MRI-compatible Nb–60Ta–2Zr alloy for vascular stents: Electrochemical corrosion behavior in simulated plasma solution

Materials Science and Engineering C 56 (2015) 205–214

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Materials Science and Engineering C journal homepage: www.elsevier.com/locate/msec

MRI-compatible Nb–60Ta–2Zr alloy for vascular stents: Electrochemical corrosion behavior in simulated plasma solution Hui-Zhe Li, Xu Zhao, Jian Xu ⁎ Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China

a r t i c l e

i n f o

Article history: Received 10 August 2014 Received in revised form 26 May 2015 Accepted 12 June 2015 Available online 16 June 2015 Keywords: Niobium Corrosion Simulated plasma Vascular stents Passive film

a b s t r a c t Using revised simulated body fluid (r-SBF), the electrochemical corrosion behavior of an Nb–60Ta–2Zr alloy for MRI compatible vascular stents was characterized in vitro. As indicated by XPS analysis, the surface passive oxide film of approximately 1.3 nm thickness was identified as a mixture of Nb2O5, Ta2O5 and ZrO2 after immersion in the r-SBF. The Nb–60Ta–2Zr alloy manifests a low corrosion rate and high polarization resistance similar to pure Nb and Ta, as shown by the potentiodynamic polarization curves and EIS. Unlike 316L stainless steel and the L605 Co–Cr alloy, no localized corrosion has been detected. Semiconducting property of passive film on the Nb–60Ta– 2Zr alloy was identified as the n-type, with growth mechanism of high-field controlled growth. The excellent corrosion resistance in simulated human blood enviroment renders the Nb-60Ta-2Zr alloy promising as stent candidate material. © 2015 Elsevier B.V. All rights reserved.

1. Introduction For the fabrication of vascular stents, widely used in percutaneous coronary intervention procedures to support the artery wall and improve blood flow, the present materials mainly consist of 316L stainless steel (SS), cobalt–chromium alloys, nickel–titanium shape memory alloys and pure tantalum [1–3]. However, these metals more or less contain leachable toxic elements, such as nickel, cobalt, chromium and molybdenum. After long-term exposure to the blood environment, these metallic ions are released into the tissues, resulting in localized allergic reactions [4,5]. In addition to their radio-opacity to X-ray fluoroscopy, the compatibility of stent materials with magnetic resonance imaging (MRI) has become increasingly significant. MRI is a non-invasive diagnostic tool that does not pose the danger of exposing patients to ionizing radiation, and it avoids using nephrotoxic contrast agents [6]. Cardiac MRI examination is also important in the diagnosis of cardiovascular disease, such as in the assessment of cardiac vessel morphology, plaque characterization [7]. However, under a magnetic field with intense strength, paramagnetic metals with high volume magnetic susceptibility (χv), such as 316L SS and Co–Cr alloys can generate artifacts in the images as a result of the distortion of the magnetic field. To minimize the image artifacts, the χv of stent metal is required to be as close as possible to that of the surrounding tissue, which is in a range of (− 11.0)–(− 7.0) × 10−6 [8]. Thus, in regard to diagnostic

⁎ Corresponding author. E-mail address: [email protected] (J. Xu).

http://dx.doi.org/10.1016/j.msec.2015.06.027 0928-4931/© 2015 Elsevier B.V. All rights reserved.

imaging, developing new stent materials with better MRI compatibility is of considerable interest. To this end, niobium-based alloys were developed recently, such as the Nb–28Ta–3.8W–1.3Zr [9,10] and the Nb–60Ta–2Zr [11,12]. Besides the permanent metallic stents, a considerable effort has been made to develop the stents composed of bioabsorbable/biodegradable material [13,14]. To this end, the magnesium [15–17], iron [18] and zinc [19] were selected as the candidate materials. Among them, the magnesium stents are more promising. However, it remains problematic how to control and to compromise the degradation and mechanical integrity of a stent after deployment at a narrowed arterial vessel, until healing and re-endothelialization is obtained [16,17]. As is well documented, human blood plasma is chemically aggressive to implanted metals, particularly due to the presence of high concentration of chloride ions (~103 mmol/L) [20,21]. Apart from the corrosion-induced metallic ion release, the cavities generated by dissolution of the stent struts become the sources of crack initiation under mechanical loading [22]. Thus, corrosion resistance of a stent material to the blood environment is vital to the success and biosafety of a stent during its lifetime. In this work, the electrochemical corrosion behavior in a simulated plasma solution of an MRI compatible alloy, Nb–60Ta–2Zr, was investigated based on the open-circuit potential, potentiodynamic polarization, potentiostatic, electrochemical impedance spectra and Mott–Schottky analysis. For the purpose of comparison, typical stent metals, such as 316L SS, the L605 Co–Cr alloy, as well as pure Ta and pure Nb, were selected to conduct parallel assessments. Finally, the nature and stability of passive film of these alloys in simulated blood environment as well as their correlation with bio-safety are discussed.

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2. Experimental 2.1. Material and sample preparation An alloy ingot with a nominal composition of Nb–60Ta–2Zr (weight percentage, wt.%) was fabricated using the electron beam melting (EBM) technique. The EBM system was equipped with an electron gun with an accelerating voltage of 25 kV and a power of 120 kW. Temperature on the surface of melted materials is higher than 3000 °C under EBM. Pressure in the vacuum chamber was maintained to be lower than 6 × 10−2 Pa. After the melting, the ingot was cooled in the furnace down to room temperature. Subsequently, it was forged after heating at 1100 °C for 1 h. The as-forged ingot was annealed at 1300 °C for 1 h in a furnace under flowing purified argon. The microstructure and mechanical properties of the annealed Nb–60Ta–2Zr alloy were presented in our previous work [11]. The annealed alloy formed a single-phase solid solution with bcc structure and averaged grain size of 30 μm, and possesses the Young's modulus of 149 GPa, yield strength of ~330 MPa and elongation of ~24%. Commercial pure Ta (99.98%), pure Nb (99.95%), 316L SS (ASTM F138) and L605 Co–Cr alloy (ASTM F90) with annealing states were obtained in plate and bulk form. The pure Ta and Nb bars were purchased from the Ningxia Orient Tantalum Industry Co., Ltd., while the 316L SS and L605 Co–Cr alloy was supplied by the Trauson (China) Medical Instrument Co. Ltd., and Shenzhen Beikehangfei Biomedical Engineering Co. Ltd., respectively. The chemical compositions (wt.%) of 316L SS and the L605 Co–Cr alloy were presented elsewhere [12]. For the electrochemical measurements, the specimens were prepared with dimensions of 10 mm × 10 mm and thickness of 3 mm or 10 mm. They were connected to copper wire, embedded in a polytetrafluoroethylene (PTFE) holder, and then mounted in epoxy resin to ensure a tight seal between the specimen and the PTFE holder in order to avoid the occurrence of crevice corrosion. The surface area of the specimens that was exposed to the test solution was approximately 1 cm2. Prior to each measurement, the sample surface was ground using abrasive paper of 1200-grit silicon carbide. The specimens were ultrasonically cleaned and degreased with ethanol and deionized water, then dried using compressed air. To determine the surface composition, two additional specimens were prepared with dimensions of 10 mm × 10 mm × 2 mm. After mechanical grinding and ultrasonic cleaning in deionized water, one of them was dried in air, and another one was immersed in the test solution for 1 h. A beaker with a solution volume of 300 mL was used for immersion, and the solution was kept at 37 ± 0.5 °C with a water bath. After immersion, the specimen was ultrasonically washed again for 5 min in deionized water and used for subsequent analysis. 2.2. XPS analysis Chemical composition of sample surface of the Nb–60Ta–2Zr alloy was analyzed using X-ray photoelectron spectrometry (XPS). The XPS measurements were performed in an ESCALAB250 surface analysis system (Thermo VG, USA) with a monochromatic Al Kα X-ray source of 1486.6 eV at a take-off angle of 90°. The beam spot size was 500 μm × 500 μm. The pass energy was 50 eV with an energy step size of 0.1 eV. The measured binding energies were calibrated with reference to the C 1s peak with a binding energy value of 284.6 eV. The background was subtracted from the measured spectrum according to Shirley's method [23]. The curve fitting of the XPS spectra was performed by determining the peak position, height, width and Gaussian/Lorentzian ratio, using a commercial software XPSPEAK4.1 for analysis. Depth profile of specimen surface was determined by argon ion beam sputtering to cover an area of 2 mm × 2 mm, using a sputtering voltage of 3 kV and sample current of 2 μA. A sputtering rate of 0.1 nm/s was used to convert the sputtering time into an approximate

sputter depth. Standard Ta2O5 film with known thickness is used to calibrate the sputtering rate. Thickness of surface oxide film on the sample was estimated in terms of the half-maximum of oxygen signal. 2.3. Electrolyte preparation Simulated body plasma solution was used as the electrolyte. The r-SBF was selected to mimic the ion concentrations of human blood plasma [20], with ion concentrations of 142.0 mmol/L Na+, 5.0 mmol/L K+, 1.5 mmol/L Mg2+, 2.5 mmol/L Ca2+, 103.0 mmol/L Clˉ, 27.0 mmol/L HCO3ˉ, 1.0 mmol/L HPO24ˉ, and 0.5 mmol/L SO24ˉ. All reagents were supplied by Alfa Aesar Johnson Matthey Company. HEPES, 2-(4-(2hydroxyethyl)-1-piperazinyl) ethanesulfonic acid, was previously dissolved in 100 mL ultra-pure water with a resistivity of 18.25 MΩ cm. The pH value of the electrolyte was adjusted to 7.4 by the addition of ~0.8 mL 1.0 mol/L NaOH aqueous solution and was measured using a PB-10 pH meter (Sartorius, Germany). 2.4. Electrochemical measurements Electrochemical measurements were performed in a standard three-electrode cell filled with approximately 300 mL r-SBF solution. A saturated calomel electrode (SCE) was used as the reference electrode, a platinum foil was used as the counter electrode, and the specimens were used as the working electrode. The temperature of the electrochemical cell was maintained at 37 ± 0.5 °C with a water bath throughout the tests. For deaeration, the electrolyte was purged with commercial gaseous nitrogen with a purity of 99.99% before and during the electrochemical measurements. All potentials are presented in this paper in terms of the values versus the SCE, with a standard potential of 0.241 V against the standard hydrogen electrode. The open-circuit potential (OCP), potentiodynamic polarization, potentiostatic, electrochemical impedance spectroscopy (EIS), and Mott–Schottky (MS) tests were performed using an electrochemical workstation (EG&G Princeton Applied Research Model 2273) with specific software (Electrochemistry Power Suite). The OCP measurements were performed with duration of 1 h. The potential evolution against immersion time was recorded. The potential of the materials reached a relatively stable state after immersion for 1 h, where the variation in potential was less than 2 mV over a period of 5 min. Prior to the potentiodynamic polarization tests, the specimens were immersed in electrolyte for 1 h, in order to attain a stable state of the OCP. Potentiodynamic polarization was measured in a range from − 0.35 V versus OCP to 1.6 V with a scanning rate of 0.167 mV/s. The corrosion potential, Ecorr, and corrosion current density, icorr, were determined from the polarization curves. The icorr was determined by the intersection of the extrapolated Tafel line at cathodic Tafel regions with the Ecorr. In that region, around 110 mV below the Ecorr, a significantly linear relationship between E and log i is present. The passive current density, ipass, and breakdown potential, Eb, were also obtained to characterize the response in the passive region. The ipass was recorded at a potential of 0 V, which is within the passive region for all of the metals. The surface topography and roughness of the samples subjected to the potentiodynamic tests in the solution were characterized using a laser scanning confocal microscope (LEXT OLS4000; Olympus, USA). The arithmetic mean deviation of the profile (Ra) was used to assess the surface roughness with an evaluation length of 4 mm, which was five times the sampling length (l = 0.8 mm). Potentiostatic tests for the Nb–60Ta–2Zr alloy were conducted at an applied passive potential of −0.2 V, and the current density evolution at this potential was measured for 1 h. Prior to the testing, an initial cathodic cleaning was carried out by applying a potential of − 1.2 V for 1 min, which is used to remove the air-formed oxide on the specimen surface.

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EIS measurements were carried out at the OCP after immersion in the solution for 1 h. The amplitude of sinusoidal AC voltage was ± 10 mV over a frequency range from 100 kHz down to 10 mHz at a rate of 10 points per decade. Both the Bode and Nyquist plots were recorded. The EIS data were fitted with appropriate equivalent circuits using ZsimpWin 3.2 software. For each alloy, all electrochemical measurements were repeated with at least three specimens taken from the bulk materials to gain the standard deviation values. MS tests for the Nb–60Ta–2Zr alloy were performed at a single frequency of 1000 Hz in the potential range from − 0.8 to 1.6 V in r-SBF solution. The amplitude of the superimposed AC signal was 10 mV, and successive steps of 10 mV were applied in the positive direction for the polarization. Before testing, the samples were immersed in the solution for 1 h to form a relatively stable film on the metal surface. 3. Results 3.1. XPS analysis To characterize the nature of oxide film spontaneously formed on the surface, Fig. 1 displays the XPS survey spectrum of the outermost surface of the as-ground Nb–60Ta–2Zr alloy in air and after immersion in r-SBF solution for 1 h. As shown in Fig. 1, peaks for Nb, Ta, Zr, O and C were detected. The C 1s peak at 284.6 eV is attributed to air contamination during the sample transfer into the XPS chamber. It is revealed that the outermost layer in both cases is mainly composed of Nb, Ta, Zr and O. For the as-ground specimen in air, the relative atomic concentration of these elements is 20.2%, 15.2%, 1.5% and 63.0%, respectively. While for the immersed specimen, it is 18.2%, 15.3%, 1.0% and 65.4%, respectively, but the elements of Na, Ca, Cl and P, which is contained in the r-SBF solution, were not detected due to their low contents. Evidently, in both cases, the niobium and tantalum oxides form on the Nb–60Ta–2Zr alloy as major components of the surface layer. Fig. 2(a)–(d) illustrates representative deconvoluted highresolution XPS spectra of Nb 3d, Ta 4f, Zr 3d and O 1s for the asground specimen in air, as the representative. The immersed specimen exhibits very similar spectra (not shown). As usual, if the thickness of oxide film on the metal surface is less than ~ 2 nm, a depth imposed by XPS signal, the signals contributed by matrix elements are inevitably present in the spectra. As a result, the metallic-state elements of Nb and Ta in the substrate are detectable since the surface oxide film on the alloy is thinner than ~2 nm. The center of the Nb 3d5/2 peak is located at a binding energy of 206.8 eV, indicative of the presence of Nb2O5.

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Meanwhile, no detectable peak appears at binding energies ranging from 202.8 to 207.6 eV, which represents the location of binding energy for the Nb 3d5/2 peak in NbO and NbO2 [24]. Thus, Nb2O5 is identified as the major Nb-bearing oxide. The binding energy of the Ta doublet peaks at 25.4 and 27.4 eV represents the Ta 4f7/2 and 4f5/2 in Ta2O5, respectively. This suggests that Ta2O5 is the major Ta-bearing oxide in the film. As shown in Fig. 2(c), the Zr 3d5/2 peak appears at 182.1 eV, proving the presence of ZrO2. By virtue of the XPS analysis combined with ion sputtering, the indepth compositions and thicknesses of the films formed on the asground and immersed Nb–60Ta–2Zr alloy are established. Fig. 3 shows the depth profiles of the oxide layer with elemental distributions of Nb, Ta, Zr and O. The C content decreases within the first ten seconds of sputtering, implying that it is present only at the surface of the sample as a contaminant. As the sputtering depth increases, the O content significantly decreases from ~ 65 at.% to less than 10 at.%, indicating that the surface oxide layer was sputtered away. The contents of Nb, Ta and Zr gradually increase and then become steady at nearly the nominal compositions. The thicknesses of oxide films on the surfaces of the asground and immersed specimens were estimated to be approximately 1.1 nm and 1.3 nm, respectively, based on the decayed O concentration at a half-maximum of the outermost surface. 3.2. Open-circuit potential response Fig. 4 displays the evolution of the OCP over 1 h for immersion of the five investigated metals in r-SBF solution. The OCP curves of Nb–60Ta– 2Zr, Nb and Ta exhibit similar features, as shown in Fig. 4. At the initial stage, the potential rapidly increased towards the more positive side within the first few minutes of the as-ground specimen being immersed in the solution. Subsequently, the potential increased at a slower rate and gradually reached a plateau when the immersion time was extended over 1 h. At this plateau, the potential variation of the Nb–60Ta–2Zr, Ta and Nb is less than 5 mV within 10 min. The trend of monotonically positive shifts in the potential suggests that a protective passive film is formed in situ on the metal surface and continuously grows with the extension of immersion time. In addition, the appearance of a stable potential is indicative of the chemically stable nature of the passive film. In contrast, the OCPs of the 316L SS and the L605 alloy manifest distinctly different behavior. For the 316L SS, the OCP slowly increased within 5 min of the initial immersion at a rate of ~ 0.6 mV/min and then gradually dropped down at the same rate. The OCP of the L605 alloy increased within the initial few minutes as well, but at a rapid rate of ~ 4 mV/min. Afterwards, the OCP dropped down at a rate between ~ 0.8 mV/min to ~ 0.3 mV/min as the immersion time was extended. Finally, the potential value remained almost unchanged. It is well known that a negative shift in potential is associated with either the breakdown and dissolution of a passive film or a lack of film formation [25]. The measured OCP values of all the investigated metals are summarized in Table 1 for comparison. The potential at the last minute of the 1 h immersion was used as the stable-state OCP, although the time-independent state was not fully achieved. As shown in Table 1, after immersion for 1 h, the OCPs of the Nb–60Ta–2Zr, Nb and Ta are substantially comparative and within a range of (−710)–(−810) mV, while the values for the L605 alloy and the 316L SS are about − 590 mV and − 340 mV, respectively. It is noteworthy that these findings are in agreement with the data measured in deaerated Hank's solution at 37 °C, as previously reported by Gurappa [25]. 3.3. Potentiodynamic polarization assessment

Fig. 1. XPS spectrum of the Nb–60Ta–2Zr alloy (a) as-ground in air and (b) immersed in rSBF for 1 h at 37 °C.

Fig. 5 shows a group of potentiodynamic polarization curves for the five investigated metals in r-SBF solution at 37 °C under deaerated conditions. The measured electrochemical properties are listed in Table 1, including the Ecorr, icorr, ipass and Eb. In the cases of Nb–60Ta–

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Fig. 2. Deconvoluted XPS spectra of the as-ground Nb–60Ta–2Zr alloy in air: (a) Nb 3d, (b) Ta 4f, (c) Zr 3d and (d) O 1s.

2Zr, Nb and Ta, the active-to-passive transition following the Tafel region in the curves was not distinct, as shown in Fig. 5. It was then accompanied by the appearance of a stable passive region. No breakdown of the oxide film was detected even at the highest polarization potential of 1.6 V. The Ecorr values for Nb–60Ta–2Zr, Ta and Nb were observed to be in the range of (− 750)–(− 830) mV. As a measure of uniform corrosion, the icorr of the Nb–60Ta–2Zr is 0.04 μA/cm2, which is similar to the value of Ta (0.04 μA/cm2), but 40% smaller than that of Nb (0.07 μA/cm2). A similar scenario is also presented for the ipass. As shown in Table 1, the ipass of Nb–60Ta–2Zr was determined to be 1.34 μA/cm2, which is situated between those of Ta (1.07 μA/cm2) and Nb (1.75 μA/cm2). On the other hand, polarization curves of the 316L SS and L605 alloy are markedly different from those of Nb–60Ta–2Zr, Nb and Ta, as illustrated in Fig. 5. Starting at the Eb potential, the current density rapidly increases as a result of the local breakdown of surface passive film. As shown in Table 1, in contrast to the L605 alloy, the 316L SS manifests not only a more positive breakdown potential (~ 1.0 V versus ~ 0.4 V) but also a wider potential range for passivation. This finding implies that the passive film on the 316L SS surface is more stable. Meanwhile, the ipass of 316L SS and L605 alloy are 0.44 and 0.64 μA/cm2 , respectively, which is slightly lower than that of the Nb–60Ta–2Zr alloy. For the five investigated metals, a reduction of the i pass follows the sequence of Nb, Nb–60Ta–2Zr, Ta, L605 Co–Cr and 316L SS. In this sense, the 316L SS manifests the highest corrosion resistance.

It is noteworthy that the icorr of the Nb–60Ta–2Zr alloy is remarkably reduced by a factor of ~ 3 and ~ 570 with respect to those of 316L SS (~0.12 μA/cm2) and the L605 alloy (~22.9 μA/cm2). As a consequence, it is evident that the uniform corrosion rate of the Nb–60Ta–2Zr alloy is significantly lower than those of 316L SS and the L605 alloy, whereas the dissolution rate is slightly higher after the steady state of passivation was established. 3.4. Potentiostatic test Fig. 6 shows the current evolution at an applied passive potential of −0.2 V on the Nb–60Ta–2Zr alloy in deaerated r-SBF solution at 37 °C. As shown in an inset in Fig. 6, when a passive potential is applied, the current density abruptly dropped to a constant value of 10− 6 A/cm2 within 50 s. The passivation kinetics of the Nb–60Ta–2Zr alloy can be expressed with the following equation: i ¼ A  t −n ;

ð2Þ

where i is the current density, A is a constant, t is the time and n is the passivation rate parameter representing the slope of the log i versus log t plot. The value of n can be considered as a measure of growth models for the oxide film on the alloy. In our case, n ≈ −1 is extracted from the plot of log i versus log t, as shown in Fig. 6. According to the typical interpretation of n in Ref. [26,27], n = − 1 is indicative of a high-field controlled growth of barrier-type passive films, which is

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typical for valve metals, defined as the metals with a large driving force to react with water or oxygen and to form a surface oxide layer [27], such as Ti, Ta, Nb, Zr or Al. In this scenario, the growth of the passive film involves a field-assisted migration of interstitial cation and/or anion vacancies across the oxide film [27–29]. 3.5. Electrochemical impedance spectroscopy

Fig. 3. Depth profiles of the surface of the Nb–60Ta–2Zr alloy with the elemental distribution of Nb, Ta, Zr and O. (a) As-ground in air, (b) immersed in r-SBF for 1 h at 37 °C.

Fig. 7(a) and (b) displays results from EIS for Nb–60Ta–2Zr, Nb and Ta in the r-SBF solution at OCP after immersion for 1 h. The Nyquist plots exhibit a typical feature of a depressed semicircular capacitive loop with a single time constant in the frequency range from 105 to 10−2, as shown in Fig. 7(a). The curvature radius of the capacitive loop for the Nb–60Ta–2Zr is comparable to that of pure Nb but is significantly smaller than that of pure Ta. This suggests that the surface film on the pure Ta is much more corrosion resistant. The Bode plots of Nb–60Ta–2Zr, Nb and Ta present nearly identical features, as shown in Fig. 7(b). In the frequency range of 105–103 Hz, the impedance magnitude, |Z|, remains constant, which corresponds to the resistance of the solution between the specimens and the reference electrodes. In the frequency range of 103 to 10−2 Hz, the |Z| value increases linearly with a slope of approximately −0.9, indicating a capacitor-like behavior at the electrode/electrolyte interface. Furthermore, at the lowest frequency, the |Z| value that is a summation of the solution resistance and polarization resistance reaches an order of 105 Ω cm2 for all three metals, reflecting a strong barrier to transport of cation and anion vacancies then responsible for higher corrosion resistance. From the phase angle curves in the Bode plots, only a single time constant can be identified, as shown in Fig. 7(b), which is consistent with the appearance as shown in Fig. 7(a). The phase angle is independent of the frequency over a wide range of 102–10−1 Hz. The maximum value is attained at approximately −80°, which indicates the formation of a passive oxide film on the soaked metal surface and a near-capacitive response of the film. As the frequency is reduced below 0.1 Hz, the phase angle gradually drops due to the resistance effect of the passive film. A Randles equivalent circuit (EC) with a single time constant was chosen to fit the EIS data, as shown in Fig. 7(c). This type of EC has been widely used for valve metals, including Ti, Nb, Ta and Zr, as reported in previous work [28,30–32]. It can be understood based on Macdonald's Point Defect Model [33]. This model provides a picture of the formation of oxide films in terms of the generation and transport of cation and anion vacancies in a constant electric field. In this EC, Rs represents the resistance of the electrolyte between the working electrode and the reference electrode, while the polarization resistance, Rp, and the constant phase element, CPE, represent the resistance and capacity of the surface layer, respectively. The oxide layer has a resistance Rp parallel to the capacity of CPE and an electrolyte resistance Rs connected in series, as displayed in Fig. 7(c). In addition, with the feature of the depressed semicircle in the Nyquist plots as shown in Fig. 7 (a), the CPE does not behave as a pure capacitor and exhibits non-ideal capacity behavior. The impedance of the CPE can be given as: Z CPE ¼

Fig. 4. Open-circuit potential evolution against immersion time within 1 h for the investigated metals in r-SBF.

1 ; Q ð jωÞn

ð3Þ

where ω is the angular frequency, and Q is a constant with a unit of Ω−1 cm−2sn. The exponent n is a coefficient with a magnitude between 0 and 1, reflecting the deviation from ideal capacitive behavior. When the value of n approaches 1, the CPE behaves as a capacitor, and Q is equivalent to the capacitance. As shown in Fig. 7(a) and (b), the measured data fit well, with a chi-squared value (χ2) of less than 5 × 10− 3 and errors of less than 10% for each component. The parameters obtained from fitting the EIS data are tabulated in Table 2. In all cases, the Rs and n are approximately 10 Ω cm2 and 0.9, respectively. The largest Rp is presented by pure Ta (8.1 × 105 Ω cm2) and is approximately 1.5 times larger than those of the Nb–60Ta–2Zr

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Table 1 Electrochemical properties of the investigated metals in r-SBF measured with OCP and potentiodynamic polarization curves. Metals

OCP (mV)

Ecorr (mV)

icorr (μA/cm2)

ipass (μA/cm2)

Eb (mV)

Nb–60Ta–2Zr Nb Ta 316L SS L605

−715 ± 25 −781 ± 2 −812 ± 7 −337 ± 10 −591 ± 15

−749 ± 33 −817 ± 7 −826 ± 10 −564 ± 7 −750 ± 1

0.04 ± 0.004 0.07 ± 0.005 0.04 ± 0.008 0.13 ± 0.008 22.9 ± 2.8

1.34 ± 0.12 1.75 ± 0.07 1.07 ± 0.03 0.44 ± 0.10 0.64 ± 0.02

1027 ± 57 432 ± 4

(3.3 × 105 Ω cm2) and pure Nb (2.9 × 105 Ω cm2). This suggests that the passive film grown on the surface of pure Ta is more protective. Based on the parallel-plate capacitor model, the oxide film thickness, dfilm, can be estimated using following equation: dfilm ¼

εε0 Aeff ; C eff

ð4Þ

where ε0 is the dielectric permittivity in vacuum (8.85 × 10−12 F m−1); A is the effective surface area and is supposed to be three times the geometric surface area [30,34,35]; Ceff is the effective capacitance of the film; and ε is the dielectric constant of the oxide. The values of ε for Ta2O5, Nb2O5 and ZrO2 are ~ 27, ~ 41 and ~ 12 [27], respectively. The oxide film on the surface of the Nb–60Ta–2Zr alloy is composed of Ta2O5, Nb2O5 and ZrO2, as indicated by XPS analysis. The ε is not available for such a mixed oxide, then it was roughly estimated based on the composition of the oxide film from the XPS results. The ε obtained for the Nb–60Ta–2Zr alloy is 33, when applying the rule of mixtures to the ε values of Ta2O5 (~52 at.%), Nb2O5 (~40 at.%) and ZrO2 (~8 at.%). Furthermore, it is well known that two mathematical expressions can be used to calculate the Ceff from the CPE parameters depending on whether the time-constant distribution is either a surface or normal distribution [36–40]. Given the circumstance of a surface distribution, as proposed by Brug et al. [39], the global admittance response for the electrode includes contributions from every part of the electrode surface. Thus, the Ceff can be calculated according to Eq. (5): C eff ¼ Q 1=n Rðs1‐nÞ=n ;

ð5Þ

where Rs is the solution resistance. On the contrary, in the case of normal distribution, as developed by Hsu and Mansfeld [40], the Rs

Fig. 5. Potentiodynamic polarization curves of the investigated metals in r-SBF from −0.35 V vs. OCP to 1.6 V with a scan rate of 0.167 mV/s in deaerated solution at 37 °C and pH = 7.4.

does not make a contribution to the CPE, and the Ceff is expressed as Eq. (6): ð1‐nÞ=n

C eff ¼ Q 1=n RP

;

ð6Þ

where the Rp is the film resistance. For comparison, Fig. 8(a) displays the Ceff calculated using the two approaches for Nb–60Ta–2Zr, Nb and Ta, denoted as C eff-Brug and C eff-H.M. , respectively. These C eff data are summarized in Table 3. The Ceff-Brug of Nb–60Ta–2Zr (3.2 × 10− 5 F cm− 2) is almost same as that of pure Ta (3.1 × 10− 5 F cm − 2) but less than that of pure Nb (5.8 × 10−5 F cm−2). Similarly, the values of Ceff-H.M. for Nb–60Ta–2Zr (9.2 × 10 − 5 F cm − 2 ), Nb (9.7 × 10 − 5 F cm − 2 ) and Ta (17.8 × 10 − 5 F cm − 2 ) show the same tendency as those of the Ceff-Brug, as illustrated in Table 3. It is noteworthy that for these three metals, the Ceff-H.M. values are approximately 2-fold larger than the Ceff-Brug values. Fig. 8(b) shows the oxide film thicknesses of these metals calculated by substituting the Ceff-Brug and Ceff-H.M. into Eq. (4); these are denoted as dfilm-Brug and dfilm-H.M., respectively. As listed in Table 3, for the Nb–60Ta–2Zr alloy, the dfilm-Brug with an assumption of surface distribution of time constants is ~ 2.6 nm, while the dfilm-H.M. with a normal distribution of time constants is ~0.9 nm, which is quite close to the film thickness of ~ 1 nm determined by XPS. Therefore, these calculations support the suggestion that a normal distribution of time constants is the case of the Nb–60Ta–2Zr alloy in r-SBF solution. 3.6. Mott–Schottky analysis The MS analysis reveals the semiconducting properties of the mixed oxide film grown on the Nb–60Ta–2Zr alloy, including the type of semiconductor, the donor concentration and the flat band potential for the

Fig. 6. Potentiostatic measurements of the Nb–60Ta–2Zr alloy after 1 h of immersion at −0.2 V in r-SBF.

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211

Table 2 EIS equivalent circuit parameters of Nb–60Ta–2Zr, Nb and Ta in deaerated r-SBF at OCP potential after 1 h of immersion. Metals

Rs (Ω cm2)

Rp (105Ω cm2)

Q (10−5 Ω−1 cm−2sn)

n (μA/cm2)

Nb–60Ta–2Zr Nb Ta

9±1 11 ± 3 10 ± 3

3.3 ± 1.1 2.9 ± 0.6 8.1 ± 3.0

6.8 ± 0.1 12.0 ± 1.8 6.5 ± 0.9

0.91 ± 0.01 0.90 ± 0.01 0.91 ± 0.01

to the MS approximation, when 1/C−2 SC is plotted versus V, the slope of the linear plot is 2/εε0eND. When kT/e is negligible (~26 mV at 37 °C), the flatband potential can be obtained by extrapolating the linear portion of the line to 1/ CSC 2 = 0 [43]. Fig. 9 represents the variation of 1/C−2 SC versus the applied potential V for the mixed oxide film on the Nb–60Ta–2Zr alloy. The MS curve comprises two regions [44]. In region I, with potential below −0.25 V, the capacity value is very high and independent of the polarization, and the film accumulates electrical charge and allows current flow. Then, in region II, with potentials of − 0.25–1.0 V, the capacity value decreases as a function of applied potential, which is a typical Mott– Schottky behavior and corresponds to the depletion of donors and the formation of a space charge layer. The positive slope of this linear region indicates an n-type semiconducting property for the film on the Nb–

Fig. 7. Electrochemical impedance spectra of Nb–60Ta–2Zr, Nb and Ta in r-SBF after immersion for 1 h (a) Nyquist diagrams, (b) Bode diagrams and (c) equivalent circuit (EC) for fitting the EIS data. The fitted spectra are presented as solid lines together with the measured data as symbols on the EIS plots.

mixed oxide. In MS tests, the capacitance C is measured as a function of potential E, as follows: C¼−

1 2π f Z

00

;

ð7Þ

where f is the test frequency, and Z″ is the imaginary component of impedance [41,42]. Under depletion conditions, the MS approximation for an n-type semiconductor has the following expression: 1 C 2SC

¼

  2 kT ; V−V FB − εε0 eND e

ð8Þ

where CSC is the space-charge capacitance, ε is the dielectric constant of the film, ε0 is the vacuum permittivity (8.85 × 10−14 F/cm), e is the electric charge (1.602 × 10− 19 C), ND is the donor density, V is the applied voltage, VFB is the flatband potential, k is the Boltzmann constant (1.38 × 10−23 J/K) and T is the temperature in K. According

Fig. 8. (a) Effective capacitance, Ceff, and (b) oxide film thickness, dfilm, of Nb–60Ta–2Zr, Nb and Ta from the assumption of time-constant surface distribution (Brug et al.) and normal distribution (Hsu and Mansfeld).

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Table 3 Calculated values of Ceff and dfilm for Nb–60Ta–2Zr, Nb and Ta based on the assumption of surface or normal distribution, respectively. Metals

Nb–60Ta–2Zr Nb Ta

Surface distribution

Normal distribution

Ceff-Brug (10−5 F cm−2)

dfilm-Brug (nm)

Ceff-H.M. (10−5 F cm−2)

dfilm-H.M. (nm)

3.2 ± 0.2 5.8 ± 1.2 3.1 ± 0.4

2.6 ± 0.2 2.0 ± 0.4 2.3 ± 0.4

9.2 ± 1.1 17.8 ± 2.3 9.7 ± 1.4

0.9 ± 0.1 0.6 ± 0.1 0.8 ± 0.1

60Ta–2Zr alloy. Additionally, the values of ND and VFB are equal to 9.2 × 1020 cm− 3 and − 0.98 V, respectively. The ND value is on the order of 1020 cm− 3 and is in good agreement with the data for the oxide film on pure Ta in 0.15 M NaCl solution, which is approximately 13 × 1020 cm−3, as reported in Ref. [41]. 4. Discussion 4.1. Nature of the surface oxide film on Nb–60Ta–2Zr alloy As shown in Section 3.2, the measured OCP shifts to the positive side during immersion in r-SBF. This suggests that the alloy underwent an anodic dissolution accompanied by the formation of oxide film on the surface, which was mainly composed of Nb2O5 and Ta2O5. Electrochemically, this process can be expressed with the following reactions. 2Nb þ 5H2 O→Nb2 O5 þ 10Hþ þ 10e ðanodic reactionÞ

ð9Þ

2Ta þ 5H2 O→Ta2 O5 þ 10Hþ þ 10e ðanodic reactionÞ

ð10Þ

2Hþ þ 2e→H2 ðcathodic reactionÞ

ð11Þ

In addition, the unstable sub-oxides of NbO and NbO2 are most likely involved in the electrochemical procedure [45–47]. However, they are easily decomposed in aqueous solution, then transform into a highvalence oxide such as Nb2O5 as the following reactions [28]. Nb þ H2 O→NbO þ 2Hþ þ 2e

ð12Þ

NbO þ H2 O→NbO2 þ 2Hþ þ 2e

ð13Þ

2NbO2 þ H2 O→Nb2 O5 þ 2Hþ þ 2e

ð14Þ

The standard Gibbs free energy of formation, ΔfGo, for the relevant oxides is obtained [48]. The ΔfGom is the change in Gibbs free energy for the formation of 1 mol of that oxide from its component elements at their standard states (298.15 K and 0.1 MPa). For NbO, NbO2, Nb2O5, TaO2 and Ta2O5, the ΔfGo are −379, −741, −1766, −180 and − 1911 kJ mol− 1, respectively [48]. In the current case, as shown by the XPS, the oxide compositions are mainly the high-valence Nb2O5 and Ta2O5, with their higher ΔfGo values. Based on the Raman spectrum of Nb exposed to air after immersion in a solution of 0.15 M NaCl, Huang et al. [49] identified the composition of the passive film formed on Nb as NbO2 and Nb2O5 at the metal/solution interface. NbO2 has the propensity to be oxidized into the more stable Nb2O5 after long-term exposure to air. In this regard, the preferential formation of Nb2O5 by Nb is consistent with the current observation for the phase selection of niobium oxide in the Nb–60Ta–2Zr alloy, as shown in Fig. 2(a). On the other hand, apart from the chemical composition, the thickness of the surface oxide layer, dfilm, is also a key issue for the protectiveness of the passive film. As is well known, the rate of the charge-transfer reaction between the metal and the solution can be slowed down due to the presence of the oxide film. For a film of approximately 1 nm thickness, the barrier for electron transfer through the tunnel between the metal and the electrolyte could be enhanced by increasing the film thickness [50]. As usual, XPS and EIS are used to determine the dfilm. In contrast to XPS, the EIS approach has the advantage of revealing the in situ response for the entire reaction. However, the dielectric constant of the mixed oxides on the alloy surface is not obtained. In the current work, the oxide layer on the Nb–60Ta–2Zr alloy is a mixture of Nb2O5 Ta2O5, and ZrO2. Its dielectric constant can be roughly estimated according to the rule of mixtures with the ε value based on the composition of the oxide film determined from the XPS results. According to the XPS analysis, the dfilm values of the native surface film for the as-ground and immersed Nb–60Ta–2Zr alloy are approximately 1.1 nm and 1.3 nm, respectively, as indicated in Fig. 3. These thicknesses are somewhat thinner than those of pure Nb and Ta measured by Tanaka et al. [51]. In their work, the thicknesses of the air-formed oxide films on Nb and Ta were approximately 2.8 nm and 3.0 nm, respectively. In addition, for Nb–28Ta–3.5W–1.3Zr stents, O'Brien et al. [10] showed that the dfilm on the stent surface can be thickened to 5.3 nm after using conventional technologies of laser cutting, dross removal and electropolishing to manufacture stents. On the other hand, the film thickness of the alloy at the initial stage after immersion in r-SBF can also be assessed using EIS. As shown in Table 3, the dfilm values of the Nb–60Ta–2Zr alloy based on the models of normal distribution and surface distribution are approximately 0.9 nm and 2.6 nm, respectively. As noted, the dfilm value of the immersed oxide film from the XPS result is quite close to the dfilm using the model with a normal distribution of time constants. Thus, this model seems more appropriate to the scenario of Nb–60Ta–2Zr alloy in r-SBF. 4.2. Uniform corrosion and its significance

Fig. 9. Mott–Schottky (MS) plot of oxide film formed on Nb–60Ta–2Zr alloy.

To reveal the corrosion mechanism of a ternary alloy, it is necessary to understand the electrochemical characteristics of the alloying component itself. As shown in Fig. 5, the polarization curve of the Nb– 60Ta–2Zr alloy in r-SBF is substantially similar in shape to those of pure Nb and Ta. The curves display a steady current plateau up to 1.6 V, suggesting that no breakdown of the oxide film on the metal surfaces takes place. The current density in the Tafel region, icorr, is significantly small, remaining at an order of magnitude of 10−2 μA/cm2. Wang et al. [30] reported that after immersion for 24 h in aerated PBS, the icorr of pure Nb was measured at ~ 0.08 μA/cm2. This value is similar to the icorr = 0.07 μA/cm2 of pure Nb under the current condition. Moreover, they noted that the addition of BSA to PBS has a twosided effect, either reducing or accelerating the icorr of Nb, which

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depends on the BSA concentration in the solution because of competitive adsorption between BSA and PO3− 4 . However, the current density in the passive region, ipass, does not change much with the BSA concentration, remaining at the magnitude of 2 μA/cm2. In our work, the ipass was determined to be 1.75 μA/cm2, which is comparable to their value. It is interesting to note that even though different simulated physiological solutions were used for in vitro assessment, no significant difference exists for the corrosion behavior of Nb in the passive region, indicating similar passivation features of the oxide film in these solutions. Wilhelmsen et al. [52] reported that tantalum exhibited a potentialindependent ipass with little variation in the electrode potential and the pH value under quasi-steady conditions after immersion. In our cases, as shown in Table 1, the ipass values increase in the sequence of Ta, Nb– 60Ta–2Zr and Nb. The difference between them is greater than 25%. These differences between the three metals are most likely attributed to the different compositions of the oxide films formed on the surface. This suggests that in r-SBF, the protectiveness of Ta2O5 formed on pure Ta is superior to the Nb2O5 formed on pure Nb. The same tendency is also supported by the present EIS results, as illustrated in Fig. 7. For these three metals, the high values of polarization resistance, Rp, at a level of 105 Ω cm2, reflect the high protectiveness of the surface film as well. It is noted that the Rp increases in the order Nb, Nb–60Ta–2Zr and Ta, in accord with the reverse order of ipass. Thus, a high concentration of Ta in the Nb–Ta–Zr series alloys was expected to enhance the corrosion resistance under physiological conditions. Zitter et al. [53] provided a range of current densities of metals in 0.9% NaCl solution and cell culture solution (MEM with fetal calf serum) at 37 °C. It was also shown that the current density is reduced in the sequence of 316L SS, L605 alloy, Nb and Ta. These findings suggest that under the condition of uniform corrosion, the present Nb–60Ta–2Zr alloy has superior corrosion resistance, promising more reliable biocompatibility. 4.3. Pitting corrosion and breakdown of the passive film As shown in Fig. 5, no pitting corrosion in the Nb–60Ta–2Zr alloy was detected with the potentiodynamic polarization measurements. This further suggests that this candidate stent metal would be inert in the blood environment. Meanwhile, materials with lower breakdown potentials are more susceptible to pitting initiation. Pitting corrosion took place in 316L SS at the potential of 1.03 V, which is much higher than the breakdown potential of 0.43 V for the L605 alloy. Obviously, the pitting resistance of these two metals is remarkably inferior to that of our Nb–60Ta–2Zr alloy, for which no localized corrosion appeared even as the potential was raised to 1.6 V. In other words, under the current simulated blood environment, the surface passive film on the Nb–60Ta–2Zr alloy is remarkablely more stable than that on the 316L SS and L605 alloy. Further investigations under in vivo situation are required to clarify the difference in corrosion resistance between these alloys. Furthermore, as a measure to assess the resistance to breakdown of the passive film, the Eb value is influenced by multiple factors, including the surface topography of the sample and testing conditions such as the solution chemistry, pH value, temperature and O2 content. As noted, Cogan et al. [54] reported that the Eb values of 316LVM SS and MP35N Co–Cr alloy under charge-injection conditions were 1.05 V and 0.45 V, respectively, and Sutow et al. [55] found that the Eb of 316L SS varied from 0.35 V to 0.9 V in Ringer's solution at 37 °C. Such differences are most likely caused by the surface configuration, due to variations in the electropolishing time, which was performed at 2.7 V, 100 mA/cm2, room temperature. Tang et al. [56] showed that pitting corrosion occurred at potentials between 0.3 V and 0.5 V for 316L SS immersed in cell culture medium for 0–7 days. The variation of Eb is partially attributed to the fact that the passivation behavior of the samples is sensitive to the test conditions.

213

To our knowledge, the corrosion potential of each of these metals under in vivo conditions has not been well studied aside from the very limited data. This potential is dependent on the location of the tissue where the material is implanted, e.g., in human bone, muscle or the cardiovascular system. To be used as implants, it is necessary for metals to have an Eb higher than the in vivo corrosion potential, which is required for bio-safety. Rosenbloom et al. [57] proposed that metals with an Eb above 0.4 V are acceptable with regards to their corrosion resistance. On the contrary, metals with an Eb below 0.4 V are not qualified candidates for use as implants. Based on this criterion, the breakdown potential of the L605 alloy (~0.4 V) is not preferable. In this sense, the current Nb–60Ta–2Zr alloy provides a remarkable advantage in resisting to localized corrosion in the blood environment compared with 316L SS and the L605 alloy. This is attributed to the high chemical stability of the Nb2O5 and Ta2O5 passive surface films. As already mentioned, the corrosion resistance of metals directly determines the severity of metallic ion release, which is responsible for the biocompatibility of metals. It can be expected that the amount of ions released from the Nb–60Ta–2Zr alloy is much less than those from 316L SS and the L605 alloy, as a consequence of its lower corrosion rate. Furthermore, due to complexity of their toxicity, tolerable doses of the released elements in the human body are another critical issue affecting bio-safety. In contrast to toxic Ni and Cr, the Nb and Ta potentially released from the Nb–60Ta–2Zr alloy have been proven to be biocompatible [58,59]. Although in vitro and in vivo assessments of the biocompatibility of the present alloy are in progress, some properties can be inferred from an alloy with similar chemistry, Nb–28Ta–3.5W–1.3Zr [10]. As illustrated, this Nb-based stent metal shows a bio-response comparable to that of 316L SS, which was supported by the adhesion and growth of human coronary artery endothelial cells. 5. Conclusions Using revised simulated body fluid (r-SBF), the electrochemical corrosion behavior of an Nb–60Ta–2Zr alloy for MRI compatible vascular stents was characterized in vitro. For the alloy immersed in the r-SBF for a short time, the surface passive oxide film of approximately 1.3 nm thickness was identified to be a mixture of Nb2O5, Ta2O5 and ZrO2, as indicated by XPS analysis. Growth mechanism of oxide film was revealed as a high-field controlled growth of barrier-type passive films by potentiostatic test, similar to typical valve metals. The Nb–60Ta–2Zr alloy manifested a lower corrosion rate and a steady passive region up to 1.6 V, similar to pure Nb and Ta, as shown by the potentiodynamic polarization behavior. Meanwhile, unlike 316L SS and the L605 Co–Cr alloy, no localized corrosion was detected due to the highly stable passive film on the Nb–60Ta–2Zr alloy surface. The passive current density (ipass) of the Nb–60Ta–2Zr alloy was situated between those of pure Nb and Ta, suggesting that in r-SBF, the protectiveness of the Ta2O5 film formed on the Ta was superior to the Nb2O5 on the Nb. Additionally, the protective nature of the passive film was also supported by their high polarization resistance (Rp) at a level of 105 Ω cm2, as measured by EIS. Semiconducting property of passive film on the Nb–60Ta–2Zr alloy was identified as the n-type, with a donor density of 9.2 × 1020 cm−3. Based on the present in vitro assessment, it is indicated that the pitting corrosion resistance of the Nb–60Ta–2Zr alloy in simulated human blood environment is significantly superior to that for current stent metals such as 316L SS and Co–Cr alloys. Together with the toxic-element-free composition, it is expected to provide much reliable biocompatibility for stents made of this material. Acknowledgments The authors gratefully acknowledge Prof. K Yang for providing the samples of 316L stainless steel and the L605 Co–Cr alloys, and Prof. H.H. Huang and Dr. T. Zhang for stimulating discussions on corrosion data.

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References [1] B. O'Brien, W. Carroll, The evolution of cardiovascular stent materials and surfaces in response to clinical drivers: a review, Acta Biomater. 5 (2009) 945–958. [2] T. Hanawa, Materials for metallic stents, J. Artif. Organs 12 (2009) 73–79. [3] G. Mani, M.D. Feldman, D. Patel, C.M. Agrawal, Coronary stents: a materials perspective, Biomaterials 28 (2007) 1689–1710. [4] Y. Okazaki, E. Gotoh, T. Manabe, K. Kobayashi, Comparison of metal concentrations in rat tibia tissues with various metallic implants, Biomaterials 25 (2004) 5913–5920. [5] R. Koster, D. Vieluf, M. Kiehn, M. Sommerauer, J. Kahler, S. Baldus, T. Meinertz, C. Hamm, Nickel and molybdenum contact allergies in patients with coronary instent restenosis, Lancet 356 (2000) 1895–1897. [6] G. Katti, S.A. Ara, A. Shireen, Magnetic resonance imaging (MRI)—a review, Int. J. Dent. Clin. 3 (2011) 65–70. [7] G. Constantine, K. Shan, S.D. Flamm, M.U. Sivananthan, Role of MRI in clinical cardiology, Lancet 363 (2004) 2162–2171. [8] J.F. Schenck, The role of magnetic susceptibility in magnetic resonance imaging: MRI magnetic compatibility of the first and second kinds, Med. Phys. 23 (1996) 815–850. [9] B. O'Brien, J. Stinson, W. Carroll, Development of a new niobium-based alloy for vascular stent applications, J. Mech. Behav. Biomed. Mater. 1 (2008) 303–312. [10] B.J. O'Brien, J.S. Stinson, D.A. Boismier, W.M. Carroll, Characterization of an NbTaWZr alloy designed for magnetic resonance angiography compatible stents, Biomaterials 29 (2008) 4540–4545. [11] H.-Z. Li, J. Xu, MRI compatible Nb–Ta–Zr alloys used for vascular stents: optimization for mechanical properties, J. Mech. Behav. Biomed. Mater. 32 (2014) 166–176. [12] X.-M. Li, H.-Z. Li, S.-P. Wang, H.-M. Huang, H.-H. Huang, H.-J. Ai, J. Xu, MRIcompatible Nb–60Ta–2Zr alloy used for vascular stents: Haemocompatibility and its correlation with protein adsorption, Mater. Sci. Eng. C 42 (2014) 385–395. [13] P. Erne, M. Schier, T. Resink, The road to bioabsorbable stents: reaching clinical reality? Cardiovasc. Intervent. Radiol. 29 (2006) 11–16. [14] D. Levi, A. Cheng, Biodegradable implants, in: E.M. Da Cruz, D. Ivy, J. Jaggers (Eds.), Pediatric and Congenital Cardiology, Cardiac Surgery and Intensive CareSpringer, London 2014, pp. 1219–1235. [15] R. Erbel, C. Di Mario, J. Bartunek, J. Bonnier, B. de Bruyne, F.R. Eberli, P. Erne, M. Haude, B. Heublein, M. Horrigan, C. Ilsley, D. Böse, J. Koolen, T.F. Lüscher, N. Weissman, R. Waksman, Temporary scaffolding of coronary arteries with bioabsorbable magnesium stents: a prospective, non-randomised multicentre trial, Lancet 369 (2007) 1869–1875. [16] R. Waksman, R. Erbel, C. Di Mario, J. Bartunek, B. de Bruyne, F.R. Eberli, P. Erne, M. Haude, M. Horrigan, C. Ilsley, D. Böse, H. Bonnier, J. Koolen, T.F. Lüscher, N.J. Weissman, Early- and long-term intravascular ultrasound and angiographic findings after bioabsorbable magnesium stent implantation in human coronary arteries, J. Am. Coll. Cardiol. Intv. 2 (2009) 312–320. [17] M. Haude, R. Erbel, P. Erne, S. Verheye, H. Degen, D. Böse, P. Vermeersch, I. Wijnbergen, N. Weissman, F. Prati, R. Waksman, J. Koolen, Safety and performance of the drug-eluting absorbable metal scaffold (DREAMS) in patients with de-novo coronary lesions: 12 month results of the prospective, multicentre, first-in-man BIOSOLVE-I trial, Lancet 381 (2013) 836–844. [18] H. Hermawan, D. Dube, D. Mantovani, Developments in metallic biodegradable stents, Acta Biomater. 6 (2010) 1693–1697. [19] P.K. Bowen, J. Drelich, J. Goldman, Zinc exhibits ideal physiological corrosion behavior for bioabsorbable stents, Adv. Mater. 25 (2013) 2577–2582. [20] A. Oyane, H.-M. Kim, T. Furuya, T. Kokubo, T. Miyazaki, T. Nakamura, Preparation and assessment of revised simulated body fluids, J. Biomed. Mater. Res. A 65 (2003) 188–195. [21] S. Virtanen, I. Milosev, E. Gomez-Barrena, R. Trebse, J. Salo, Y.T. Konttinen, Special modes of corrosion under physiological and simulated physiological conditions, Acta Biomater. 4 (2008) 468–476. [22] D.O. Halwani, P.G. Anderson, B.C. Brott, A.S. Anayiotos, J.E. Lemons, Clinical devicerelated article surface characterization of explanted endovascular stents: evidence of in vivo corrosion, J. Biomed. Mater. Res. B Appl. Biomater. 95 (2010) 225–238. [23] D.A. Shirley, High-resolution X-ray photoemission spectrum of the valence bands of gold, Phys. Rev. B 5 (1972) 4709–4714. [24] NIST X-ray Photoelectron Spectroscopy Database, NIST Standard Reference Database 20 (2012) Version 4.1. [25] I. Gurappa, Characterization of different materials for corrosion resistance under simulated body fluid conditions, Mater. Charact. 49 (2002) 73–79. [26] J.W. Schultze, M.M. Lohrengel, D. Ross, Nucleation and growth of anodic oxide films, Electrochim. Acta 28 (1983) 973–984. [27] M.M. Lohrengel, Thin anodic oxide layers on aluminium and other valve metals: high field regime, Mater. Sci. Eng. R 11 (1993) 243–294. [28] M. Metikos-Hukovic, A. Kwokal, J. Piljac, The influence of niobium and vanadium on passivity of titanium-based implants in physiological solution, Biomaterials 24 (2003) 3765–3775.

[29] A. Gebert, P.F. Gostin, L. Schultz, Effect of surface finishing of a Zr-based bulk metallic glass on its corrosion behaviour, Corros. Sci. 52 (2010) 1711–1720. [30] W. Wang, F. Mohammadi, A. Alfantazi, Corrosion behaviour of niobium in phosphate buffered saline solutions with different concentrations of bovine serum albumin, Corros. Sci. 57 (2012) 11–21. [31] D. Mareci, R. Chelariu, D.M. Gordin, G. Ungureanu, T. Gloriant, Comparative corrosion study of Ti–Ta alloys for dental applications, Acta Biomater. 5 (2009) 3625–3639. [32] G. Bolat, J. Izquierdo, J.J. Santana, D. Mareci, R.M. Souto, Electrochemical characterization of ZrTi alloys for biomedical applications, Electrochim. Acta 88 (2013) 447–456. [33] D.D. Macdonald, The point defect model for the passive state, J. Electrochem. Soc. 139 (1992) 3434–3449. [34] A.T. Kuhn, Techniques in Electrochemistry, Corrosion and Metal Finishing, John Wiley & Sons, 1987. [35] D. Wallinder, J. Pan, C. Leygraf, A. Delblanc-Bauer, Eis and XPS study of surface modification of 316LVM stainless steel after passivation, Corros. Sci. 41 (1998) 275–289. [36] B. Hirschorn, M.E. Orazem, B. Tribollet, V. Vivier, I. Frateur, M. Musiani, Determination of effective capacitance and film thickness from constant-phase-element parameters, Electrochim. Acta 55 (2010) 6218–6227. [37] F. Mohammadi, T. Nickchi, M.M. Attar, A. Alfantazi, EIS study of potentiostatically formed passive film on 304 stainless steel, Electrochim. Acta 56 (2011) 8727–8733. [38] B. Hirschorn, M.E. Orazem, B. Tribollet, V. Vivier, I. Frateur, M. Musiani, Constantphase-element behavior caused by resistivity distributions in films, J. Electrochem. Soc. 157 (2010) C458. [39] G.J. Brug, A.L.G. van den Eeden, M. Sluyters-Rehbach, J.H. Sluyters, The analysis of electrode impedances complicated by the presence of a constant phase element, J. Electroanal. Chem. Interfacial Electrochem. 176 (1984) 275–295. [40] C.H. Hsu, F. Mansfeld, Technical note: concerning the conversion of the constant phase element parameter Y0 into a capacitance, Corrosion 57 (2001) 747–748. [41] L.V. Taveira, M.F. Montemor, M. Da Cunha Belo, M.G. Ferreira, L.F.P. Dick, Influence of incorporated Mo and Nb on the Mott–Schottky behaviour of anodic films formed on AISI 304L, Corros. Sci. 52 (2010) 2813–2818. [42] K.S. Raja, D.A. Jones, Effects of dissolved oxygen on passive behavior of stainless alloys, Corros. Sci. 48 (2006) 1623–1638. [43] S.P. Harrington, T.M. Devine, Relation between the semiconducting properties of a passive film and reduction reaction rates, J. Electrochem. Soc. 156 (2009) C154. [44] R.A. Silva, M. Walls, B. Rondot, M. Da Cunha Belo, R. Guidoin, Electrochemical and microstructural studies of tantalum and its oxide films for biomedical applications in endovascular surgery, J. Mater. Sci. Mater. Med. 13 (2002) 495–500. [45] C.V. D'Alkaine, L.M.M. de Souza, F.C. Nart, The anodic behaviour of niobium—I. The state of the art, Corros. Sci. 34 (1993) 109–115. [46] C.V. D'Alkaine, L.M.M. de Souza, F.C. Nart, The anodic behaviour of niobium—II. General experimental electrochemical aspects, Corros. Sci. 34 (1993) 117–127. [47] C.O.A. Olsson, D. Landolt, Atmospheric oxidation of a Nb–Zr alloy studied with XPS, Corros. Sci. 46 (2004) 213–224. [48] D. Wagman, W. Evans, V. Parker, R. Schumm, I. Halow, S. Bailey, K. Churney, R. Nuttall, The NBS tables of chemical thermodynamic properties, J. Phys. Chem. Ref. Data 11 (Supplement 2) (1982). [49] B.X. Huang, K. Wang, J.S. Church, Y.-S. Li, Characterization of oxides on niobium by raman and infrared spectroscopy, Electrochim. Acta 44 (1999) 2571–2577. [50] K.E. Heusler, M. Schulze, Electron-transfer reactions at semiconducting anodic niobium oxide films, Electrochim. Acta 20 (1975) 237–244. [51] Y. Tanaka, M. Nakai, T. Akahori, M. Niinomi, Y. Tsutsumi, H. Doi, T. Hanawa, Characterization of air-formed surface oxide film on Ti–29Nb–13Ta–4.6Zr alloy surface using XPS and AES, Corros. Sci. 50 (2008) 2111–2116. [52] W. Wilhelmsen, Passive behaviour of tantalum, Electrochim. Acta 33 (1988) 63–67. [53] H. Zitter, H. Plenk, The electrochemical behavior of metallic implant materials as an indicator of their biocompatibility, J. Biomed. Mater. Res. 21 (1987) 881–896. [54] S.F. Cogan, G.S. Jones, D.V. Hills, J.S. Walter, L.W. Riedy, Comparison of 316LVM and MP35N alloys as charge injection electrodes, J. Biomed. Mater. Res. 28 (1994) 233–240. [55] E.J. Sutow, S.R. Pollack, E. Korostoff, An in vitro investigation of the anodic polarization and capacitance behavior of 316-L stainless steel, J. Biomed. Mater. Res. 10 (1976) 671–693. [56] Y.C. Tang, S. Katsuma, S. Fujimoto, S. Hiromoto, Electrochemical study of Type 304 and 316L stainless steels in simulated body fluids and cell cultures, Acta Biomater. 2 (2006) 709–715. [57] S.N. Rosenbloom, R.A. Corbett, An assessment of ASTM F 2129 electrochemical testing of small medical implants, NACE International, 2007. (Paper No. 07674). [58] E. Eisenbarth, D. Velten, M. Muller, R. Thull, J. Breme, Biocompatibility of betastabilizing elements of titanium alloys, Biomaterials 25 (2004) 5705–5713. [59] H. Matsuno, A. Yokoyama, F. Watari, M. Uo, T. Kawasaki, Biocompatibility and osteogenesis of refractory metal implants, titanium, hafnium, niobium, tantalum and rhenium, Biomaterials 22 (2001) 1253–1262.