titanium. Oxygen diffusion creates an oxygen-enriched layer due to the high solubility of oxygen in the titanium and the oxygen stabilizing effect in the crystalline titanium structure [26]. In solids, the most likely atomic diffusion mechanism is a vacancy or interstitial mechanism, i.e., the motion of atoms occurs as the consequence of the presence of imperfections [27]. The interstitial diffusion mechanism is typical for low atomic radius atoms such as hydrogen, oxygen, carbon, and nitrogen.

The great advantage of the oxide layers produced on titanium and its alloys is their capability of repassivation, and some released ions depend on regeneration. Hanawa et al. [11], while measuring repassivation potentials, estimated the recovery rate of the oxide layer in 0.9% physiological saline: for 316L steel as 35.3 min, for Ti6Al4V as 8.2 min, and Co28Cr6Mo as 12.7 min. The research conducted by Hanawa et al. [28] showed that in Hanks' solution, the rate of repassivation was lower than in 0.9% saline solution.

Metallic elements have a different tendency to release ions, and even trace amounts of elements in the alloy composition should not be neglected [11]. There are data on the significant contents of some alloying elements of Ti6Al4V within the tissue around the implanted alloy. So far [29] reports on the consequences of ion release into the body have focused on the importance of the possible impact of released ions on biomolecules and the initiation of adverse biological reactions as the titanium ions could quickly react with water molecules or inorganic anions, easily binding with body fluids.

Osseointegration involves a series of biological events influenced by multiple factors. Among them, the porous-structured Ti alloys have shown to allow rapid bone ingrowth and improved osseointegration by increasing the bone-implant interface area [30,31]. Such conditions are achieved for dental implants by micro-arc oxidation, which brings out the rough surface [19,32–35]. The bioactivitity is usually increased by anodic oxidation in an electrolyte containing calcium phosphates [36,37], and wear resistance by incorporation of tough nanoparticles [35].

We have put a hypothesis that the damage of oxide coating can be sometimes or often, the main cause for the degradation of material and removal of the implant. The purpose of the study was to characterize the processes which allow for titanium dissolution from dental implants. To achieve that, surface examinations of new and applied dental implants were carried out. The dissolution rates of titanium dioxide (rutile) into two simulating body fluids at different pH values were performed. The titanium transport through the oxide coating was also calculated.

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

#### *2.1. Microstructural Characterization of Surfaces of Implants*

The first stage of research was the qualitative analysis of the surface as well as of the cross-sections for new and used (removed) dental implants. The tests were carried out on groups of samples:


The examinations of the surfaces and cross-sections of new and used dental implants were carried out at the Gdansk University of Technology using a scanning electron microscope (SEM; JEOL JSM-7600F, JEOL Ltd., Tokyo, Japan). Before observation the new and used surfaces of dental implants were cleaned in methanol. The cross-section samples were cut from the implants and were ground with abrasive papers (No. 2500 as the last, Struers Inc., Cleveland, OH, USA).

To obtain detailed information on the geometry and presence of cracks or delaminations of the coatings on the dental implants, a computer microtomography (CT) technique was used. The CT examinations were made with the μCT (General Electric, Lewistown, PA, USA) phoenix v-tome-x s using an X-ray "direct tube" with a set power of 17 W (70 kV, 100 μA). One thousand radiographs (2D X-rays) for each tomogram were made with 360◦ rotation and an exposure time of 333 ms (for a single radiogram). 3D tomograms were reconstructed from radiographs using the phoenix datos-x2 reconstruction program and a standard reconstruction algorithm. The reconstructed samples had a resolution of 2.413 μm/Voxel and were analyzed using the commercial VGStudio Max package.

#### *2.2. Investigation of Dissolution Rate of Oxide Coatings*

In the second stage, the tests of the dissolution rate of rutile (titanium oxide) were performed in two simulated body fluids (SBF). The starting material was the titanium oxide powder (purity of 98.0%–100%) delivered by Acros Organics (Morris Plains, NJ, USA). Cylindrical samples, of dimensions 6 mm × 3 mm (diameter × length), were prepared using the classical powder metallurgy method without a filler. The material was formed in a single-axis pressing process using a force of 2 kN acting on the stamp for 60 s at the position shown in Figure 1.

**Figure 1.** The cold pressing scheme for manufacturing the titanium dioxide specimens by powder metallurgy.

The sintering process was carried out in a chamber oven (Type 22 MRT/1300, Conbest Ltd., Kraków, Poland) at 1300 ◦C for 2 h in an air atmosphere. The samples were heated at a rate 0.5 ◦C/min up to 100 ◦C and then at a rate 3 ◦C/min. Such two-stage heating significantly limited the appearance of cracks and delaminations were not observed.

Two simulated body fluids were used. Ringer's solution was prepared based on commercial tablets (Ringer's tablet, Merck, Germany) and artificial saliva according to the composition shown in Table 1. Hydrochloric acid (4M) (HCl) was used to prepare solutions of the appropriate pH values of 3, 5, and 7. The Elmetron CPI-505 pH meter was used to measure the pH values.


**Table 1.** Chemical composition of the artificial saliva.

The samples were cleaned in an ultrasonic washer and immersed in the prepared solutions for 3 and 12 months. After this time the test solution samples were analyzed for total titanium content at the Centre of Biological and Chemical Sciences, the University of Warsaw (Warsaw, Poland). The inductively coupled plasma mass spectrometry (ICP-MS, Perkin Elmer, Inc., Waltham, MA, USA) was carried out by the norm E2371-13 to determine the titanium content in the solutions. The ions were separated using a special mass analyzer, distributing the ions according to the value of their mass-to-charge ratio. The ICP-MS was calibrated using an external calibration curve, which was prepared using 1% nitric acid and a titanium pattern. The quadrupole mass spectrometer, Elan 9000 Perkin Elmer ICP-MS, with conjugated plasma induction excitation was used for the study.

The solutions were mineralized before measurements in a closed microwave system. Approximately 1 g of the solution and 1 mL of 30% nitric acid were used for mineralization. The samples were then diluted to 15 mL.

#### *2.3. Calculations of Di*ff*usion of Titanium Ions through Oxide Coatings*

The last stage was the calculations of the theoretical diffusion rate of titanium atoms through the rutile crystalline structure. The calculations were based on Arrhenius and Fick's laws [38] and the earlier high-temperature measurements [39,40].

#### **3. Results**

#### *3.1. Examinations of Implants*

Figure 2 presents the surfaces of different three new implants, produced by two companies, A and B. The layer discontinuities, material defects appearing in the layer, unevenness, numerous rolling scratches, and material allowances resulting from the surface treatment processes are visible before implantation.

The used implants were obtained from surgeons from the Warsaw Medical University. Only removed implants, among all, which demonstrated clear signs of damage, were selected. When examinations of the used implants (Figure 3) were made, two areas of surfaces could be distinguished. The first was the top of the threads, with characteristic flattened, rubbed bumps, with clearly visible pits on the surface. The second type was the bottom of the threads, in which there is a detachment of the material of the layer from the ground, and numerous deep cracks are visible likely arisen as a result of stress concentration, which are potential places of corrosion

initiation and development of corrosion processes in the environment of body fluids and particularly aggressive saliva.

**Figure 2.** Surfaces of new dental implants: (**a**) A company; (**b**) B company; (**c**) C company; (**d**) D company. SEM.

**Figure 3.** Surfaces of used dental implants: (**a**) A company; (**b**) B company; (**c**) C company; (**d**) D company. SEM.

The surfaces of the used implants were subjected to purification or sterilization processes, thanks to which we can observe traces of organic residues on the surface—bacteria and tissues. Surprisingly, the largest clusters of bacterial colonies are located between the tip and the bottom of the thread, on the lateral surfaces. Perhaps this phenomenon is caused by adverse conditions at the tops and bottoms of the threads. Pits are visible on the surface of the thread tops, numerous and deep cracks in the thread cavities, which may indicate a significant impact of the environment and continuous operation of the implant—the influence of tensile forces and friction forces that affect the implant placed in the bone. On the implant surfaces, discontinuities of bone formation (Figure 4a,c) and bacterial colony residues (Figure 4b,d) can be observed, which tightly cover the material.

**Figure 4.** Organic remains on the surfaces of used dental implants: (**a**) A company; (**b**) B company; (**c**) C company; (**d**) D company. SEM.

The implant cross-sections (Figure 5) illustrate surface unevenness, numerous discontinuities characterizing the layers produced, material stratification, a significant number of cracks of varying lengths, and arrangement occurring in the coating. They can be ideal for corrosion progress.

**Figure 5.** Cross-sections of used dental implants: (**a**) A company; (**b**) B company; (**c**) C company; (**d**) D company. SEM.

Using computed tomography, the layer thickness distribution on dental implants was depicted (Figure 6). The analysis shows that the thickness of the layers formed on the surface of the implants is diverse and does not evenly distribute. The surfaces of the vertices and thread cavities are characterized by a larger thickness of the coating. The images of the layer thickness distributions on the surface of the dental implants were very similar.

**Figure 6.** Images obtained by the computed tomography study of the sections: —longitudinal: (**a**) D company; (**b**) A company; —transverse dental implants: (**c**) D company; (**d**) A company.

#### *3.2. The Dissolution of Rutile*

Knowing the characteristics and defects occurring in the layers covering dental implants, in the second stage of research, the determination of the rate of penetration of titanium ions into the solution was undertaken. Titanium oxide powder samples were prepared using powder metallurgy processes, which were immersed in Ringer's solution and artificial saliva for a period of 3 and 12 months.

The surface of the samples produced, featuring a slight degree of porosity, is shown in Figure 7. The samples vary in grain size from 1.429 to 8.184 μm.

**Figure 7.** SEM microstructures of the sample surface after pressing with a force of 2 kN.

The analyses of the titanium ions carried out three months after their exposure to both SBFs at pH 3, 5, and 7 (Table 2), showed negligible solubility of the titanium dioxide, below the 0.080 mg/kg limit of determination. In solution adjusted to pH 3, titanium dissolution was distinctly higher, but only after 12 months, 0.093 mg/kg.



\* GO—Limit of identification of titanium 0.080 mg/kg.

Slightly different results were achieved for the artificial saliva solution. The contents of titanium ions for artificial saliva of different pH were higher than those for Ringer's solution and about 0.107 mg/kg for pH 5 and 7, while for pH 3 the titanium ion level exceeded 0.137 mg/kg.

These experiments show the importance of the pH value of body fluids and the possible dissolution and penetration of titanium ions into the human body. The increasing dissolutions of ferrous oxides [41] and copper oxides [42] with decreasing pH, and rutenium–titanium oxide coating at pH 2 [43] are in accordance with obtained results showing the important effect at the lowest pH value. The more distinct dissolution of rutile in artificial saliva is likely due to the higher chloride concentration in saliva than in blood, and the susceptibility of the oxides to the pitting. It is worth noting that the solutions used only simulated natural human body fluids and did not contain various biological substances such as enzymes, which can create an even more aggressive environment. The important conclusion, however, is that a titanium dental implant in the mouth is more susceptible to dissolution than in other tissues. The drastic lowering of the pH value of the solution accelerates the process of removal of titanium ions from rutile, always present on the titanium surface. Thus, changes in the pH value, occurring during some inflammatory reactions in the living organism, may significantly influence the condition of the oxide layer, and consequently the status of the implant.

In these long-term, expensive tests, we have analyzed the trends: the effect of decreasing pH and test solution on the dissolution rate. Taking into account the extremely low values of dissolubility, the precision of the spectrometric measurements, and the use of slightly porous materials, we have concluded that a significant number of specimens should be applied to obtain the homogenous sample and low standard deviations. Taking this into account, our purpose has been only to recognize at least the row of the magnitude of dissolubility and how the pH effects what has been reached. The results clearly show that the rutile ceramics dissolubility is extremely low and it decreases with decreasing pH, as it is for metallic substances.

#### *3.3. Di*ff*usion of Titanium Ions in The Oxide Layer*

To calculate the distance of the diffusion of titanium ions in rutile lattice, data of two references were taken into account. In [40], the random tetravalent titanium atoms were assumed to be the predominant defects evident from self-diffusion. The enthalpy of motion was determined as Δ*H*<sup>m</sup> = 57.03% ± 4.9% kcal/mole. In another report [39], for diffusing the radio-isotope titanium-44 into single crystal rutile at temperatures in the range of 900 to 1300 ◦C, the activation energy was found to be 61,400 calories per mole and the frequency factor was calculated to be 6.4 <sup>×</sup> <sup>10</sup>−<sup>4</sup> m2/s. Assuming the diffusion enthalpy at 59.2 kcal/mol and *D*<sup>0</sup> at the above value, the titanium diffusion coefficient at room temperature (293 K) was calculated at about 10−<sup>49</sup> m2/s. That following, the diffusion distance at this temperature in one year is about 3 <sup>×</sup> <sup>10</sup>−<sup>34</sup> nm.

#### **4. Discussion**

All new implants were made of the Ti6Al4V alloy by casting and milling (likely CNC). As a rule, such implants are assumed to have a perfect surface, at designed roughness achieved by mechanical treatment or chemical acidic (SLA implants) or alkaline treatment. Some of the commercial implants have deposited coatings (Osseotite and Nanotite implants). The detailed surface treatment is not disclosed. The majority of dental implants are likely subjected to micro-arc oxidation in phosphate solutions.

So far, the imperfections visible on new implants are attributed to the forces acting during implantation surgery. For example, the grooves and abraded facets, and loose titanium particles at the interface were reported for dental implants and attributed to the surgical procedure [44].

It is a damage that certainly locally destroys the titanium coatings. However, our investigations showed that several implants before any implantation possessed already imperfect surfaces with such damage forms of the oxide coating as the layer discontinuities (holes/pits), large unevenness, and rolling scratches. Such defects may initiate the local degradation of an implant, in particular the cracking and pitting corrosion.

The examinations of implants used for a relatively short time confirmed the above assumption, even indirectly. The pits are visible on the tops of the threads, and the detachment of the coating on the bottoms. The numerous deep cracks are likely arisen as a result of stress concentration, and can serve as potential places of corrosion initiation and development in the environment of body fluids and particularly aggressive saliva.

The traces of organic residues on the surface were between the tops and the bottoms of the threads. Perhaps this phenomenon is caused by adverse conditions at the tops and bottoms of the threads due to a significant impact of the environment and continuous influence of tensile and friction forces that affect the implant placed in the bone. The detailed mechanism of this phenomenon cannot, however, be proposed at the moment.

Three possible mechanisms of the release of titanium ions can be proposed as already shown. These results demonstrate that the most significant is corrosion initiation and propagation of corrosion in a presence of local damage of oxide coating and, on the other hand, complex stresses imposed on the screw implants. However, it is necessary to consider whether two other mechanisms can also operate and be comparable.

The dissolution of rutile may occur, but at an extremely low rate and only in strongly acidic environments. Such conditions may occur only in inflammation conditions at which pH may reach highly acidic values. Even if so, the dissolution rate achieves 0.136 mg/kg in 12 months, such results means that the oxide coating even 100 μm thick (after anodic oxidation) decreases less than 1 nm. Such a mechanism is then impossible and must be rejected.

The third mechanism, which can be considered, is diffusion of titanium ions through the rutile lattice. The diffusion of titanium at the temperature of a human body in the rutile crystalline structure seems unlikely. There is no such data even for high temperatures so that it seems desirable to consider the titanium diffusion in other structures. The performed calculations showed the titanium diffusion coefficient as extremely low, even below that for the diffusion of titanium in yttria-stabilized zirconia, the diffusion coefficient at room temperature is below 10−<sup>30</sup> m2/s [45]. It means that the time necessary to diffuse through a 10 nm thick oxide layer would be as high as 1035 s.

Summarizing, it can be said that the only origin of the degradation processes resulting in, among other causes, in a necessity of removal of the dental implant, is the damage of the oxide coating. Such degradation may be attributed to the forces during implantation surgery, but they are likely initiated by the cracks, crevices, and discontinuities already appearing at the manufacturing stage.

#### **5. Conclusions**

The titanium dissolution occurs only by the corrosion tunnels in the oxide layer. The tunnels may be formed by cracks or discontinuities. Such potential corrosion initiation and development sites are already present in new implants and they become operative in applied implants during their use.

The present results demonstrate that among three possible mechanisms such as (i) diffusion of the liquid environment into the cracks and crevices in oxide coating; (ii) chemical dissolution of the titanium oxide layer; and (iii) diffusion of titanium atoms through the oxide layer, the two last processes are very unlikely to cause the damage of dental implants.

The both dissolution of rutile and titanium diffusion through the perfect oxide structure are negligible at the temperature of the human body. However, when the pH value at the implant surface and in the environment of saliva falls locally, the oxide layer starts to dissolve, but even at pH = 3, only a small fraction, 10−<sup>8</sup> of the rutile oxide, may dissolve during 12 months.

**Author Contributions:** Methodology, A.O.; validation, A.O.; investigation, A.O.; resources, A.O.; original draft preparation, A.O.; formal analysis, A.Z.; writing—conceptualization, A.O. and A.Z.; writing—review & editing, A.O. and A.Z. All authors have read and agreed to the published version of the manuscript.

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

**Acknowledgments:** We are grateful to Grzegorz Gajowiec (GUT) for his examinations of oxidized surfaces with the SEM, M.Eng. Gabriel Strugała for his examinations of oxidized surfaces with the CT, Eliza Kurek from Biological and Chemical Research Center University of Warsaw for the ICP-MS studies, and Andrzej Wojtowicz from the Department of Dental Surgery, the Medical University of Warsaw for delivery of the implants.

**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* **Influence of Two-Stage Anodization on Properties of the Oxide Coatings on the Ti–13Nb–13Zr Alloy**

#### **Agnieszka Ossowska 1,\*, Andrzej Zieli ´nski 1, Jean-Marc Olive 2, Andrzej Wojtowicz <sup>3</sup> and Piotr Szweda <sup>4</sup>**


Received: 2 June 2020; Accepted: 20 July 2020; Published: 22 July 2020

**Abstract:** The increasing demand for titanium and its alloys used for implants results in the need for innovative surface treatments that may both increase corrosion resistance and biocompatibility and demonstrate antibacterial protection at no cytotoxicity. The purpose of this research was to characterize the effect of two-stage anodization—performed for 30 min in phosphoric acid—in the presence of hydrofluoric acid in the second stage. Scanning electron microscopy, atomic force microscopy, energy-dispersive X-ray spectroscopy, X-ray diffraction, Raman spectroscopy, glow discharge optical emission spectroscopy, nanoindentation and nano-scratch tests, potentiodynamic corrosion studies, and water contact angle measurements were performed to characterize microstructure, mechanical, chemical and physical properties. The biologic examinations were carried out to determine the cytotoxicity and antibacterial effects of oxide coatings. The research results demonstrate that two-stage oxidation affects several features and, in particular, improves mechanical and chemical behavior. The processes influencing the formation and properties of the oxide coating are discussed.

**Keywords:** titanium alloys; electrochemical oxidation; nanotubular oxide layers; microstructure; nanomechanical properties; corrosion resistance; wettability; antibacterial protection; cytotoxicity

#### **1. Introduction**

Titanium and its alloys—due to their mechanical properties—excellent corrosion resistance, and a high strength/density ratio, are nowadays the most appropriate materials for load-bearing implants and biomedical materials [1,2] used, e.g., in arthroplasty [2,3], as dental implants [4–6] and dental prostheses [7]. The titanium and its alloys proposed for medicine, after their oxidation, include medical titanium [8–10], Ti–6Al–4V [11,12], Ti–6Al–7Nb and Ti–13Nb–13Zr [13] alloys. The most commonly used Ti–6Al–4V alloy contains alloying elements, which may provoke undesirable tissue reactions damaging nerves cells, softening the bones, and, as a consequence, resulting in the appearance of diseases of the circulatory and central nervous systems [14,15]. Therefore, Ti-13Zr-13Nb alloy was chosen for this research as it has no harmful elements and possesses a low (76 GPa) Young's modulus, similar to that of cortical bone, providing better stress distribution at the implant–bone contact zone and preventing against loosening and damage of the implant.

The different surface modification methods of titanium alloys such as deposition of coatings, oxidation, ion beam surface modification, ion implantation, titanium plasma spraying, acid etching, grits blasting, sandblasting followed by acidic etching, electropolishing and laser melting were applied for titanium and its alloys for biomedical applications [16]. In particular, adhesion of the cells has been shown better on rough than on smooth surface [17,18]. Such surface characteristics may be achieved by the laser treatment [19,20], surface mechanical attrition [21,22], acid etching [23,24], deposition of phosphate [25–29] and composite coatings [30–32]. Among those modifying approaches, the oxidation remains essential as it can either form nanotubular oxide structures or rough oxide surfaces, enhancing the adhesion of osteoblasts, if alone, and deposition of coatings, if used as an interlayer.

Oxidation plays an essential role among possible surface engineering methods, and even a spontaneously formed titanium oxide layer is a barrier limiting the entry of metal ions into tissues [33]. Moreover, the oxide layer may influence the osteoinduction processes by a change in the architectural features and chemical composition of the oxides [34]. The techniques used for this purpose include the low voltage anodization, micro-arc oxidation (MAO) [35], thermal oxidation of titanium biomaterials [36], less often the oxidation using hydrogen peroxide [37] and laser-enhanced oxidation [38].

The MAO is used as a technique for creating a multiporous or highly developed surface, often implemented with different ions [39–42]. A novel "cortex-like" micro/nano dual-scale structured TiO2 coating was prepared in such a way in tetraborate electrolytes [43]. The MAO in ammonium acetate was resulted in a multiporous, crystalline titanium oxide layer demonstrating the apatite forming ability [9]. The antibacterial activity may be achieved by the MAO performed in electrolytes comprising Ag, Cu or Zn [40]. However, this technique needs a high voltage and results in thick oxide coatings.

The low potential electrochemical oxidation may result in either the compact oxide [44] or nanotubular oxide coatings, depending on the type of electrolyte and anodization parameters [45–55]. The TiO2 nanotubular surface provided topography favorable for improving the clinical performance of implants when comparing to the sand-blasted acid-etched topography [56]. The individual nanotubes can be filled with antibiotics or nanometals for introducing the antibacterial ability. The release rate of nanosilver depends on its placement: relatively fast release was observed for nanoAg inside the nanotubes and gradual release, for Ag inside the cavities [57]. The functionalization of titanium dioxide nanotubes with some biomolecules was developed for biomedical applications [58] and the osteogenic differentiation can be modulated by various additional treatments of nanotube coatings on Ti–6Al–4V implants [12]. Superhydrophobic titanium oxide nanotube arrays may serve as the drug reservoir, and ultrasonic waves may trigger the drug release [59]. Such a superhydrophobic Ti surface was fabricated by subsequent anodization in H2O2 followed by aging [60]. Hierarchical structures were obtained, applying two nanotexturing surface treatments onto titanium coatings, anodic oxidation and alkaline treatments, and the simultaneous presence of micro-/nano-roughness resulted in a distinct increase in cell proliferation [61].

Different composite coatings were also developed. The ion implantation of helium ions was made on the oxide film obtained by previous anodization to improve hydrophilic properties [8]. The decoration of previous titanium oxide nanotubes with MnO increased the ability to form apatite [62]. Osseointegration was enhanced by coating the titanium implants with a nanostructured thin film comprised of titanium carbide and titanium oxides clustered around graphitic carbon [63].

The oxides obtained by low voltage anodization can be in the form of thin, compact coatings or nanotubular layers. In the past, the bi-layer coating was prepared [47] by gas oxidation of titanium alloy and then electrochemical oxidation resulting in nanotubular layers grown on the previous compact oxide layer. Such treatment brought out the highly corrosion-resistant coatings but possessing the relatively short nanotubes. Therefore, the present research has applied the two-stage electrochemical oxidation assuming that such procedure may positively affect some properties of the oxide coatings, in which the nanotubular layer is formed not in the bare metal, but in the compact oxide layer. In particular, the mechanical and chemical behavior have been expected to improve.

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

The study was performed on a two-phase titanium alloy Ti–13Nb–13Zr of chemical composition listed in Table 1. The microstructure of the investigated alloy (Baoji SeaBird Metal Materials Co., Ltd., Baoji, China) is shown in Figure 1. It is a β-phase structure comprising of the α' phase being a supersaturated solution with a slightly stubborn effect and a martensitic structure, which is formed as a result of rapid cooling from the temperature of the β-phase stability or as a result of plastic deformation.

**Table 1.** Chemical composition of the Ti–13Nb–13Zr alloy, weight percent (according to manufacturer's certificate).

**Figure 1.** Microstructure of the Ti–13Nb–13Zr alloy after etching with the Kroll solution composed of 2 mL HF (40 wt.%), 2 mL HNO3 (55 wt.%) + 96 mL H2O.

The specimens of dimensions 15 mm × 10 mm × 4 mm were cut from the alloy sheet of initial thickness of 4.2 mm. Then the samples were ground with abrasive papers (No. 2500 as the last, Struers, Inc., Cleveland, OH, USA). Afterward, the specimens were cleaned in an ultrasonic chamber (Sonic-2, Polsonic Palczynski Sp. J., Warsaw, Poland) with isopropanol, methanol (Avantor Performance Materials Poland S.A., Gliwice, Poland) and distilled water, subsequently, for 5 min in each batch and finally dried in cold air.

The tests were performed in a standard circuit composed of an electrochemical cell, power supply (SPN-110-1C, MPC Lab Electronics, Nijmegen, The Netherlands), Pt electrode as the polarizing electrode and the tested metallic electrode. Neither stirring, aeration nor deaeration were applied. All measurements were performed at room temperature. The anodization parameters were set up based on some earlier investigations [47]; in particular, even if the electrochemical oxidation time has a small effect on the oxide thickness, the 30 min period was assumed necessary to perform the electrochemical oxidation at equilibrium conditions.

The anodization was carried out at the first stage electrochemically in 1 M orthophosphoric acid (H3PO4) at the potential value of 40 V, in one step, at 20 ◦C, for 30 min (samples obtained in such a way are here designated as EO1). The electrochemical oxidation was repeated in 1 M orthophosphoric acid with an addition of 0.3 vol.% of hydrofluoric acid (HF) (designation EO2). The process was performed again at 20 ◦C, at a potential value of 20 V, in one step, for 30 min. The coatings were also obtained

by two-stage oxidation—first EO1, then EO2 (designation EO1 + EO2). After each of the processes, the samples were rinsed in distilled water and dried in cold air. The samples were heat-treated after oxidation at 400 ◦C for 2 h in the air (humidity <70%).

The surfaces of specimens and their cross-sections after each form of oxidation were examined with the scanning electron microscope (SEM JEOL JSM-7600 F, JEOL, Ltd., Tokyo, Japan), equipped with a LED detector, at 5 kV acceleration voltage. The chemical composition of the coatings was determined using an X-ray energy-dispersive spectrometer (EDS, Edax, Inc., Mahwah, NJ, USA).

The surface examinations, with used linear roughness measurement, were performed with the atomic force microscopy (MFP-3D, Oxford Instruments Asylum Research Inc., Santa Barbara, CA, USA) at the Université Bordeaux, France. The surface topography was assessed in the noncontact mode at a force 50 mN. The roughness index *R*<sup>a</sup> was estimated within an area of 5.0 μm × 5.0 μm.

The X-ray diffraction studies were carried out at the Gda ´nsk University of Technology, Faculty of Applied Physics and Mathematics, with the use of X-ray diffractometer (Philips X'Pert Pro–MPD, Brighton, UK) system with a vertical T–T goniometer (190-mm radius). The X-ray source was a long-fine-focus, ceramic X-ray tube with Cu anode. The standard operating power was 40 kV, 50 mA (2.0 kW). The system optics consisted of programmable divergence, anti-scatter and receiving slits, incident and diffracted beam Soller slits, curved graphite diffracted beam monochromator and a proportional counter detector (Bragg–Brentano parafocusing geometry (2θ ca. 5◦–100◦). The spectroscopic examinations of the grown oxide layers were performed with the Raman spectrometer (Horiba Jobin Yvon Gmbh, Bensheim, Germany) at the Max Bergmann Centrum of Biomaterials, Dresden Technical University.

The glow discharge optical emission spectroscopy (GDOES) tests were carried out at the University of Bordeaux, using the GD-Profiler 2 (Horiba Jobin Yvon IBH Ltd., Glasgow, UK). The measurements were performed using the following process parameters: a glow discharge source (argon plasma) at 700 Pa and 30 W, measurement time 120 s.

The nanoindentation tests were performed with the NanoTest Vantage (Micro Materials, Wrexham, UK) equipment using a Berkovich three-sided pyramidal diamond. The maximum applied force was equal to 5 mN, the loading and unloading times were set at 20 s, the dwell period at full load was 10 s. The distances between the subsequent indents were 50 μm. During the indent, the load–displacement curves were determined using the Oliver and Pharr method. Based on the load–penetration curves, the surface hardness (*H*) and reduced Young's modulus (*E*) were calculated using the integrated software. The critical process parameters included the maximum force, holding time and test rate. In calculating Young's modulus (*E*), a Poisson's ratio of 0.3 was assumed for the titanium oxide layer. The measurements were processed in randomly selected five points for each surface, and the results were averaged.

The electrochemical measurements of corrosion parameters were performed by a potentiodynamic mode in the Ringer's solution. The simulated body fluid was obtained by dissolving a Ringer's tablet (Merck KGaA, Darmstadt, Germany; each tablet contained 1.125 g NaCl, 0.0525 g KCl, 0.03 g CaCl2 and 0.025 g NaHCO3) in 0.5 L of distilled water at 20 ◦C. Different pH levels were obtained by adding the hydrochloric acid (5 wt.% to the solution. Lowering of pH even to 3 resembled acidic environmental conditions during inflammation [64], so the test was carried for pH ranging from 7 (normal physical state) through 5 to 3 (inflammatory state). A standard three-electrode electrochemical cell was used comprising of a saturated calomel electrode (SCE) as the reference electrode, a platinum electrode as the counter electrode and the sample as the working electrode (anode). All experiments were performed using a potentiostat/galvanostat (VersaSTAT 4, Ametek Scientific Instrumentation, Leicester, UK). Before the test, the samples were stabilized at their open circuit potential (OCP) for 0.5 h. Potentiodynamic polarization tests were carried out at a potential change rate of 10 mV/min, within a scan range from −2 to 2.5 V. The corrosion potential *E*corr and corrosion current density *i*corr were determined from the polarization curves using the Tafel extrapolation method.

The water contact angle (wettability) measurements were taken for the reference Ti–13Nb–13Zr alloy and oxidized specimens using a contact angle goniometer (Attension Thete Lite, Dyne Technology, Lichfield, UK) at room temperature. All analyses were repeated three times for each sample.

Studies of antibacterial activity of nanotubular surfaces were carried out at the Gda ´nsk University of Technology, Faculty of Chemistry, with the *Staphylococcus aureus* ATCC25923 strain. The samples were put in 5 mL of bacterial suspension (containing at least 10<sup>6</sup> colony forming units (CFU) in 1 cm3) prepared in phosphate buffered saline (PBS, chemical composition 8.0-g/L NaCl, 0.2-g/L KCl, 1.44-g/L Na2HPO4, 0.24-g/L KH2PO4)), in which they stayed for 1 min. This step of the procedure aimed to allow the bacterial cells to adsorb on the surfaces of the tested materials. Next, the samples (with bacteria adsorbed on their surfaces) were transferred to 5 mL of sterile TSB medium placed in the 8-well microplates. The samples were incubated at 37 ◦C for 24 h (one day) or 120 h (5 days). Subsequently, the samples were removed carefully from TSB medium and rinsed by submersion three times in a sterile saline solution (0.9% NaCl). Afterward, the samples were placed in the wells of a new titration plate containing 5 mL of MTT (3-(4,5-dimethyl-2-thiazolyl)-2,5-diphenyl-2*H*-tetrazolium bromide) solution (0.3%) in PBS. The living cells of bacteria reduce MTT to insoluble in water violet formazan crystals, and the amount of formed formazan is proportional to the number of live bacteria that are still present (in the form of biofilm) on the surfaces. Following 2 h incubation at 37 ◦C in the dark, the solution of MTT in PBS was carefully removed from the wells and replaced with 5 mL of DMSO for dissolving formed formazan crystals. The optical density of the obtained solutions was measured at 540 nm using a Victor<sup>3</sup> microtiter reader (PerkinElmer, Waltham, MA, USA).

Cytotoxicity tests were performed at the Warsaw Medical University, Department of Dental Surgery. They were carried out on the titanium alloy and the oxidized surfaces of the samples. Experiments were performed on fibroblasts obtained from neonatal rat Lewis Op/Op after the third passage. A small microscope slide was placed into small plastic plates with a diameter of 35 mm (430165, Corning Manufacturer, Corning, NY, USA). For all of them, except for control plates, single titanium samples were filled with a suspension of cells in the culture medium. All plates received 100,000 cells suspended in 2.0 mL medium. After five days, the slides with the cells deposited on them were rinsed with physiological saline and preserved in a mixture of methanol and acetic acid (3:1) for 5 min, then stained with hematoxylin and eosin. The preparations were dehydrated with DPX (a mix of distyrene, a plasticizer and xylene), dried and subjected to microscopic evaluation. The density of cultured cells and their morphologic features, as well as the presence of forms of mitotic divisions, were assessed. The evaluation of each sample was carried out three times.

#### **3. Results**

#### *3.1. Microstructure, Surface Topography, Phase and Chemical Compositions*

Figure 2 presents the morphology of the oxide coatings. They all were homogenous and transparent, but the interference of reflected light resulted in a color effect related to the applied voltage and resultant thickness of the oxide layer and its structure. The samples after EO1 treatment showed a blue color (Figure 2a), typical of titanium oxidized at 40 V and resulting in the thickness of about 74 nm [65]. The oxide coatings obtained after EO2 treatment and EO1 + EO2 modification were matt-gray as expected for nanotubular layers.

The observations of the surfaces of oxide coatings revealed a homogenous and even surface after EO1 treatment (Figure 2a) and an appearance of nanotubes after EO2 (Figure 2c) and EO1 + EO2 (Figure 2e) surface modifications.

The measurement of the thickness of the thin EO1 coating was challenging as before oxidation, the surface was mirror-like, and there were reflections from the surface. The results of measurements based on the cross-sections of the samples (Figure 2b,d,f) showed that the EO1 coating was about 80-nm-thick (Figure 2b), in perfect accordance with the previously cited report, the EO2 coating was about 1000-nm-thick (Figure 2d), and the EO1 + EO2 coating had similar thickness. However, the last

coating could be supposed to compose of two zones: typical nanotubular outer layer and an inner layer of presumably different view as discussed later (Figure 2f).

**Figure 2.** Images SEM - comparison of (**a**) alloy after amorphous layers (EO1) treatment (surface); (**b**) alloy after EO1 treatment (cross-section); (**c**) alloy after nanotubular layer (EO2) treatment (surface); (**d**) alloy after EO2 treatment (cross-section); (**e**) alloy after EO1 + EO2 treatment (surface), (**f**) alloy after EO1 + EO2 treatment (cross-section).

Figure 3 shows the surface coating topography after different electrochemical oxidation and Table 2 presents the roughness of coatings. After EO1 treatment, the smoothest layer, even compared to the polished material, was observed. The EO2 (not shown in figure) and EO1 + EO2 coatings were characterized by slightly increased roughness than the EO1 and substrate material.

**Figure 3.** Atomic Force Microscope (AFM) images of the surface topography. (**a**) Substrate titanium alloy; (**b**) alloy after EO1 treatment; (**c**) alloy after EO1 + EO2 treatment; (**d**) a single nanotube.

**Table 2.** Roughness parameters (*R*a) in the area of 200 <sup>×</sup> 200 nm2.


The chemical composition of the layers determined by EDS measurements is demonstrated in Table 3. However, because of the oxide volume and thickness examined by the EDS, which exceeds that of the oxide coating of the EO1 sample, the data for this specific case could be result from both thin oxide coating and the alloy. The oxygen content in each coating was determined from stoichiometry, assuming that it formed the stoichiometric oxides.



(\*) quantities, in this case, must be regarded as only informative.

The Raman spectra of the titanium alloy are shown in Figure 4. According to previous research [66], the Raman spectra for the EO1 coating should display clear signs of the anatase phase four-peak pattern with peaks at 575 cm−<sup>1</sup> deriving from ν<sup>1</sup> vibrations being the strongest and the other peaks at 144, 198 and 406 cm−<sup>1</sup> being much weaker. These peaks come from anatase [67,68]. However, here, the small rutile band was observed at 238 and 612 cm<sup>−</sup>1. For the EO1 + EO2 coating, the intensity of

the band 313 cm−<sup>1</sup> increased and moved to higher frequencies; these peaks originated from the TiO2 band [68]. A very similar situation was noticed for the 198 cm−<sup>1</sup> peak, which also came from anatase. These findings are in agreement with XRD characterization showing TiO2-specific peaks (for anatase and rutile phases). The Raman spectra of here examined samples are similar to those reported for titanate crystal formed of nanotubes [69].

**Figure 4.** Raman spectra of the oxide coatings for a non-oxidized substrate and after different electrochemical oxidation.

As a result of the GDOES measurements, the values of wavelengths emitted by the excitation of atoms appear for all here present elements, as shown in Figure 5. For the Ti–13Nb–13Zr alloy (Figure 5a), the distribution of elements with erosion (sputtering) time was abrupt and remained at a certain level. For the oxidized EO1 sample (Figure 5b), the maximum intensity for oxygen occurred in the initial phase of the measurement, significantly exceeding the value of the peak derived from titanium, which decayed very quickly. The distribution of the intensity values of particular elements for the EO2 (Figure 5c) and EO1 + EO2 (Figure 5d) coatings was different. In the case of the EO1 + EO2 coating, in the initial phase of the study, there were distinct fluctuations in the intensity of the main alloying elements: Ti, Nb and Zr. The differences in the intensity of the elemental distribution with erosion time are visible, which may be due to different thicknesses of the tested coatings.

**Figure 5.** Glow discharge optical emission spectroscopy (GDOES) analysis results. (**a**) Ti–13Nb–13Zr alloy; (**b**) alloy after EO1 treatment; (**c**) alloy after EO2 treatment; (**d**) alloy after EO1 + EO2 treatment.

Analyzing the XRD diagrams (Figure 6) and based on the literature data [70–72], for each sample the characteristic reflexes corresponding to the positions of the α-Ti phase and β-Ti can be found (Figure 6a). Depending on the sample, they differ in intensity and width. In all tested samples, several reflexes from both crystallographic structures were found at appropriate angles.

For oxidized specimens, the reflexes from anatase were observed at three positions, and from rutile–only at one. Due to the overlapping of peaks from titanium and titanium oxide, the crystalline structure of the titania nanotubes could not be determined. A slight decrease in the intensity of the peaks for the EO1 (Figure 6b) sample was observed. In the case of EO1 + EO2 coating (Figure 6d), it became necessary to reduce the reflex intensities of both phases (α-Ti and β-Ti) in comparison with the EO2 coating (Figure 6c). A much more significant decrease in the intensity of the reflexes from the β-Ti phase was observed compared to the reduction of the α-Ti phase, perhaps due to the presence of oxygen. Most likely, the reflection intensity depends on the thickness of the layer; the thicker the layer, the higher the reflex intensity. This result appeared for all oxidized specimens for which a significant increase in reflex intensity was observed. It is worth noting that the primary reflexes from the oxide phase were distinctly the highest for the EO1 + EO2 sample.

#### *3.2. Nanomechanical Properties*

The nanohardness, microhardness and Young's modulus values of the specimens are shown in Table 4. The tests showed an increase in both hardness values and Young's modulus for oxidized samples. When considering Young's modulus, the highest value was obtained for the EO1 coating. Similar results were obtained for the nanoparticle layer and hybrid coating.

**Figure 6.** XRD spectra. (**a**) Alloy Ti–13Nb–13Zr; (**b**) alloy after EO1 treatment; (**c**) alloy after EO2 treatment; (**d**) alloy after EO1 + EO2.


**Table 4.** Mechanical properties of the tested specimens.

An increase in hardness of one order of magnitude was observed for samples with the compact oxide layer. For samples with the EO2 and EO1 + EO2 coatings, the values were similar, much higher compared to polished alloy or the compact layer. The main factor influencing the change in hardness is the thickness of the oxide layer [73–75]. The small thickness of the compact solid layer (80 nm) could cause some errors in the nanoindentation test. It is well known that the response is not only be given by the first indented layer, but the substrate or the subsequent layers may also contribute to the

indentation response [76]. In both cases, the thickness of the nanotubular layer (EO2) was so large that there was no response from the substrate.

#### *3.3. Wettability*

The results of measurements of water contact angle are presented in Table 5. The decrease of the contact angle was small for the EO1 and significant for the other coatings. The created surfaces were hydrophilic. The most desirable value of contact angle for regeneration applications in hard tissues ranges from 35◦ to 80◦ [77].


**Table 5.** Contact angle for the water droplet for the tested specimens.

#### *3.4. Corrosion Properties*

The corrosion test results are presented in Figure 7. The polarization curves were S-shaped. The decrease in pH value always resulted in a shift of corrosion potential to a more active area. However, the appearance of thin, compact oxide coating slightly worsened the corrosion behavior, and the EO1 + EO2 treatment caused an opposite effect—a shift of the corrosion potential to the more noble area. The corrosion current values can be determined if the Tafel straight lines in potentiodynamic curves (at logarithmic scale) are sufficiently long (at least two decades). Here, this condition has not been fulfilled. Therefore, there is no sufficient base to calculate the corrosion current densities.

**Figure 7.** Potentiodynamic polarization curves at different pH. (**a**) Ti–13Nb–13Zr; (**b**) EO1 sample; (**c**) EO2 sample; (**d**) EO1 + EO2 sample.

#### *3.5. Antibacterial Properties*

Figure 8 presents the images illustrating the intensity of biofilm formation on the surfaces of materials, measured by absorbance values of solutions of formazan (diluted in DMSO) produced by live cells of bacteria from MTT. After one-day exposure, the lowest absorbance values were observed for the reference and the EO1 samples. The presence of nanotubular surface distinctly increased biofilm formation. In contrast, even the biofilm increased in five days, the lowest levels were attained for the EO2 and EO1 + EO2 samples, for which only a slight difference was noticed between the first and fifth days.

**Figure 8.** Growth of bacterial film of *Staphylococcus aureus* strain on the tested samples surfaces.

#### *3.6. Cytotoxicity*

In these tests (Figure 9), five-day-old cultures formed a reasonably even layer of cells (monolayers), with a large number of figures of mitotic divisions: prophase and metaphase, a small number of polymorphs with no signs of cell damage. The images of cell culture on the surface of the non-oxidized alloy, as well as of the oxidized samples, show that none of the studied surfaces deteriorated the behavior of osteoblasts.

**Figure 9.** Cytotoxic tests of fibroblastic cells. (**a**) Ti–13Nb–13Zr alloy; (**b**) EO1 sample; (**c**) EO2 sample; (**d**) EO1 + EO2 sample.

#### **4. Discussion**

The two-stage oxidation could result in the bi-layer ("sandwich" layer) oxide coating as shown in [47] in which the alloy was subjected to gaseous oxidation and then to the electrochemical oxidation. The last method performed in the presence of HF acid caused the transformation of the upper part of the compact oxide coating into a nanotubular layer, resulting in both highly corrosion-resistant and bioactive coating. However, such a mechanism is possible if oxidation is performed at a high temperature at which oxygen diffusion is fast. On the other hand, with increasing compact oxide thickness, the thickness of the nanotubular layer decreases to zero because of increasing electrical resistance.

Therefore, we have attempted to create the compact oxide layer by the electrochemical method and transform it into a nanotubular layer. The thickness of the oxide coating appearing at room temperature is low because of slow oxygen diffusion, and the thickness of the nanotubular layer creating in the second stage overpasses that of the compact oxide. Therefore, at applied oxidation parameters, the creation of the bilayer coating is not possible. The inner layer shown in Figure 2f may be simply a part of the nanotubular layer, likely, as sometimes observed [71], the interface between the oxide layer and bare metal.

The mechanism of creation of oxide coating during two-stage anodization involves the appearance of a nanotubular layer in the second stage; initially, within the compact oxide layer and then, when the compact layer is fully transformed into the nanotubular structure, within the bare alloy structure. No double structure was observed in any image; however, it affected mechanical and chemical, but not biologic properties.

In the beginning, let us consider the similarities and differences between behavior of specimens subjected to EO2 treatment (nanotubular oxide on bare metal) and EO1 + EO2 oxidation. Neglecting the doubtful inner layer in the EO1 + EO2 sample, it may be said that the form of applied procedure does not affect the coating thickness and roughness. Moreover, the Raman spectra for EO1 + EO2 and EO2 also look very similar in the shape presenting the peaks at the same positions. Raman spectra confirm that the heat treatment of the nanotubes transforms them from the amorphous to the crystalline structure. However, the thickness of the compact oxide layer is much lower than that of the nanotubular layer. It is because the single electrochemical oxidation is determined only by oxygen diffusion, and the growth of the nanotubular layer is much faster as determined by the chemical reaction of etching, being relatively quick. The appearance of the nanotubular structure is an obvious explanation of surface roughness distinctly higher than that of the compact oxide.

The question is, what causes the difference between EO2 is and EO1 + EO2 coatings in some their properties? It may be microstructure, but such results are very difficult to observe. We believe that it is an enrichment with oxygen of the upper zone of the oxide coating after EO1 + EO treatment. The EDS examinations do not confirm this assumption, but their precision is low. However, the GDOES experiments—in which the slower erosion rate and higher oxygen intensity in the EO1 + EO2 coatings compared to EO2 coating—may be an evidence of above proposed phenomenon. Physically, oxygen in the previous oxide layer may occupy also interstices making a microstructure more resistant to diffusion of other elements, more mechanically resistant and more resistant to corrosion by creating the oxides in existing imperfections.

The XRD results are similar for both EO2 and EO1 + EO2 oxidation procedures. The appearance of both rutile and anatase in the oxide layers has been detected what may be surprising as such transformation may occur at temperatures beginning from 450 ◦C [72], 500 ◦C [49] or even 550 ◦C [71] for titanium. Here the heating temperature was 400 ◦C, close to the suggested beginning of transformation.

The nanoindentation tests show that the presence of nanotubular structure significantly increases hardness and Young's modulus and decreases plastic work. Such results are supported by earlier research [78]. They may result from the specific microstructure of the nanotubular layer, which is comprised of very hard nanotubular oxides, flexible and readily underwent slight deformation. The difference between those values for the EO2 sample and EO1 + EO2 coating may be attributed to the effect of higher oxygen/oxides content in the last sample surface.

The contact angle measurements show an increase in wettability for both EO2 and EO1 + EO2 coatings. The present values permit to classify both surfaces as potentially biocompatible, which may attract proteins and pre-osteoblasts. There have been some assumptions about the best values contact angle for cell attachment assessed at 55◦ and for bone regeneration from 35◦ to 80◦ [70,77]. Here found contact angle values, about 29◦ for nanotubular layer and 49◦ for hybrid coating are favorable for biochemical adsorption processes.

The corrosion resistance of titanium alloys is well-known to increase after oxidation. The behavior of the potential beginning from a fast decrease rate followed by a slow rise is typical for a partial stabilization of current density and formation of the highly protective passive film. The current density stabilization suggests the passive film breakdown, in a way similar to what occurs during pitting nucleation and repassivation. Electrochemical potentiodynamic studies have shown how significantly the corrosion quality is affected by the surface quality and thickness of the obtained layers [79,80]. An increase in corrosion resistance of samples covered with amorphous layers (EO1) is visible. The compact, uniform layer provides better corrosion resistance, comparable with spontaneous oxide layers. In the case of EO2 layers (Figure 7c), no decrease in resistance to corrosion was expected. The nanotubes are hollow, and there are voids between them, which were corrosion tunnels - potential places for corrosion development. Saji et al. [81] have observed an increase in the resistance of the oxide layer with a nanotube structure. However, for the Ti–Nb–Ta–Zr alloy, the same authors observed a decrease in corrosion resistance of the nanotubular layer [82]. The thickness of such a layer may play an important role, as it is a barrier to the progression of corrosion [80]. From performed research it follows that a nanotubular layer does not form a suitable protection against corrosion. However, the EO1 + EO2 coating obtained by the electrochemical method shows nobler corrosion potential and likely better corrosion resistance compared to the nanotubular EO2 layer.

The attachment of bacteria to the solid surfaces of different chemical compositions, including alloy surfaces, is affected by the electrostatic double layer, hydrophobicity, roughness and various other factors [83–85]. Bacteria need to overcome the energy barriers to reach the negative energy regions, thereby facilitating the bacterial attachment [86]. In our experiment, we observed a rather moderate development of the population of bacterial cells attached (a bacterial biofilm) to the surfaces of all materials tested after 24 h of incubation. It confirms that some cells of *S. aureus* were able to adsorb on the surfaces of all materials. The MTT assay revealed a bit higher number of bacteria on surfaces of EO2 and EO1 + EO2 compared to EO1 and Ti–13Nb–13Zr. However, no drastic differences were observed between oxidized and non-oxidized specimens in the level of bacterial content after 1st day of incubation.

In contrast, the bacteria development on the non-oxidized alloy drastically increased during the following five days of incubation up to their highest level. In the case of this material, the value of absorbance in MTT assay, which is a consequence of the number of bacteria and the amount of reduced MTT to formazan, reached a value of about 1.5 compared to 0.6, for EO1 and about 0.4 for EO2 and EO1 + EO2. Moreover, in the case of the samples subject to EO2 or EO1 + EO2 oxidation, the levels of bacterial content (biofilms) after one and five days were comparable, which confirms bacterial growth inhibition on the surfaces of these materials which is their essential advantage. The observed influence of the presence of a nanotubular oxide layer on the bacteria attachment or growth inhibition was already reported [87,88]. It may be attributed to the influence of surface topography on the adhesion of bacteria, and it is an evidence that the presence of oxide nanotubes prevents to some extent, thanks to the specific layer microstructure, the danger of bacteria inflammation. The different strategies to avoid infection onto titanium surfaces have been reported: surface modification and coatings by antibiotics, antimicrobial peptides, inorganic antibacterial metal elements and antibacterial polymers [89], but a presence of nanotubular crystalline titanium dioxide also could be useful.

The introduction of nanomaterials and nanostructures may affect the osseointegration processes [90], but also may develop cytotoxicity as several nanomaterials. In our tests (Figure 9), after five days, no cytotoxic effects against the osteoblasts were noticed. These results are following some previous reports [34].

Summarizing it can be said that the application of oxidation joined with etching by HF acid has a different effect on mechanical and chemical properties depending on whether the alloy or oxide coating form the surface zone. The difference can be attributed to the different microstructure within the surface zone, even if the processes leading to such a result cannot be precisely described yet. It may be a saturation of oxides with oxygen or an appearance of more close-packed nanotubes or other not recognized phenomena. This problem will be investigated in detail in the near future.

#### **5. Conclusions**

Two-stage Ti–13Nb–13Zr electrochemical oxidation with the use of orthophosphoric acid and subsequently hydrofluoric acid results in an improvement of several nanomechanical and chemical properties such as hardness, Young's modulus and corrosion resistance. No significant effects on biologic properties are observed.

The observed influences may be attributed mainly to the change in the chemical composition and microstructure of the upper zone of the nanotubular layer, inside which the formation of nanotubes occurs not inside the bare alloy, but in the previous anatase layer.

Following the positive effects of present tests, future research will be aimed at recognizing and modeling processes that occur during the formation of titanium oxide nanotubes on previously oxidized alloy, developing the nanotubes of thickness comparable to that of the compact oxide layer, without loss of bioactivity.

**Author Contributions:** Conceptualization, A.O.; methodology, A.O., J.-M.O., A.W., P.S.; validation, A.O., J.-M.O., P.S.; formal analysis, A.Z.; investigation, A.O., J.-M.O., P.S.; resources, A.O., J.-M.O., A.W., P.S.; writing—original draft preparation, A.O.; writing—review and editing, A.O. and A.Z.; All authors have read and agreed to the published version of the manuscript.

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

**Acknowledgments:** We are grateful to Grzegorz Gajowiec (GUT) for his examinations of oxidized surfaces with the SEM and EDS, Maria Gazda (GUT) for the research and with XRD, Dieter Scharnweber and his Group from Max-Bergmann-Centrum of Biomaterials, Dresden Technical University, for Raman spectroscopy and the research staff of the Institut de Mécanique et d'Ingénierie de Bordeaux for the AFM and GDOES tests.

**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* **Nanotubular Oxide Layer Formed on Helix Surfaces of Dental Screw Implants**

#### **Magdalena Ja ˙zd ˙zewska \* and Michał Bartma ´nski**

Faculty of Mechanical Engineering and Ship Technology, Gda ´nsk University of Technology, Narutowicza 11/12, 80-233 Gda ´nsk, Poland; michal.bartmanski@pg.edu.pl

**\*** Correspondence: magdalena.jazdzewska@pg.edu.pl; Tel.: +48-58-347-17-96

Received: 11 December 2020; Accepted: 18 January 2021; Published: 20 January 2021

**Abstract:** Surface modification is used to extend the life of implants. To increase the corrosion resistance and improve the biocompatibility of metal implant materials, oxidation of the Ti-13Nb-13Zr titanium alloy was used. The samples used for the research had the shape of a helix with a metric thread, with their geometry imitating a dental implant. The oxide layer was produced by a standard electrochemical method in an environment of 1M H3PO4 + 0.3% HF for 20 min, at a constant voltage of 30 V. The oxidized samples were analyzed with a scanning electron microscope. Nanotubular oxide layers with internal diameters of 30–80 nm were found. An analysis of the surface topography was performed using an optical microscope, and the Sa parameter was determined for the top of the helix and for the bottom, where a significant difference in value was observed. The presence of the modification layer, visible at the bottom of the helix, was confirmed by analyzing the sample cross-sections using computed tomography. Corrosion tests performed in the artificial saliva solution demonstrated higher corrosion current and less noble corrosion potential due to incomplete surface coverage and pitting. Necessary improved oxidation parameters will be applied in future work.

**Keywords:** nanotubular oxide; helix surfaces; dental implants; roughness; corrosion properties

#### **1. Introduction**

Titanium and its alloys are nowadays among the most popular biomaterials, called the "gold standard" for endosseous dental implants, even if some adverse reactions may be expected. They possess a lot of important properties, such as their low density, suitable fatigue strength, Young's modulus and specific tensile strength, high resistance to brittle cracking, high corrosion resistance, and the best biocompatibility. Despite that, titanium and its alloys need surface modifications for early osseointegration [1,2]. The type of commercial implant determines surface topography and differences in geometry [3].

Surface modification is nowadays an obligatory treatment of dental implants. Bioactivity of the surface resulting in adhesion of osteoblasts and bone ingrowth can be achieved by the development of surface roughness, creation of bioactive films, and deposition of coatings [2]. Many different methods have been used to change the surface roughness of dental implants, including mechanical techniques such as grinding, polishing, machining, sandblasting and attrition, chemical etching in acids, alkali and fluorides, electrophoretic deposition, and laser treatments [4–11].

The interaction of cells and adsorption of proteins depends on surface structure and is significant in the presence of nanometric pores, which increase the rate of osseointegration and biomechanical fixation [2,7,8,12–17]. A significantly higher bone contact of 27% (*p* < 0.05) was observed in nanotextured compared to machined implants [18]. However, reproducibility of nanoscale surface profiles of titanium with chemical modifications such as acid-etching is quite difficult to achieve and unreliable, and knowledge on the ideal surface roughness parameters for rapid osseointegration is still lacking [19,20].

Among various surface treatments, artificial oxidation seems particularly plausible for titanium dental implants resulting in high corrosion resistance and biocompatibility [21]. The oxidation required for dental implants is currently mostly applied by micro-arc oxidation (MAO) [22,23] and electrochemical oxidation [24]. Gaseous oxidation has also been proposed [25]. MAO induced titanium oxide formation in the anatase crystalline phase and also incorporated Ca, P, and Mg in the film [26–33]. An oxide thickness of 600–1000 nm demonstrated significantly stronger bone responses than that of 17 or 200 nm [30]. The coatings comprising nano TiO2 and nanohydroxyapatite (nanoHAp) demonstrated a torque value of coated screws significantly greater than that of nanoHAp covered screws [34].

The creation of nanotubular oxide layers on titanium and its alloys is well-known. The formation of nanotubular oxide structures on dental implants has not often been investigated and developed. The overly short life of dental implants observed proves the ineffectiveness of the applied surface modifications and provides prompts for further research. The anodization depends to a great extent on the geometry and structure of the surfaces involved. Indeed, the formation of titanium dioxide nanotubes on flat titanium surfaces, provided by well-known suppliers, does not have the same effect on titanium implants, mainly due to the geometry of the implant, which changes the priority, intensity, and interconnection of the electrochemical processes [35]. Nanotubular oxide layers have been reported to increase the bioactivity of titanium implants [36–38], the nucleation and growth of hydroxyapatite coatings [39], and to introduce antibacterial effects after loading the nanotubes with drugs [40,41]. Such a type of surface was already fabricated on the nontoxic Ti-13Zr-13Nb alloy investigated here [42,43]. This research was aimed at an assessment of the creation of nanotubular oxide layers on screw fixed dental implants and the characterization of the layers obtained on the tops and bottoms of helices of implants.

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

#### *2.1. Material*

The biphase α + β Ti-13Zr-13Nb alloy (SeaBird Metal Materials Co., Baoji, China) with chemical composition presented in Table 1 was investigated in the as-received state.

**Table 1.** The chemical composition of the Ti-13Zr-13Nb alloy, wt.%. (based on the manufacturer's certificate).


#### *2.2. Preparation of Specimens*

Round specimens of height 9 mm and diameter 8 mm were prepared by precision milling. The metric thread was cut on all specimens. The remaining impurities were cleaned and the surface was prepared by sand blasting with corundum for 15 s. The cleaning was performed at Aesculap Chifa Ltd. in Nowy Tomysl, Poland. Immediately before oxidation, the specimens were washed in an ultrasonic bath (Sonic 3, Polsonic, Warsaw, Poland) in isopropanol (POCH, 99.8%, Gliwice, Poland) for 10 min, in distilled water for 3 min, and methanol (POCH, 99.8%, Gliwice, Poland) for 10 min.

#### *2.3. Electrochemical Oxidation*

The oxidation was performed using a direct current power supply (MCM/SPN110-01C, Shanghai MCP Corp., Shanghai, China). The specimen tested was connected to the power supply as an anode and the Pt electrode was used as a cathode. The electrolytic bath contained a solution of 150 mL of distilled water, 20 mL of 1 M H3PO4, and 1.5 mL of 0.3% HF (both from POCH, Gliwice, Poland). The distance between the electrode tested and the Pt electrode was 15 mm. The solution was neither aerated nor deaerated, and non-stirred. The experiments were carried out at ambient temperature. The experiments were performed at a voltage of 20 v for 30 min based on previously conducted experiments [42].

#### *2.4. SEM Surface Examination*

Scanning electron microscopy (SEM JEOL JSM-7800 F, JEOL Ltd., Tokyo, Japan) instrument equipped with EDS chemical analyzer (Edax Inc., Mahwah, NJ, USA).

#### *2.5. Light Microscopy Assessment of Roughness*

A light microscope (VHX-7000, Keyence, Osaka, Japan) was applied to examine the surface topography. Roughness parameters determined by the 3D Form Measurement software were applied to the Sa area.

#### *2.6. Computer Tomography*

Tomographic images were obtained using Phoenix v/Tome/xs computer tomography (General Electric, Lewistown, PA, USA).

#### *2.7. Corrosion Examinations in Simulated Body Fluid*

Corrosion tests were performed by a potentiodynamic method in simulated body fluid (SBF) at a temperature of 38 ◦C. The electrochemical measurements were achieved by using a potentiostat/galvanostat (Atlas 0531, Atlas Sollich, Gda ´nsk, Poland). An artificial saliva solution (SBF) was prepared according to EN ISO 10993-15 [44] by dissolving reagent grade chemical (NH2)2CO (0.13 gL-1), NaCl (0.7 gL-1), NaHCO3 (1.5 gL-1), Na2HPO4 (0.26 gL-1), K2HPO4 (0.2 gL-1), KSCN (0.33 gL-1), KCl (1.2 gL-1) (POCH, Gliwice, Poland). The potential was measured vs. a saturated calomel electrode (SCE) located in the Haber-Luggin capillary. As a counter electrode, a standard platinum electrode was used. The test specimen was stabilized in a solution of artificial saliva for 30 min at open circuit potential OCP. The potential change rate was 1 mV/s within a scan range of −2000 to 1000 mV. The solutions were agitated with a magnetic stirrer. Using the Tafel extrapolation method, the corrosion potential (*E*corr) and corrosion current density (*i*corr) values were determined.

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

#### *3.1. Substrate Specimens*

The surface of the non-oxidized alloy is shown in Figure 1 at two different magnifications. The relatively smooth surface and screw lines can be seen.

**Figure 1.** Surface of reference specimen at different magnification: (**a**) 130×, (**b**) 1700×.

#### *3.2. Oxidized Specimens*

Figure 2 shows the appearance of the nanotubular oxide layer only in the area at the bottom of the helix. The pores created are spherical and longitudinal. They possess a diameter ranging between 30 and 80 nm. The layer is well adjacent to the substrate and it has a small number of cracks and surface defects. The gradual decrease of the nanoporous layer and its absence at the top of the helix may be attributed to different current densities, different electrochemical potential, and as a consequence a different course of chemical reactions. The current is screened at the bottom at a given potential and the resultant value is sufficient for electrochemical oxidation to occur. At the tops, the current density is too high and the nanotubes formed undergo fast oxidation, its rate exceeding that of chemical dissolution resulting in nanotubes. The current density is higher at the tops of such surfaces, with the effect attributed to the difficult transport of oxygen to and reaction products from this area, and stepwise depolarization of the area close to the bottom followed by a change in open circuit and corrosion potentials. These processes can shift the current and potential values beyond those necessary to form nanotubular oxide layers.

**Figure 2.** Surfaces of specimens oxidized in an electrochemical way: (**a**) view of specimen, (**b**) surface of helix top, (**c**) surface between top and bottom of the helix, (**d**) bottom of the helix (with different magnifications—d', with the result of measuring the diameter of nanotubes).

The EDS examination results presented in Table 2 suggest the obtaining of a layer of titanium oxide on the surface, which is confirmed mainly by the content of titanium and oxygen. High P content results from the absorption of HPO4 <sup>2</sup><sup>−</sup> anion within the layer pores and it is desired for better bioactivity of the surface. Trace amounts of Ca, K, Fe are observed, which most likely were impurities in the distilled water used.

Results of topography tests are presented in Figure 3. The surface of the top oxide layer is rough and well developed. The roughness profile is 630 nm, the Sa average value is 1.39 ± 0.79 μm on the tip of the helix, and 5.69 ± 2.98 μm on the bottom of the helix (Table 3). Such values in the nanometric range are also useful. The differentiation of the area comprising small nanotubes and rough pores is important. Surface modification led to smoothing the tip of the helix as a result of dissolving roughness peaks. The influence of roughness on oseointegration has been proven. In the case of long-term implants, a positive osteoblast response is required. With increasing roughness, the possibility of osteoblasts settling increases [45]. High roughness also carries the risk of biofilm formation [46]. The topography results confirmed the obtaining of a surface with a high surface roughness value. The lowest values were obtained for the oxidized sample.


**Table 2.** The EDS examination results of the chemical composition of the oxide layer.

**Figure 3.** The topography of reference Ti-13Zr-13Nb before (**a**) and after (**b**) sand blasting and oxidized Ti-13Zr-13Nb surfaces (**c**) obtained by light microscopy; the bottom of the helix (**left**) and top (**right**).

**Table 3.** Sa roughness parameters results.


The CT investigations showed the appearance of modifications at the bottoms of the helix and not at the tops (Figure 4). The area of modification can be observed as grey and red areas at the bottoms and base alloy as white metal.

**Figure 4.** CT images: (**a**) horizontal cross-section with regard to y axis, (**b**,**c**) vertical cross-sections, (**d**) 3D sample model.

The corrosion results are presented in Figure 5 and Table 4. The creation of the oxide layer became difficult because of much higher current values and a presumed shift of electrochemical potential into more anodic values resulting in the dissolution of metal rather than the oxidation of the surface. The local appearance of the nanotubular and highly rough surface is evidence that some microcells are formed due to change in potential. The local changes in pH value influence the anodization rate, the thickness of the oxide layer and its structure, or even its formation. In case of too low or too high acidity, the oxide layer is unable to achieve the nanotubular structure [47]. Here the anodization was made at the proper HF concentration enhancing the stabilization of the appropriate low pH value and resulting in a short oxidation time, thin nanotubes, a short distance between them, and scarce surface cracks. The roughness of the oxidized surface was close to that observed in similar experiments [48].

The open-circuit potential (OCP) of the non-oxidized specimen was about −199 mV(SCE). The anodic polarization exhibits a narrow plateau between 300 and 1150 mV, which can be attributed to the presence of natural titanium oxide on a specimen surface. The passive current value in this area ranged between 200 and 300 μA. For the previously oxidized specimen, the OCP was about −616 mV. The anodic curve shows a very stable passive region between −200 and 2000 mV. The passive current was about 200 μA in the entire region. However, despite high passive regions, the increase in corrosion current density after oxidation shows that the surface has not been uniformly covered with oxide layers and many microcells could appear in these oxidation conditions. The titanium dioxide layer formed on the surface of the titanium can provide increased corrosion resistance only if it is continuous over the entire surface of the alloy. The layer presented in the paper is characterized by cracks and a lack of continuity. This results in the formation of so-called "corrosion channels", which accelerate the degradation of the material. A similar effect was obtained in research [49]. The occurrence of this phenomenon may explain the deterioration of the corrosive properties compared to the reference sample.

The microscopic investigation reference specimens after corrosion tests showed effects of pitting, some discontinuity of material, and a heterogeneous structure at the bottoms of the helix and at the

tops (Figure 6). The microscopic investigation specimens oxidized in an electrochemical way and showed a network of cracks in the surface of the helix top and corrosion pitting in the bottom of the helix (Figure 7), it is probably related to the grater thickness of the obtained modification, which was confirmed by CT tests—Figure 4b.

**Figure 5.** Potentiodynamic polarization curve of reference and oxidized Ti-13Zr-13Nb specimens.

**Table 4.** Corrosion properties of reference and oxidized Ti-13Zr-13Nb specimens.


**Figure 6.** Surface of reference specimen after corrosion test at different magnifications: (**a**) surface of helix top ×200, (**b**) surface of helix top ×1000, (**c**) bottom of the helix ×200, (**d**) bottom of the helix ×1000.

**Figure 7.** Surfaces of specimens oxidized in an electrochemical way after corrosion tests: (**a**) surface of helix top ×200, (**b**) surface of helix top ×1000, (**c**) bottom of the helix ×200, (**d**) bottom of the helix ×1000.

#### **4. Conclusions**

In summary, nanotubular oxidation on the helix lines of titanium dental implants is possible, but it depends heavily on the geometric shape of the implant, anodization parameters, and environment composition. The parameters proposed here make it possible to obtain the nanotubes on the bottom of the helix and distinctly roughen almost all remaining surfaces. However, the applied conditions applied indicate that future investigations be oriented towards oxidation of the whole surface by introducing slightly higher HF contents and lower current values and mixing the electrolyte bath.

**Author Contributions:** Conceptualization, M.J.; methodology, M.J. and M.B.; formal analysis, M.J. and M.B.; investigation, M.J. and M.B.; writing—original draft preparation, M.J.; writing—review and editing, M.J. and M.B.; supervision, M.J. All authors have read and agreed to the published version of the manuscript.

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

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

**Informed Consent Statement:** Not applicable.

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

**Acknowledgments:** Authors thank the students—M. Get, M. Karczmarczyk, and K. G ˛asiorowska for their technical assistance in some tests. The helpful comments of Andrzej Zielinski are gratefully acknowledged.

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

#### **References**


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