Electrochemical Corrosion Behavior of Nanocrystalline Î²-Ta Coating

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Electrochemical corrosion behavior of nanocrystalline #-Ta coating for biomedical applications Linlin Liu, Jiang Xu, Xiaolin Lu, Paul R. Munroe, and Zonghan Xie ACS Biomater. Sci. Eng., Just Accepted Manuscript • DOI: 10.1021/ acsbiomaterials.5b00552 • Publication Date (Web): 25 Feb 2016 Downloaded from http://pubs.acs.org on February 28, 2016

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ACS Biomaterials Science & Engineering

Electrochemical corrosion behavior of nanocrystalline β-Ta coating for biomedical applications Lin lin Liua, a.

Jiang Xua,b*,

Xiaolin Lua, Paul Munroec,

Zong-Han Xieb,d

Department of Material Science and Engineering, Nanjing University of Aeronautics and Astronautics, 29 Yudao Street, Nanjing 210016, PR China

b. School of Mechanical & Electrical Engineering, Wuhan Institute of Technology,693 Xiongchu Avenue, Wuhan 430073, P. R. China c. School of Materials Science and Engineering, University of New South Wales, NSW 2052, Australia d. School of Mechanical Engineering, University of Adelaide, SA 5005, Australia

Abstract To explore its potential as a highly corrosion-resistant coating for biomedical titanium alloys, a novel β-Ta nanocrystalline coating, composed of equiaxed β-Ta grains with an average grain size ~22 nm, was deposited onto Ti–6Al–4V substrate using a double glow discharge plasma technique. The newly developed coating exhibited an extremely dense and homogeneous microstructure, exhibiting a strong (002) preferred orientation. The electrochemical behavior and semiconducting properties, such as donor density, flat-band potential, and diffusivity of point defects (Do), of the passive film formed on the β-Ta coating were compared to those for both uncoated Ti–6Al–4V and commercially pure Ta in Ringer’s physiological solution at 37 °C, using an array of complementary electrochemical techniques. The results showed that the β-Ta coating not only provided Ti-6Al-4V with good corrosion protection, but also endowed a higher resistance to corrosive attack than commercially pure Ta in in Ringer’s physiological solution at 37 °C. Mott-Schottky analysis revealed that the passive film formed on the β-Ta coating had a lower donor density and the flat-band potential than commercially pure Ta. The calculated values of Do for the β-Ta coating (1.45×10-16cm2/s) are comparable to that for commercially pure Ta (1.13×10-16 cm2/s), both of which are one order of magnitude lower than that for uncoated Ti-6Al-4V (2.73×10-15 cm2/s). Compared with commercially pure Ta, the higher corrosion resistance of the β-Ta coating results from its nanosized grains and crystallographic orientation, 1

*Corresponding authors. ACS Paragon(J. Plus E-mail addresses: [email protected] Xu)Environment


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which promotes the rapid formation of a robust passive film with a more compact structure. Key words: β-Ta; nanocrystalline; preferred orientation; Mott-Schottky; PDM

1. Introduction Among the metallic biomaterials, titanium and its alloys have emerged as one of the most attractive metals for applications in orthopedic implants, dental implants and medical devices, due to the combination of their outstanding characteristics such as low density, high specific strength, high corrosion resistance and good compatibility with biological materials [1, 2]. Of the various types of titanium alloys available for implant applications, the α+β alloy Ti–6Al–4V was the first Ti alloy registered as an implant material in ASTM standards and remains the most frequently used titanium alloy for metallic prostheses [3]. In general, excellent corrosion resistance in body fluids and good biocompatibility are the key prerequisites for biomedical applications, and the latter is strongly dependent on the former. The success of titanium alloys as an implant materials derives from their good corrosion resistance in several media through the spontaneous formation of a protective oxide film with a thin thickness (typically 3–10 nm), consisting mainly of amorphous titanium dioxide (TiO2) [1]. In spite of these alloys showing an excellent reputation for corrosion resistance in different environments, their long-term performance has raised some concerns regarding the interaction between them and the biological environment. Human body fluid is extremely aggressive chloride containing electrolyte, typically having a concentration of ~ 1 wt.% NaCl. When implants are exposed to human fluids containing high chloride ion concentrations, the oxide film may become unstable, leading to the occurrence of localized breakdown on a highly microscopic scale [4]. Subsequently, the degradation of titanium alloys exposes the underlying bare metal to corrosive medium, which result in dissolution of the base alloy. Moreover, scratches, dents, and fretting might also lead to rupture or weakening of the passive oxide films, resulting from mechanically accelerated electrochemical processes such as such as stress corrosion, corrosion fatigue and fretting corrosion [5]. In the case of Ti–6Al–4V, the dissolved non-compatible component metallic ions (e.g., Al and V) released into surrounding tissue of human body, might not only cause implant failure, but also threaten long-term health, due to a negative immunological response of organism. In vivo animal testing has indicted that an elevated level of aluminum was observed in the peripheral tissues surrounding Ti-6Al-4V implant 2

ACS Paragon Plus Environment

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[6], and the excess accumulation of Al ions in the tissue could cause neurological disorders and senile dementia of the Alzheimer type [7, 8]. Likewise, vanadium both in the elemental state and in the form of oxides, exhibits significant toxicity and adverse reactions with the human body evaluated [9]. Furthermore, owing to the inadequate wear resistance of Ti-6Al-4V under repeated loading conditions, especially for joint replacement prostheses, the generated wear debris caused by mechanical friction induces a significantly higher release of proinflammatory and osteolytic mediators, which ultimately result in implant loosening and periprosthetic soft tissue reactions [10, 11]. In view of the above, the surface properties of Ti-6Al-4V implant play a significant role in wear and corrosion processes, which are the key to suppressing the potential metal release and the formation of wear debris. Therefore, considerable efforts have been made to modify the surface properties of titanium alloys by different surface treatment techniques to prolong service life [12-14]. Tantalum is a semi-precious refractory metal with excellent physical and chemical properties, including a high melting point (Tm = 2996 °C), low ductile-to-brittle transition temperature, good ductility and biocompatibility and has extensively used in many different domains, such as electronic, chemical, mechanical and biomedical applications [15]. Similar to titanium, a highly stable and tenaciously adherent Ta2O5 oxide film that naturally grows on a Ta surface provides extreme chemical inertness, providing Ta not only with better corrosion resistance than Ti alloys in biological environment, but also exceptional biocompatibility, osseointegration properties, hemocompatibility and high radiopacity [16]. Stents made of tantalum with excellent radio-opacity property have already been applied to endovascular surgery and may be a promising alternative to the low-temperature isotropic pyrolytic carbon (LTIC) that is currently the most widely accepted biomedical material for artificial heart valves [17]. Recently, porous tantalum with a characteristic appearance similar to cancellous bone, has shown strong bone-bonding properties and has been currently used in a wide array of orthopedic applications (e.g. hip and knee arthroplasty, spine surgery, and bone graft substitute) [18, 19]. One major limitation of these porous tantalum implants is their significantly lower fatigue strength in comparison with their equivalent fully dense materials. However, the high density and cost of manufacture are the main obstacles to metallic Ta gaining widespread acceptance of bulk structural Ta metal for implant applications. A great deal of attention has been focused on the fabrication of protective Ta coatings 3



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by combining the favorable bulk properties of the substrate with desired surface properties to perform the function for which it was originally designed. Although previous studies has shown that Ta coatings on Ti and porous Ti-6Al-4V prepared via different surface modification methods had better biocompatibility and osseointegration properties compared with uncoated substrates [20, 21], little work, to our knowledge, has been attempted to investigate the electrochemical behavior of Ta coatings on titanium alloys in simulated physiological media. Essentially, the microstructural characteristics of as-deposited coatings are an important factor in controlling their corrosion resistance. Koivuluoto et al. [22] found that the denseness of Ta coatings on carbon steel substrates prepared by high-pressure cold spraying was closely related to the corrosion resistance of these coatings, and moreover, corrosion resistance of a dense Ta coating was comparable to that of the bulk Ta foil. In practice, it is generally not feasible to achieve such dense coatings by many deposition techniques, because the presence of defects or pinholes in the coatings cannot be completely eliminated. The presence of channel defects connecting the substrate and environment through the entire film thickness, leads to severe localized corrosion due to the galvanic action when the coating is cathodic with respect to the substrate [23, 24]. Furthermore, nanocrystalline structures in metals or alloys have been reported to impart significantly higher corrosion resistance than conventional coarse grain counterparts of the same chemical composition. For example, Li et al. [25] compared the electrochemical behavior of coarse-grained (CG) and nanostructured (NS) Ti2448 alloys in 0.9% NaCl solution at 37 °C. The corrosion resistance of the NS alloy was superior to that of the CG alloy, which was attributed to the fact that the nanostructuring can improve the stability of passive thin films on biomaterials. Our recent studies have demonstrated that double cathode glow discharge plasma technique is a simple, yet effective coating method to enhance the surface properties of Ti-6Al-4V alloy. The metal silicide or metal nitride coatings prepared using this technique exhibited a dense, nanocrystalline structure, with a strong adhesion with the substrate [26, 27]. In the present study, a double cathode glow discharge technique was employed to fabricate a novel nanocrystalline β-Ta coating onto a Ti-6Al-4V substrate, with the aim of obtaining a more adherent, dense and corrosion resistant as-deposited coating. The microstructure of the coating was characterized by X-ray diffraction (XRD), scanning electron microscopy (SEM) and transmission electron 4


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microscopy (TEM). In view of its biomedical applications, the electrochemical behavior of the coating was investigated in Ringer’s physiological solution at 37 °C, using an array of complementary

electrochemical

techniques

including

open

circuit

potential

(OCP),

potentiodynamic polarization, potentiostatic polarization, EIS, Mott-Schottky analysis and the point defect model. For comparative purposes, the same measurements were also performed on commercially pure Ta and uncoated Ti–6Al–4V.

2. Materials and methods 2.1. Coating specimen preparation Disk-shaped substrates, 40 mm in diameter and 3 mm in thickness, were machined from a medical grade Ti-6Al-4V alloy bar (Baoji Titanium Industry Co., Ltd, China). The nominal composition of this alloy in wt.% is given as: Al, 6.42; V, 4.19; Fe, 0.198; O, 0.101; C, 0.011; N, 0.006 and Ti, the balance. The Ti–6Al–4V substrates were ground using with SiC abrasives of sequentially finer (240–2000) grit, followed by polishing with diamond paste. Subsequently, the polished substrates were ultrasonically cleaned in acetone, alcohol and distilled water, successively, and then dried in cold air. Deposition was performed using a double cathode glow discharge apparatus with a 100 mm diameter Ta target (99.99% pure). During the deposition process, one cathode is used as the target, and the other cathode is the substrate material. When two different voltages are applied to the two cathodes, glow discharge occurs, as described in detail elsewhere [28].Temperature of the substrate was monitored by a thermocouple and mass-flow controllers regulated the flow of argon (99.999% pure) to the chamber. Prior to deposition, the substrates were sputter etched with Ar at a pressure of 20 Pa and a DC bias of −650 V for 10 min to remove any residual surface contaminants. The detailed processing parameters are described as follows: base pressure, 5 × 10-3 Pa; target electrode bias voltage with direct current, -850 V; substrate bias voltage with impulse current, -250 V; substrate temperature, 800 °C; target substrate distance, 10 mm and treatment time 2 h.

2.2 Microstructural characterization The crystalline structure of the as-deposited coating was characterized by X-ray diffraction (XRD, Bruker AXS Inc., Germany) using a D8 ADVANCE diffractometer with Cu Kα irradiation 5



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(λ = 0.154060 nm) operating at 35 kV and 40 mA. X-ray spectra were collected over scanning angles ranging from 20 to 80° with a step rate of 0.05° s-1. The surface and cross-sectional morphology as well as elemental analysis were determined using a field emission scanning electron microscopy (FE-SEM; Hitachi, S-4800, Japan) equipped with an energy dispersive spectrometer (EDS). To reveal microstructural features, the polished cross-section of samples were etched with a 10 ml HNO3, 4 ml HF and 86 ml H2O) for 15-25 s. The details of the microstructure, including grain size and its distribution, were examined using a transmission electron microscope (TEM, JEM-2100, JEOL, Japan) operating at an accelerating voltage of 200 kV. The plan-view samples for TEM observation were prepared following the standard procedure of slicing, mechanical grinding, dimpling and finally jet electropolishing from the untreated side of the substrate. Cross-sectional TEM specimens were prepared using a focused ion beam (FIB) microscope (FEI xP200, FEI Company, Hillsboro, OR). X-ray photoelectron spectroscopy (XPS) measurements were carried out on a Kratos AXIS Ultra DLD spectrometer using a monochromatic Al Kα (hν = 1486.6 eV) X-ray source with an emission current of 15 mA and a voltage of 13 kV. The Shirley background, Gaussian and mixed Gaussian/Lorentzian functions, and least-square routine were used for the peak fitting procedure. All spectral positions were corrected for possible charging shifts by referencing the C 1s peak from adventitious carbon at 284.8 eV.

2.3. Electrochemical measurements Electrochemical experiments carried out using a CHI660C electrochemical workstation (Shanghai ChenHua Instruments Inc., China) using a standard three-electrode configuration, which includes a test specimen with an exposed area of 1.0 cm2 as the working electrode, a saturated calomel electrode (SCE) as reference, and a large area platinum sheet as a counter electrode. The working electrolyte was a naturally aerated aqueous Ringer’s physiological solution (NaCl 8.61 g l-1, CaCl2 0.49 g l-1, KCl 0.30 g l-1). The temperature of the aqueous solution was controlled to be at 37±0.5°C (i.e., normal body temperature). Throughout this paper, all potentials were referred to the SCE. Prior to the electrochemical experiment, working electrode was pretreated cathodically at -0.80 V for 10 min to create reproducible initial conditions. Monitoring of the open circuit potential (EOCP) evolution was made during 60 min immersion. Afterwards, the electrochemical impedance spectroscopy (EIS) measurements were conducted over a frequency 6


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range from 100 kHz to 10 mHz with a 10 mV (peak to peak) amplitude signal at the respective EOCP. The EIS experimental data were analyzed and fitted to appropriate electrical equivalent circuit (EEC) using a ZSimpWin software package. The potentiodynamic polarization curves were recorded from -0.30 V vs. EOCP to 1.50 V at a scan rate of 20 mV min-1. Before undertaking capacitance measurements for Mott-Schottky analysis, the specimens were potentiostatically polarized at a sequence of appointed potentials for 60 min to form a steady-state passive film. Subsequently, the capacitance of passive film covered specimen was measured by sweeping the potential in the negative direction from the film formation value with potential steps of 25 mV, while simultaneously imposing a sinusoidal ac perturbation of 10 mV amplitude at a frequency of 1000 Hz. All of the reported results were verified by repeating the experiments, where only minor changes in the reported values were noted. Commercially pure Ta (99.99% pure) and bare Ti–6Al– 4V were used as references for electrochemical tests. A SEM image showing the microstructure of the commercially pure Ta specimen is shown in Fig.1. It is clear that the commercially pure Ta exhibits a fully annealed equiaxed grain structure and its grain size measured using the linear intercept method, was 300±40 μm.

3. Results 3.1. Microstructure and phase analysis Fig.2 shows typical XRD patterns recorded from the as-deposited Ta coating grown on Ti–6Al– 4V and for commercially pure Ta. As shown in Fig.2, the XRD peaks from commercially pure Ta are very similar to those of Ta powder and can be indexed as α-Ta, with a body-centered cubic structure, according to the relevant JCPDS card (No.04-0788). The XRD spectrum also shows a preferred (110) orientation. For the as-deposited Ta coating, three distinct peaks located at 2θ= 33.7°, 37.9° and 70.8° are, respectively, identified as the (002), (202) and (004) reflections for tetragonal β-Ta (JCPDS 25-1280), indicating that β-Ta is the only phase present in the as-deposited Ta coating. The diffraction peak intensity of the (002) reflection is significantly stronger than that of the corresponding theoretical powder diffraction pattern, suggesting that this coating shows a strong (002) preferred orientation. According to the intensity data, texture coefficients (TChkl) of both commercially pure Ta and the as-deposited Ta coating were calculated through the following equation [29]: 7



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𝑇𝐶𝑕𝑘𝑙 =

𝐼𝑚 (ℎ𝑘𝑙)/𝐼0 (ℎ𝑘𝑙) 1/𝑛 ∑𝑛1 𝐼𝑚 (ℎ𝑘𝑙)/𝐼0 (ℎ𝑘𝑙)

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(1)

where Im(hkl) is the measured X-ray relative intensity of the (hkl) plane, I0(hkl) is the relative intensity in the powder pattern, (hkl) represents the indices of the reflection plane and n is the number of reflection planes. For a TChkl value greater than 1, a preferred orientation is developed towards a specific crystalline plane; while a TChkl value close to 1 signifies a more random orientation and a TChkl value in the range from 0 to 1 denotes a lack of grain orientation for the specific plane under consideration [30]. A slight (110) texture for the commercially pure Ta and a strong (002) texture for the as-deposited Ta coating are further confirmed by the calculated results shown in Table 1. Such a (002) texture is often observed on β-Ta coatings prepared by the other sputter deposition techniques [31, 32]. Furthermore, the full width at half-maximum (FWHM) of the diffraction peaks in as-deposited Ta coating are relatively wide and, thus, indicative of very fine grains. After the correction of instrumental line broadening, the average grain size of the coating was estimated to be 18 nm by applying the Scherrer formula for peak broadening of the low angle (002) reflection [33]. Note that β-Ta is a thermodynamically metastable phase, which is frequently synthetized under non-equilibrium conditions such as sputter deposition or electro-deposition and can further transform into α-Ta through annealing [34]. As shown in Fig. 3(a), the surface of the Ta coating shows a smooth morphology and uniform coverage of the Ti-6Al-4V substrate without any obvious defects, such as pinholes, macro-droplets or micro-cracks. From the cross-sectional image shown in Fig. 3(b), the Ta coating exhibits an extremely dense and homogeneous microstructure with a thickness of ~40 μm and adheres to the Ti-6Al-4V substrate very well. Such a microstructural feature is intimately associated with the deposition conditions employed in the present study, including the high deposition temperature and substrate bias, which contributes to an increase of adatom surface mobility and bulk diffusion rates [35]. The deposited atoms have sufficient energy to overcome the self-shadowing effect exerted by the previously deposited atoms, thereby resulting in coating densification and surface smoothing. Moreover, coating porosity is in general inversely proportional to coating thickness. The β-Ta coating is so thick that it is more likely to obtain a defect-free coating due to the void masking effect, i.e., covering the voids (closing pores) by the depositing material [24]. It is clear from the EDS elemental maps that the Ta coating consists of only Ta, except that a small amount 8


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of Ti is present in 4μm thick diffusion layer, presumably originating from interdiffusion of Ta and Ti at the interface between the coating and substrate. This result indicates the presence of metallurgical bonding at the coating/substrate interface, which is beneficial to improving the adhesion strength between the coating and substrate. To obtain the phase structure, grain size distribution and the nature of the coating/substrate interface, a detailed investigation on the microstructure of the β-Ta coating was performed by combining the longitudinal and transverse cross-sectional and planar sections through TEM analyses, as shown in Fig. 4. As shown in Fig. 4 (a) and (b), the plan-view TEM examination suggests that the microstructure of the coating is characterized by uniformly distributed nanoscale grains with equiaxed shapes and a grain size, ranging from ~11 to ~32 nm in diameter, with an average grain size ~22 nm (shown in Fig. 4(c)) . This average grain size is slightly larger than the values deduced from the Scherrer equation based on XRD data, which is attributed to the presence of structural defects (i.e. stacking faults or low-angle grain boundaries) and the disordered character of grain boundaries [36, 37]. The selected area electron diffraction (inset in Fig. 4(a)) shows a ring pattern consistent with the (002) and (004) lattice planes of tetragonal β-Ta, and the strong diffracted intensity of the (002) ring provides further evidence of the β-Ta coating with a (002) texture, in agreement with the XRD data. A high resolution TEM (HRTEM) image of the Ta coating is shown in Fig. 4(d). Fast Fourier transform (FFT) analysis is used to accurately determine the interplanar spacing. Nanocrystallites with lattice-fringe contrast are clearly observed, and the spacing of the lattice fringes in two selected square can be calculated as 0.262 and 0.237 nm, corresponding to the (002) and (202) crystallographic planes of the tetragonal β-Ta phase. Fig. 5 shows typical bright-field cross-sectional TEM images taken from the intermediate layer and coating/substrate interface marked as squares A and B in Fig. 3(b). As can be observed in Fig. 5 (a), the intermediate layer has an elongated the columnar structure with a width of about 20~30 nm, and the growth direction of the columnar structure is mostly normal to the coating surface, as indicated by a white arrow. The columnar structure consists of equiaxed grains in its cross-section from plan-view TEM observation. From the interfaces of Ta deposition layer/diffusion layer and diffusion layer/substrate in Fig.5(b), a smooth transition of grain sizes from nano-scale to micro-scale is observed without the presence of structural defects at each interface.

9



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3.2. Electrochemical measurements The EOCP–time curves can reflect the variation of the nature of the metal–solution interface with increasing immersion time, and EOCP can be used as a criterion for corrosion behavior. Fig. 6 presents the evolution of the open circuit potentials (EOCP) for the β-Ta coating as a function of exposure time in Ringer’s physiological solution at 37 °C. For the sake of comparison, experiments were also conducted on uncoated Ti–6Al–4V and commercially pure Ta under the same condition. As can be seen, all the EOCP vs. time curves are similar in shape, viz., the EOCP shifts rapidly towards more positive potential at initial stages of immersion, and subsequently, varies slowly towards a quasi-stationary value. A similar behavior was also observed for Ti–Mo [38] and Ti–Ta alloys [39] in Ringer’s physiological solution. An initial increase in the EOCP is attributable to the spontaneous protective oxide films formed on the tested samples in this medium. These spontaneously formed oxides are electrochemical stable at quasi-stationary values of the EOCP, as sustained by Ti–H2O and Ta–H2O Pourbaix diagrams, and can act as barrier layers to hinder metal ion release. As shown in Fig.6, the EOCP vs. time curves can be fitted to an exponential decay function with two time constants, which is associated with two distinguishable stages on the specimen surfaces during immersion. In terms of the fitted results, it is evident that the first time constant is within the range of 1 to 3 min, associating with the formation and growth of oxide films, and the second time constant increases by almost one order of magnitude ranging from about 16 to 26 min, in relation to thickening of the oxide films. The first time constant for the β-Ta coating (1.34 min) is significantly shorter than those for commercially pure Ta (2.91 min) and uncoated Ti-6Al-4V (2.60 min), indicating that the β-Ta coating exhibits a higher growth rate of oxide film as compared to the two coarse-grained reference samples. This issue will be discussed in detail in later. By comparing the quasi-stationary values of the EOCP after 60 min immersion, it can be observed that the β-Ta coating has the most positive values, representing its oxide film with a higher thermodynamic stability. Fig.7 shows representative impedance spectra of the β-Ta coating, bare Ti–6Al–4V and commercially pure Ta measured at their representative EOCP in Ringer’s physiological solution at 37 °C. As seen in Fig. 7(a), the Nyquist plots for the three tested samples display similar features, characterized by somewhat unfinished capacitive semicircles with different diameters and Z″/Zʹ ratios. The β-Ta coating exhibits a much larger Z″/Zʹ ratio and diameter of the capacitive 10


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semicircle than the two reference materials, suggesting that it has a higher electrochemical corrosion resistance with an enhanced capacitive behaviour for the solid/ liquid interface. As shown in Fig.7(b), the shape of the Bode plots for the three tested samples is also very similar. The impedance magnitude is almost independent of frequency at high frequency (105–103Hz), indicative of a resistive response, and in the frequency range of 103–10−2 Hz, the log |Z| vs. log f plots exhibit a single linear region with a slope close to -1, which corresponds to plateaus of phase angles close to –90° in the Bode-phase plots. The β-Ta coating shows the largest impedance modulus at the low-frequency limit (|Z|f→0) and phase angle maximum, as well as the widest phase angle plateau among the three tested samples, denoting that the passive film formed on the β-Ta coating is more protective than those for the uncoated Ti–6Al–4V and commercially pure Ta in the Ringer’s physiological solution. After inspection of the impedance spectra in Fig.7, only one time constant was observed and thus the equivalent circuit model (Rs(QpRp)) shown in Fig.8 was applied to simulate the experimental impedance data obtained from the tested specimens. The selection of the circuit was based on the principle of obtaining a reasonable fitting to the experimental data using minimum components in the equivalent circuit. The physical significance of the circuit elements is as follows: Rs represents the solution resistance, Rp is the resistance of the barrier layer and charge transfer, and Qp is a constant phase element related to the capacitance of the barrier layer and the electric double layer. Because the barrier layer never exhibits the theoretically expected phase shift of –90° and a slope of -1 for an ideal dielectric, the CPE is used to account for the non-ideal behaviour of the capacitive elements, due to the distribution of relaxation times as a consequence of heterogeneities on the electrode surface [40]. The impedance, ZCPE, of CPE is defined by ZCPE = [Q(jω)n]-1, where Q is the frequency independent parameter (Ω-1 cm-2 sn), j is the imaginary number, ω is the angular frequency (rad s-1), and the factor n, defined as a CPE power, is an adjustable parameters related to the dispersive behavior with values between 0 and 1. The fitting parameters of the circuital elements are given in Table 2. As illustrated in Fig. 7, a better agreement between the experimental (dots) and simulated (full lines) data was obtained and the chi-squared (χ2) values are of the order of 10-4, indicating that the proposed model is suitable for explaining the behavior of the tested samples in Ringer’s physiological solution. The parameter Rp is directly proportional to the corrosion resistance of the tested specimens [41]. As shown in Table 11



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2, the Rp of the β-Ta coating (2.78 × 106 Ω cm2) is more than twice that of commercially pure Ta (1.21×106 Ω cm2), both of which are an order of magnitude higher than that of the uncoated Ti-6Al-4V (5.75×105 Ω cm2). Unlike the resistance (R) values, the capacitance (C) values are independent of the solution used and can provide information about the dielectric property of the passive film [42]. The effective capacitance (Cp) is calculated from Qp based on the formula developed by Brug et al. [43]

𝐶𝑝 = 𝑄𝑝1/𝑛 (𝑅𝑠 −1 + 𝑅𝑝 −1 )(𝑛−1)/𝑛

(2)

The values of Cp for the β-Ta coating, uncoated Ti-6Al-4V and commercially pure Ta are 5.10, 8.43 and 6.19 μF cm-2, respectively. The β-Ta coating has a lower Cp value, implying that the passive film formed on the β-Ta coating shows a higher insulating or dielectric property than those on the uncoated Ti-6Al-4V and commercially pure Ta. The parameters R and C are area dependent and cannot be directly used to evaluate corrosion rate of the tested samples, whereas the values of the two parameters can be multiplied to obtain the product C × R equivalent with the time-constant (τ) of the charge transfer reaction [44, 45]. The reciprocal of the obtained product values can describe a rate of charge transfer reactions equivalent with the corrosion rate of the tested samples. From Table 2, τ increases from 4.85 s for uncoated Ti-6Al-4V to 7.49 s for commercially pure Ta and further to 14.18 s for the β-Ta coating. This result shows that the passive film formed on the β-Ta coating has a lower rate of charge transfer as compared with the two reference samples. Fig.9 (a) compares typical potentiodynamic polarization curves of the β-Ta coating, uncoated Ti-6Al-4V and commercially pure Ta in Ringer’s physiological solution at 37 °C, and the values of the related electrochemical parameters determined from the polarization curves are summarized in Table 3. As shown in Fig.9 (a), all the tested specimens are spontaneously passivated in this electrolyte, with a distinctive passive plateau over wide potential range in the anodic domain. Obviously, the β-Ta coating is characterized by the noblest corrosion potential (Ecorr), the lowest corrosion current density (icorr) and passive current density (ipass), indicating that the coating exhibits the highest corrosion resistance among the three tested specimens. This is in agreement with the EIS results shown in Fig.7. In contrast to the other two samples that maintain the passivation until the upper limit of applied potential (1.5 V), commercially pure Ta shows an 12


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apparent increase in current at an applied potential above 1.2 V. A similar phenomenon has been previously observed on Ta electrodes with different surface treatments by Silva et al. [46], who attributed it to the localized dissolution of Ta due to the breakdown of the passive film. From Table 3, the cathodic Tafel slopes (-βc) of the tested specimens are in the range of 112~118 mV/decade, close to the theoretical value of 120 mV/decade for the hydrogen evolution reaction involving a rate-limiting one-electron transfer step [47]. However, the anodic Tafel slopes (βa) for the tested samples are larger than their respective cathodic counterparts (-βc), suggesting that the electrochemical corrosion process of the tested samples is under anodic control [48]. As illustrated in Fig. 9(b), the nobler Ecorr and lower icorr observed for the β-Ta coating are derived from its much larger βa value compared to uncoated Ti-6Al-4V.

3.3 Passive film investigations As one key factor that influences the protective ability of the passive film, the compact property of the passive films on the tested samples was evaluated by potentiostatic polarization measurements. According to the literature [49], if the contribution of the double layer charge is negligible, the initial decrease in current density can be related to the growth of a passive film on the electrode surface. In order to assess the compact property of the passive film, the variation of the current density with immersion time for each specimen after cathodic reduction was measured at a constant potential (0.8 V). The current decreases with time following formula [50]:

𝑖 = 10−(𝐴+𝑘𝑙𝑔𝑡)

(3)

where i represents the current density, t is time, A is constant, and k represents the slope of the double-log plot for potentiostatic polarization, reflecting the compact property of the passive film; k = -1 signifies the formation of a compact, highly protective passive film, while k = -0.5 indicates the presence of a porous and diffusion-controlled film. Fig.10 shows the double-log plots of current-time for the β-Ta coating, uncoated Ti-6Al-4V and commercially pure Ta potentiostatically polarized at 0.8 V in Ringer’s physiological solution. The values of k for the passive films grown on the β-Ta coating, uncoated Ti-6Al-4V and commercially pure Ta are -0.98, -0.88 and -0.96, respectively. The results show that the compact property of the passive film on the β-Ta coating is slightly superior to that of commercially pure Ta, but both values are markedly higher than that for uncoated Ti-6Al-4V. 13



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In order to gain insight into the relevant information about the formed passive films in passive region, impedance measurements were performed on the three tested specimens after potentiostatic polarization at 0.8 V for 1 h in Ringer’s physiological solution, and the measured impedance spectra are shown in Fig.11. In comparison with the EIS spectra shown in Fig.7, negligible change in impedance behavior was detectable for the β-Ta coating, whereas both uncoated Ti-6Al-4V and commercially pure Ta display a more depressed capacitive loop in the Nyquist plot, and narrower phase plateaus at intermediate frequencies and a smaller low-frequency limit for the impedance modulus |Z|, indicating a less capacitive response and a reduction in corrosion resistance for the two reference samples. The impedance data of the tested samples were also fitted with the equivalent circuit model shown in Fig.8 and the values of the fitted equivalent circuit parameters are presented in Table 4. It is obvious that the β-Ta coating shows the largest value of Rp (2.47 × 106 Ω cm2), followed by uncoated Ti-6Al-4V (6.08 × 105 Ω cm2) and commercially pure Ta has the smallest value of Rp (1.43 × 105 Ω cm2). Considering the passive film as a parallel plate capacitor, an approximate estimate for the passive film thickness (dp) can be calculated by Eq.(4) [51]: 𝑑𝑝 =

𝜀0 𝜀𝑟 A 𝐶𝑝

(4)

where Cp is the capacitance of the passive film converted from the constant phase element (Qp), εr is the vacuum permittivity (8.854 × 10-14 F cm-1), εr represents the dielectric constant of the passive film which is taken as 60 [52] for the passive film formed on the uncoated Ti-6Al-4V and 25 [53] for that of the β-Ta coating and commercially pure Ta, and A is the sample surface area. The calculated thicknesses of the passive films formed on the β-Ta coating, uncoated Ti-6Al-4V and commercially pure Ta are 15.36, 23.39 and 3.80 nm, respectively. From above results, the formed passive film of commercially pure Ta was significantly thinner than that of the β-Ta coating and the uncoated Ti-6Al-4V, which may adversely affect its protective ability. To unlock the composition and chemical state of the passive film, XPS analysis was carried out on the β-Ta coating after potentiostatic polarization at a potential of 0.8 V for 1 h in Ringer’s physiological solution at 37 °C. As shown in Fig.12(a), the wide scan spectrum of the β-Ta coating comprises of peaks for C, O and Ta. Fig.12 (b) and (c) show the high-resolution XPS spectra for Ta 4f and O 1s peaks collected from the passive film formed on the β-Ta coating, respectively. The 14


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Ta 4f core level spectrum (Fig. 12(b)) exhibits a doublet with Ta 4f7/2 and Ta 4f5/2 peaks at 26.4 and 28.3 eV, respectively, corresponding to stoichiometric Ta2O5 [54]. The O 1s core level spectrum (Fig. 12(c)) show a single broad peak, which can be deconvoluted into three overlapping peaks. A strong intensity peak located at 530.2 eV is attributed to the Ta−O bonds in tantalum oxide, and the two weak peaks located at 531.5 and 532.7 eV are assigned to oxygen in the form of the adsorbed OH-impurities on the sample surface from the ambient atmosphere and bound water, respectively [54]. Therefore, XPS analysis demonstrates that the passive film formed on the β-Ta coating consists mostly of Ta2O5. Fig.13 shows representative C-E curves for the passive films formed on the three tested specimens after an immersion time of 1 h at 0.8 V in Ringer’s physiological solution at 37 °C. Evidently, the C-E curves for the passive films on three tested samples bear a generally strong resemblance. At low potential region, the measured capacitance is inversely proportional to the applied potentials, and attains a stationary value, nearly independent of potential, when the applied potentials exceed certain values. Similar features of potential dependence of the capacitance were also observed for the electrochemically formed oxide films on Nb [55, 56]. The variation of capacitance with potential shown in Fig.13 can be interpreted with respect to the band structure model of n-type semiconducting passive film [57]. The decrease of the capacitance in the low potential region is attributable to the increase of the thickness of the depletion layer with increasing the band bending according to a typical Mott-Schottky behavior. As the applied potentials above certain values, the constant capacitance can be interpreted with the exhaustion of all donors in the passive film, which behaves like an insulator. Remarkably, as can be seen from the enlarged part of capacitance curve for commercially pure Ta, the capacitance increases with the applied potentials, when the applied potential is above 0.2 V. This phenomenon is connected with the formation of an inversion layer stemming from an increase of hole concentration in the valence band [57]. As is well-known, the formation and breakdown of the passive films is mainly regulated by the concentration and migration of point defects within the passive film, which are, in part, determined by the electronic properties of the passive films themselves. Mott–Schottky analysis, based on impedance measurements is a powerful in situ tool in probing the electronic properties of passive films. Because the charge distribution at the semiconducting passive film/electrolyte 15



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interface can be determined by Mott-Schottky analysis, the capacitance response is controlled by the band bending and can be described by the variation of the space-charge capacitance under depletion conditions [58, 59]. In this work, the space-charge capacitance (Csc) was assumed to be equal to the electrode capacitance (C) at sufficiently high frequencies. According to the Mott– Schottky theory, the space charge capacitance (Csc) of an n-type or a p-type semiconductor is given by Eq.(5): 1 Csc

2

=

2 εr ε0 qNq

(E − Efb −

kT q

)

(5)

where CSC is space charge capacitance, εr is the dielectric constant of the passive film (~60 [52] for TiO2 and ~25 [53] for Ta2O5), ε0 is the vacuum permittivity (8.854 × 10-14 F cm-1), q is the elementary charge (+e for electrons and –e for holes), Nq is the density of charge carriers (Nd for donors and Na for acceptors), E is the applied potential, Efb the flat band potential, k is the Boltzmann constant (1.38 × 10-23 J K-1), and T is the absolute temperature (310 K). Fig.14 shows the Mott-Schottky plots for the passive films formed on three tested samples after being polarized at different film formation potentials (Ef) for 1 h in Ringer’s physiological solution at 37 °C. The shape of the 1/Csc2 vs. E curves for the passive films formed on the three tested samples is very similar to that for the passive films formed on some valve metals, such as niobium [60, 61], zirconium [62] and tungsten [63]. In Fig.14, there is a break point that separates Mott–Schottky plots into two linear regions over the entire potential range and the slopes of two linear regions in each plot are always positive, indicating that the passive films formed on the three tested samples at different Ef are characteristic of n-type semiconductor. Moreover, it can be observed that the magnitudes of Csc2 and the slope of Csc2 vs. E increase with the Ef, indicating a decrease of the donor density in the passive films with the increase of the Ef. These two linear regions in the Mott–Schottky plots have also been observed for the other metal oxides in different media [62-65]. Some researchers have suggested that the two linear regions of Mott-Schottky plots in low and high potential were attributed to the semiconducting nature and dielectric behavior for passive films formed on zirconium [62], titanium [65] and tantalum [66], respectively. Our experimental results seem to support this view for the following reasons: (1) it can be seen from the C-E curves (Fig.13) that the constant capacitance have low values (of the order of 3~6 μF/cm2) in the high potential region, indicating that a full depletion situation is established and the 16


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passive film exhibits high dielectric (insulating) property; (2) the linear regions at high potential are parallel to each other for the tested samples at different Ef. This indicates that the donor density should be independent of Ef, which, however, is in conflict with the theoretical prediction of the point defect model (PDM) [67]. It is worth noting that there is a linear region with negative slope observed on commercially pure Ta at film formation potentials of 0.6, 0.8 and 1.0 V, as shown in Fig. 14(b). Generally, inversion behavior [68, 69], surface state [70] and potential-dependent film changes [71] have been proposed to explain this phenomenon. As discussed by Harrington and Devine [68], the above three physical mechanisms can be differentiated by their time dependencies. Surface states charge/discharge at finite rate and the change in film composition or structure is slow. In contrast, the formation of an inversion layer is a rapid electronic transition. The sweeping rate of 25 mV/step employed here in the capacitance measurements is sufficiently rapid to ensure that the composition, structure, and defect concentration of the passive film are “frozen-in”. Therefore, the abnormal linear regions appeared in the Mott-Schottky plots of commercially pure Ta are related to the development of an inversion layer, which is also shown by the C-E curves for commercially pure Ta. The calculated donor density (Nd) and flat band potential (Efb) for the passive films formed at different potentials for 1 h in Ringer’s physiological solution from Mott-Schottky plots, are shown in Table 5. It can be observed that with increasing Ef , Nd decreases and Efb slightly increases for the tested samples. At a given Ef, the passive film on commercially pure Ta has the largest value of Nd, of the order of 1020 cm-3, which is significantly greater than those for the passive films formed on both uncoated Ti-6Al-4V and the β-Ta coating. Since the values of the carrier density reflects the number of the defects in the space charge layer, the difference in Nd among the three tested specimens represents the variation in the amount of oxygen vacancies and/or interstitials in their passive films. The defects can act as the charge carrier and thus the higher the carrier density, the greater the passive film conductivity [72, 73]. According to the point defect model (PDM) [74], in chloride-containing solutions, negatively charged chloride ions can preferentially occupy positively charged oxygen vacancies at the passive film/solution interface. The processes of the chloride ion absorption is autocatalytic, leading to an increase in the electromigration-dominated flux of cation vacancies from the oxide/solution interface to the metal/oxide interface, where they are annihilated by an oxidative injection of cations from the metal into the passive film. If the 17



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annihilation reaction cannot accommodate the enhanced flux of cation vacancies, the excess vacancies will condense to form voids, thereby resulting in thinning and subsequent breakdown of the local passive film. Due to the presence of a concentration of chloride ions in Ringer’s physiological solution, the passive film with a higher donor density can induce a larger cation vacancy flux toward the metal/oxide interface, resulting in a higher susceptibility to local breakdown as compared to that with a lower donor density. The results provide a reasonable explanation why the passive film on commercially pure Ta is more susceptible to corrosion damage than that for the β-Ta coating, although the compact property of passive film formed on commercially pure Ta is comparable to that of the β-Ta coating. The flat band potential (Efb) is a critical parameter used to determine the positions of the semiconductor energy bands with respect to the redox potentials of electro active ions in the electrolyte. These positions are governed mainly by the charge transfer across the semiconductor/electrolyte interface, the contact potential between semiconductor and electrolyte, and the thermodynamic stability of the semiconductor. Hence, Efb provides a useful guide to the understanding of electrochemical behavior of semiconducting passive films. As shown in Table 5, at an identical Ef, the Efb of the tested samples decreases in the order of uncoated Ti-6Al-4V > commercially pure Ta > β-Ta coating. According to the energy-band model [75], at potentials more negative than Efb, a sample covered with an n-type semiconducting passive film behaves as a conductor, because electrons can readily move through this layer (accumulation mode). In contrast, at potentials anodic to the Efb, the extraction of electrons from the passive film surface creates a barrier to electron transport (depletion mode),derived from upward bending of the conduction and valence band borders, inhibiting the oxidation reaction at the electrode surface. Therefore, the passive film with a lower Efb has dielectric behavior at a larger potential range and exhibits a higher corrosion resistance. Furthermore, the Efb of a semiconductor material can be expressed as [58]:

𝐸𝑓𝑏 = −𝐸𝑓0 /𝑞 + 𝛥𝛷𝐻

(6)

where 𝐸𝑓0 is the Fermi level of the semiconductor at the potential of zero charge (PZC), 𝐸𝑓0 /𝑞 stands for the potential drop of space charge layer, ΔΦH is the potential drop across the Helmholtz double layer. At the flat band potential, the potential drop in the space charge layer of the 18


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semiconductor electrode is zero, and hence the difference of potential distribution on the Helmholtz double layer is responsible for the change of flat band potential. Under constant temperature, the adsorption of anions on the electrode surface can cause the change of ΔΦH, resulting in the change of Efb accordingly. Cheng et al. [76] found that the anion concentration near the surface of stainless steel varied directly with the Efb. Therefore, among the three tested specimens, the β-Ta coating shows the lowest anion adsorptive capacity, which is conducive to enhancing the stability of the formed passive film. Under anodic polarization, the space charge layer formed in the n-type semiconducting passive film surface serves as an effective barrier to charge carriers transfer, and thus controls its corrosion resistance. The thickness of the space charge layer (δsc) for an n-type semiconductor can be estimated by the following equation [77]: 2𝜀 𝜀

𝛿𝑠𝑐 = [ 𝑒𝑁𝑟 0 (𝐸 − 𝐸𝑓𝑏 − 𝑑

𝑘𝑇 𝑒

)]1/2

(7)

From Table 5, it is seen that the values of δsc increase with the Ef for all specimens, because the higher potentiostatic polarization potential is expected to promote the formation of thicker passive film [78]. At a given Ef, uncoated Ti-6Al-4V possesses the thickest δsc, followed by β-Ta coating and then commercially pure Ta. Commonly, the thickness of the space charge layer scales with the passive film thickness [79] and a relatively thin film has more lattice defects and imperfection than a thicker film [80]. The thickness trend of the δsc for the tested samples is also consistent with that of the calculated passive film thickness grown at 0.8 V shown in Table 4. In accordance with the results reported elsewhere [80], the thinnest passive film on commercially pure Ta shows the highest donor density, while the thickest passive film on uncoated Ti-6Al-4V exhibits the lowest donor density. Besides the carrier density (Nd or Na), the diffusivity (Do) of the point defects in the passive film is another key parameter responsible for the growth and breakdown of the passive film. This parameter qualitatively describes the transport of the point defects in the film from a microscopic perspective based on point defect model (PDM) [81]. The diffusivity (Do) of the point defects in the passive film can be calculated by using two available equations, i.e. low-field equation [82] and high-field equation [83]. Considering the very high field strength (the order of magnitude of ~106 V/cm) for the passive films on the Ti and Ta [84], the high-field equation is more adaptable 19



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to calculate the diffusivity of the point defects in this work [60, 61, 64]:

𝐷𝑜 =

𝑧𝛼𝑖𝑠𝑠 𝑒𝜔2 exp(𝑧𝛼𝜀̅𝐹 /𝑅𝑇)

(8)

where z is the charge number of the mobile point defects in the passive film (z = 2 is used in this work), α is the half-jump distance of migrating point defects, assuming α = 0.25 nm [64], iss is the steady state passive current density, which can be determined from the polarization curves, e is the charge of an electron (1.6 × 10-19C), ω2 is a unknown constant that can be acquired by exponentially fitting the Nd-Ef plot (shown in Fig. 15), 𝜀̅ is the mean field strength of the passive film (approximately 3.0 × 106 V/cm [84]), F is the Faraday constant (96.485 C/mol), R is the gas constant (8.314 J/K·mol), and T is the temperature in Kelvin (310 K). In order to gain the value of ω2, the dependence of Nd on Ef was fitted using a first-order exponential decay function [60, 64, 67]:

𝑁𝑑 = 𝜔1 𝑒𝑥𝑝(−𝑏𝐸𝑓 ) + 𝜔2

(9)

Fig. 15 presents the variation of Nd for the tested specimens as a function of Ef . It is apparent that Nd exponentially decays with the Ef for the three tested samples. The values of ω2 for the passive films formed on uncoated Ti-6Al-4V, commercially pure Ta and the β-Ta coating are 1.08×1019 cm-3, 9.62×1019 cm-3 and 7.88×1019 cm-3, respectively. Then, the variables iss, ω2, z, e, 𝜀̅, F, R, T, and α can be substituted into Eq. (8) to calculate the diffusivity of the point defects in the passive film, which are shown in Table 5. The calculated values of Do for the β-Ta coating (1.45×10-16 cm2/s) is comparable to that for commercially pure Ta (1.13×10-16 cm2/s), both of which is one order of magnitude lower than that for uncoated Ti-6Al-4V (2.73×10-15 cm2/s). The results offer a valid explanation why uncoated Ti-6Al-4V exhibits a lower donor density and thicker passive films, but shows a lower corrosion resistance as compared to the β-Ta coating.

4. Discussion From the above results of potentiodynamic polarization and EIS measurements, it is apparent that the application of the β-Ta coating not only yields good corrosion protection for Ti–6Al–4V substrate, but also shows a higher resistance to corrosion attack than the coarse-grained pure Ta. Both the nanocrystalline β-Ta coating and its coarse-grained counterpart have the same chemical composition of the passive film, however, they are characteristic of different electrochemical corrosion behaviors in identical medium, which is intimately connected with the nanoscale 20


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microstructure and preferred crystallographic orientation. Unlike traditional coarse-grained materials, nanocrystalline materials have unique physical, chemical and mechanical properties because of their high volume fraction of grain boundaries resulting from the nanoscale grain size [31, 35, 85]. Generally, nanoscale structures are expected to have both positive and negative impacts on corrosion resistance, depending on whether they contain sufficient concentrations of passive elements that assist in the growth of an insoluble passive film [86, 87]. If the formed passive film is soluble, the corrosion resistance would be decreased by nano-structuring, and vice versa. In this study, the nanoscale structure is beneficial to the enhanced corrosion resistance of the β-Ta coating, since the high density of the grain boundaries can act as short diffusion paths for passive elements and facilitate rapid formation of a stable passive layer. Secondly, in chloride-containing media, chloride is the most commonly encountered aggressive anion causing pitting corrosion for many metals and alloys. Several studies have already confirmed that nanoscale materials with special surface conditions hindered the adsorption ability of chloride ions from passive films formed on their surface, thereby lowering the occurrence of local corrosion [86, 88, 89]. This is in accordance with the results of the Efb analysis obtained from the passive films on both the β-Ta coating and commercially pure Ta. Thirdly, a number of investigations have also shown that nanoscale structures decreased the carrier density of the passive film formed on various metals and alloys [88-90]. For example, Ye et al. [90] reported that improvements in pitting resistance of nanostrucuted 309 stainless steel in acidic NaCl was associated with the formation of a more compact and stable passive film with a lower donor density. This is also consistent with our results which show that the passive film on the β-Ta coating with nanoscale structure has a lower carrier density than that on commercially pure Ta. Previous studies have shown that crystallographic orientation also played an important role in the kinetics of oxide growth and the semiconductor properties of the passive films formed on various metals and alloys [51, 61, 91, 92]. In terms of geometric configuration, the close packed planes with a higher atomic coordination have maximum bond energies in the plane and a minimum for a constituent out of the plane. Therefore, the activation energy for dissolution of the close-packed planes is higher than the relatively less densely planes, denoting that the more loosely packed planes show a higher surface reactivity and are more susceptible to dissolution 21



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than the close-packed planes [93]. The faster dissolution of the less densely packed planes is kinetically more appropriate for taking place the oxygen evolution reaction and exhibits faster oxide film growth. Furthermore, some researchers have found that for both hexagonal Ti and cubic Nb [61, 91] the passive film formed on the less densely packed planes exhibited a lower donor density than that of those planes with a higher atomic surface density, which lead to a lower corrosion rate for the less densely packed planes. As shown in Fig. 2, the tetragonal β-Ta coating and commercially pure α-Ta, with body-centered cubic structure, exhibit (002) and (110) preferred orientations, respectively. Table 6 compares the crystallographic properties of the β-Ta coating with those of commercially pure α-Ta. From Table 6, it can be seen that the preferred (002) plane of the β-Ta coating has an atomic density of 9.63 × 1013 atoms/cm2, which is more than one order of magnitude lower than that for the preferred (110) plane for commercially pure α-Ta (1.29 × 1015 atoms/cm2). Thus, the (002) plane for the β-Ta coating is more favorable for rapid growth of a robust passive film with a lower donor density as compared to the (110) plane for commercially pure α-Ta. This is also consistent with the data obtained from the OCP and Mott-Schottky measurements. The finding obtained from this study demonstrates the importance of nano-structuring and crystal orientation control in developing highly corrosion-resistant Ta coatings for biomedical applications. Such advanced understanding could be applied to guide the design of high performance biomedical coatings, enabled by the processing route established in this work.

5. Conclusions A novel β-Ta nanocrystalline coating with a thickness of ~40 μm were prepared onto a Ti– 6Al–4V substrate by reactive sputter-deposition using a double glow discharge plasma technique. Following phase and microstructural characterization, the electrochemical corrosion behavior, compact property, donor density, flatband potential, and diffusivity of the point defect of the passive film formed on the β-Ta nanocrystalline coating in Ringer’s physiological solution at 37 °C were investigated by open circuit potential (OCP), potentiodynamic polarization, potentiostatic polarization, EIS, Mott-Schottky analysis, and the point defect model. Based on the experimental results, the following conclusions were obtained: (1)The EOCP measurements indicated that the β-Ta coating exhibited a higher oxide film growth 22


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rate compared with the two coarse grained materials: Ti–6Al–4V and commercially pure Ta, suggesting that nanoscale structure and the (002) preferred orientation with a lower atomic surface density could facilitate rapid formation of a stable passive film. (2) The results of potentiodynamic polarization and EIS measurements revealed that the β-Ta coating exhibited a higher electrochemical corrosion resistance than uncoated Ti–6Al–4V substrate and commercially pure Ta. (3) Potentiostatic polarization analysis suggested that the compact property of the passive film on the β-Ta coating was slightly superior to that on commercially pure Ta, but both of which were markedly higher than that on uncoated Ti-6Al-4V. (4) The Mott-Schottky analysis revealed that the passive film formed on the β-Ta coating exhibited n-type semiconductor behavior with a donor density of the order of 1019cm−3, which was significantly smaller than that for the passive film formed on commercially pure Ta. The Efb of the β-Ta coating is lower than that of uncoated Ti-6Al-4V alloy and commercially pure Ta, implying that the β-Ta coating has a lower anion adsorptive capacity as compared to the two reference materials. (5) The diffusivity of the point defect, Do, was calculated based on the point defect model. The calculated values of Do for the β-Ta coating (1.45×10-16 cm2/s) are comparable to that for commercially pure Ta (1.13×10-16 cm2/s), both of which is one order of magnitude lower than that for uncoated Ti-6Al-4V (2.73×10-15 cm2/s).

Acknowledgements The authors acknowledge the financial support from the National Nature Science Foundation of China under Grant No. 51374130. This work is supported by Funding of Jiangsu Innovation Program for Graduate Education, the Fundamental Research Funds for the Central Universities (CXLX13-151). This study is also supported by the Australian Research Council Discovery Project (DP150102417).

References (1) Long, M.; Rack, H. J. Titanium alloys in total joint replacement — a materials science 23



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For Table of Contents Use Only

(004) (002)

5 1/nm

10 μm 100 nm

Electrochemical corrosion behavior of nanocrystalline β-Ta coating for biomedical applications Lin lin Liu,

Jiang Xu, Xiaolin Lu, Paul Munroe, Zong-Han Xie

B


A

32

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1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 1. SEM image of the microstructure of the commercially pure Ta sample.

32


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(202)β

(004)β

As-deposited coating

(110) a

(200)a

(211)a

Pure Ta

20

30

40

50

60

70

(220)a

Intensity (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

(002)β


80

90

2θ (degree) Figure 2. XRD patterns recorded from the as-deposited Ta coating and the commercially pure Ta sample.

33


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1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(a)

(b) A

B 10 μm

(c)

20 μm

Deposition layer

Diffusion layer

Substrate

(d)

Figure 3. (a) SEM image showing the surface morphology, (b) SEM image and (c, d) the corresponding 2D EDS elemental maps for Ta and Ti for the as-deposited Ta coating.

34



(004)

(a)

(002)

(b)

5 1/nm

100 nm

100 nm 30 25

Percentage (%)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 36 of 53

20

(c) 125 grains dave = 22 nm

15 10 5 0

12 14 16 18 20 22 24 26 28 30 32

Grain size (nm)

(d) 0.262 nm —— lattice spacing of β-Ta (002) plane

0.237 nm —— lattice spacing of β-Ta (202) plane

5 nm

Figure 4. (a) Bright-field and (b) dark-field plan-view TEM images, (c) corresponding statistical histogram of the Ta grain sizes and (d) HRTEM image for the as-deposited Ta coating. The inset in (a) is a selected area electron diffraction (SAED) pattern.

35


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(a)

200 nm

(b)

200 nm

Figure 5. Bright-field cross-sectional TEM images taken from (a) the intermediate coating layer and (b) coating/substrate interface marked as squares A and B in Fig. 3(b).

36



E = -0.30e^(-t/1.34) - 0.12e^(-t/16.38) - 0.02

0.0

EOCP (V vs. SCE)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 38 of 53

-0.1 E = -0.16e^(-t/2.91) - 0.19e^(-t/17.69) - 0.08

-0.2 -0.3

E = -0.07e^(-t/2.60) - 0.18e^(-t/25.22) - 0.24

-0.4

Ti-6Al-4V

-0.5 -0.6

Pure Ta

β-Ta coating 0

10

20

30

40

50

60

Time (min) Figure 6. Evolution of the open circuit potential (EOCP) for uncoated Ti-6Al-4V, commercially pure Ta and β-Ta coating in Ringer’s physiological solution at 37 °C.

37



1.0x10

6

8.0x10

5

(a)

4.0x10

5

Ti-6Al-4V 2.0x10

10

5

10

4

10

3

10

2

10

1

2

5

6

(b)

80 70

|Z| (Ω cm )

6.0x10

10

90

Pure Ta

5

2.0x10

5

4.0x10

5

6.0x10

5

8.0x10

5

1.0x10

6

50 40

β-Ta coating 0.0

60

10

-2

Ti-6Al-4V

30

Pure Ta

20

β-Ta coating

10

10

-1

10

0

10

1

10

2

10

3

10

4

0

10

5

Frequency (Hz)

Z'(Ω cm2)

Figure 7. (a) Nyquist and (b) Bode plots of uncoated Ti-6Al-4V, commercially pure Ta and β-Ta coating at respective open circuit potentials in Ringer’s physiological solution at 37 °C. Symbols are experimental data and solid lines are fitted results.

38


-Phase angle (degrees)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

-Z''(Ω cm2)

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Figure 8. Electronic equivalent circuit (EEC) used in the fitting procedure of the EIS experimental data.

39



10

-3

10

-4

10

-5

10

-6

10

-7

10

-8

10

-9

10

(a)

i

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

i (A cm-2)

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(b)

βa1

βc

βa2

icorr1 Ti-6Al-4V

icorr2

Pure Ta

βa2 > βa1

β-Ta coating

Ecorr2

-10

-0.6

-0.3

0.0

0.3

0.6

0.9

1.2

1.5

Ecorr1

E

E (V vs. SCE) Figure 9. (a) Potentiodynamic polarization curves of uncoated Ti-6Al-4V, commercially pure Ta and β-Ta coating in Ringer’s physiological solution at 37 °C; (b) The Evans diagram depicting changes in Ecorr and icorr.

40


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

i (A cm-2)


10

-1

10

-2

10

-3

10

-4

10

-5

10

-6

10

-7

Page 42 of 53

Ti-6Al-4V Pure Ta

β-Ta coating y= y= y=

1

10

-0.9

-0.9

8x -

6x

-0.8

8x 1

.96

- 2.

58

3.3

9

100

1000

Time (s) Figure 10. Double-log plots of current-time for uncoated Ti-6Al-4V, commercially pure Ta and β-Ta coating potentiostatically polarizated at 0.8 V in Ringer’s physiological solution.

41



1.0x10

6

8.0x10

5

6.0x10

5

4.0x10

5

2.0x10

5

(a)

10

7

10

6

10

5

10

4

10

3

10

2

10

1

90

(b)

80 70 60

2

6

|Z| (Ω cm )

1.2x10

Ti-6Al-4V Pure Ta

β-Ta coating

0.0 5

5

5

5

6

2.0x10 4.0x10 6.0x10 8.0x10 1.0x10 1.2x10

6

50 40 30

Ti-6Al-4V

10

-2

10

-1

Pure Ta

20

β-Ta coating

10

10

0

10

1

10

2

10

3

10

0

4

10

5

Frequency (Hz)

Z'(Ω cm2)

Figure 11. (a) Nyquist and (b) Bode plots of uncoated Ti-6Al-4V, commercially pure Ta and β-Ta coating after potentiostatic polarization at 0.8 V for 1 h in Ringer’s physiological solution. Symbols are experimental data and solid lines are fitted results.

42


-Phase angle (degrees)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

-Z''(Ω cm2)

Page 43 of 53


Ta4f

Ta4d

Ta4p

OKLL

Intensity (a.u.)

C1s

O1s

(a)

1000 900 800 700 600 500 400 300 200 100

0

Binding energy (eV)

(b) Intensity (a.u.)

Ta 4f

32

31

30

29

28

27

26

25

24

23

22

Binding energy (eV)

(c)

O 1s

Intensity (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 44 of 53

538

536

534

532

530

528

526

Binding energy (eV) Figure 12. (a) XPS survey spectrum and (b, c)high-resolution XPS spectra for Ta 4f and O 1s for the passive film formed on the β-Ta coating after potentiostatic polarization at 0.8 V for 1 h in Ringer’s physiological solution.

43


-2

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


C (μF cm )

Page 45 of 53

35 30 25 20 15 10 5 13 12 11 10 9 8 7 6 5

Ti-6Al-4V

6.0 5.8 5.6

Pure Ta

5.4 5.2 0.0

0.2

0.4

0.6

9 8 7 6 5 4 3

-2.0

0.8

β-Ta coating

-1.5

-1.0

-0.5

0.0

0.5

1.0

E (V vs. SCE) Figure 13. C-E curves for the passive films formed on uncoated Ti-6Al-4V, commercially pure Ta and β-Ta coating after potentiostatic polarization at 0.8 V for 1 h in Ringer’s physiological solution.

44


4 -2

10

5x10

10

4x10

10

3x10

10

2x10

10

1x10

10

(a)

semconducting behavior

dielectric behavior

-2

6x10

Page 46 of 53

0.4 V 0.6V 0.8 V 1.0 V

0 -1.2

-0.8

-0.4

0.0

0.4

0.8

1.2

3.5x10

10

3.0x10

10

2.5x10

10

2.0x10

10

1.5x10

10

1.0x10

10

(b) dielectric behavior inversion behavior

-2

-2

4

Csc (F cm )

E (V vs. SCE)

5.0x10

0.4 V 0.6 V 0.8 V 1.0 V

semconducting behavior

9

-1.6

-1.2

-0.8

-0.4

0.0

0.4

0.8

1.2

E (V vs. SCE) 1.2x10

11

1.0x10

11

8.0x10

10

6.0x10

10

4.0x10

10

2.0x10

10

dielectric behavior

(c)

0.4 V 0.6 V 0.8 V 1.0 V

-2

-2

4

Csc (F cm )

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Csc (F cm )


0.0 -2.0

semconducting behavior -1.6

-1.2

-0.8

-0.4

0.0

0.4

0.8

1.2

E (V vs. SCE) Figure 14. Mott-Schottky plots of the passive films formed at different potentials (0.4, 0.6, 0.8, 1.0 V) on (a) uncoated Ti-6Al-4V, (b) commercially pure Ta and (c) β-Ta coating in Ringer’s physiological solution.

45


-3

1.4x10

20

1.2x10

20

1.0x10

20

8.0x10

19

6.0x10

19

4.0x10

19

2.0x10

19

(a)

-3

Nd (cm )

1.80x10

20

1.65x10

20

1.50x10

20

1.35x10

20

1.20x10

20

1.05x10

20

20

Nd = 4.33×10 exp(-3.65Ef) + 1.08×10

19

Ti-6Al-4V

0.3

0.4

0.5

0.6

0.7

0.8

0.9

1.0

1.1

E (V vs. SCE)

(b)

20

Nd = 2.80×10 exp(-3.24Ef) + 9.62×10

19

Pure Ta

0.3

0.4

0.5

0.6

0.7

0.8

0.9

1.0

1.1

Nd (cm )

E (V vs. SCE)

-3

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Nd (cm )

Page 47 of 53

1x10

20

9x10

19

8x10

19

7x10

19

6x10

19

5x10

19

4x10

19

3x10

19

(c)

20

Nd = 1.90×10 exp(-1.64Ef) + 7.88×10

19

β-Ta coating

0.3

0.4

0.5

0.6

0.7

0.8

0.9

1.0

1.1

E (V vs. SCE) Figure 15. Donor density (Nd) in the passive films formed on (a) uncoated Ti-6Al-4V, (b) commercially pure Ta and (c) β-Ta coating in Ringer’s physiological solution as a function of film formation potential (Ef). The solid lines are the exponential fit of Nd on Ef.

46



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 48 of 53

Table 1. The calculated texture coefficients of different planes for the commercially pure Ta sample and β-Ta coating. Commercially pure Ta

Crystal plane T(hkl)

β-Ta coating

Crystal plane T(hkl)

(110) 1.34

(200) 0.87 (002) 1.76

(211) 0.75

(220) 1.03 (202) 0.25

47


Page 49 of 53

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Table 2. Electrochemical parameters derived from impedance fitting for investigated specimens at their respective open circuit potentials in Ringer’s physiological solution at 37 °C.

Samples Rs (Ω cm2) Qp (Ω-1 cm-2 sn) n Rp (Ω cm2) Cp (μF cm-2) τ (s) χ2

Ti-6Al-4V 26.25±0.17 (1.15±0.01) × 10-5 0.911±0.001 (5.75±0.15) × 105 8.43 4.85 9.35 × 10-4

Pure Ta 25.21±0.20 (1.26±0.01) × 10-5 0.919±0.001 (1.21±0.04) × 106 6.19 7.49 5.61 × 10-4

β-Ta coating 13.13±0.10 (1.00±0.01) × 10-5 0.930±0.001 (2.78±0.14) × 106 5.10 14.18 7.10 × 10-4

48



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 50 of 53

Table 3. Electrochemical parameters extracted from potentiodynamic curves of investigated specimens in Ringer’s physiological solution at 37 °C.

Samples Ecorr (V vs. SCE) βa (mV/decade) -βc (mV/decade) icorr (A cm-2) ipass1 (A cm-2) 1

Ti-6Al-4V -0.24 158.08 116.48 4.20 × 10-7 9.86 × 10-6

Pure Ta -0.21 249.25 118.04 2.99 × 10-7 3.63 × 10-6

β-Ta coating -0.15 248.02 112.54 1.73 × 10-7 9.24 × 10-7

The passive current densities were derived at 0.60 V vs. SCE.

49


Page 51 of 53

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Table 4. Electrochemical parameters derived from impedance fitting for investigated specimens after potentiostatic polarization at 0.8 V for 1 h in Ringer’s physiological solution.

Samples Rs (Ω cm2) Qp (Ω-1 cm-2 sn) n Rp (Ω cm2) Cp (μF cm-2) dp (nm) χ2

Ti-6Al-4V 21.92±0.32 (7.46±0.10) × 10-6 0.880±0.002 (6.08±0.18) × 105 2.27 23.39 1.34 × 10-3

Pure Ta 22.55±0.28 (1.36±0.02) × 10-5 0.905±0.003 (1.43±0.03) × 105 5.82 3.80 1.40 × 10-3

β-Ta coating 22.75±0.24 (4.05±0.04) × 10-6 0.900±0.001 (2.47±0.10) × 106 1.44 15.36 1.22 × 10-3

50



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 52 of 53

Table 5. Summary of parameters derived from capacitance measurements and PDM analysis for the passive films formed on the investigated specimens in Ringer’s physiological solution. Ti-6Al-4V

Pure Ta

β-Ta coating

0.4 V

Nd (×1019cm-3) Efb (V) δsc (nm)

11.13 -1.15 9.61

17.26 -1.66 5.74

9.95 -2.03 8.21

0.6 V


5.92 -0.90 12.96

13.75 -1.57 6.60

7.06 -1.92 9.93

0.8 V


3.42 -0.80 17.61

11.53 -1.53 7.47

5.31 -1.83 11.70

1.0 V


2.20 -0.78 23.16

10.81 -1.52 8.02

3.70 -1.74 14.30

1.08 9.50 27.33

9.62 3.49 1.13

7.88 0.92 1.45

Samples

ω2(×1019 cm-3) iss (×10-6A·cm-2) Do (×10-16cm2/s)

51


Page 53 of 53

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Table 6. Comparison of crystallographic properties for commercially pure α-Ta and β-Ta coating. Phase

α-Ta

β-Ta

JCPDS no.

04-0788

25-1280

Structure

BCC, S.G. Im3m

Tetragonal, S.G. P42/mnm

Lattice parameters

a=b=c=0.3305 nm

Plane

a=b=1.019 nm

(110)

c=0.5313 nm

(002)

Surface area /cm2

1.54×10-15

1.04×10-14

Number of atoms /surface plane

2

1

Atomic surface density /number of atoms cm-2

1.29×1015

9.63×1013

52


Electrochemical Corrosion Behavior of Nanocrystalline Î²-Ta Coating

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