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Electrochemical and corrosion behavior of Ti–xAl–yFe alloys prepared by direct metal deposition method
Electrochimica Acta 51 (2006) 2042–2049
Electrochemical and corrosion behavior of Ti–xAl–yFe alloys prepared by direct metal deposition method Natalia V. Pimenova, Thomas L. Starr ∗ Chemical Engineering Department, J.B. Speed School of Engineering, University of Louisville, Louisville, KY 40292, USA Received 29 November 2004; received in revised form 7 July 2005; accepted 7 July 2005 Available online 25 August 2005
Abstract The electrochemical and corrosion behavior of Ti-based alloys was investigated. The direct metal deposition technique was used to fabricate 21 alloys with different ratio of metals (0 ≤ Al ≤ 27 wt.%, 0 ≤ Fe ≤ 25 wt.%). Corrosion resistance of each alloy was evaluated both qualitatively and quantitatively by voltammetric measurements in the simulated human body fluid conditions (Hank’s solution). The corrosion rates of the materials were compared in Hank’s solution using Tafel extrapolation method. Among the Ti–xAl–yFe alloys the Ti–7Al–4Fe alloy exhibited the slowest corrosion rate of 7.7 × 10−4 mm/year and the least value of passive current density (6.3 × 10−3 A/m2 ). The alloy is resistant to pitting corrosion as well. © 2005 Elsevier Ltd. All rights reserved. Keywords: Titanium-based alloy; Polarization; Passivity; Pitting corrosion; Direct metal disposition
1. Introduction Titanium demonstrates an excellent corrosion resistance for long periods in the physiological fluids. The titanium dioxide film (TiO2 ) formed on its surface instantly and spontaneously in the presence of oxygen and moisture is extremely stable and protective. A damaged oxide film can usually recuperate itself relatively easily. Due to its excellent biocompatibility, mechanical properties and corrosion resistance in aggressive media titanium is the most important candidate for use as human body implant material. Although pure titanium offers better corrosion resistance and tissue tolerance, than all other implant materials [1–3], its comparatively low strength restricts its usage to certain application, such as pacemaker cases and reconstruction devices [4,5]. When aluminum and vanadium are added to titanium in small quantities, the strength of the alloy is much increased compared with that of titanium [1]. For example, the Ti–6Al–4V alloy demonstrates higher strength and lower ∗
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elastic modulus than that of pure titanium. The alloy is being used as implant material in high stress-bearing situations, such as hip and knee prostheses, as well as trauma fixation devices for the last two decades. However, since the early 1980’s the use of the alloy has been questioned because of black debris containing high levels of titanium, vanadium and aluminum that are found in surrounding tissue under conditions of high wear, such as in knee and hip implants [4]. Although no toxic effects have been linked to this black debris, safety concerns about the vanadium and aluminum have been raised in literature [2,6–8]. Aluminum [2,6] is poorly absorbed within the gastrointestinal tract. Very little amount gets into the blood stream, and hence it appears nontoxic. Vanadium [7,8] can alter the kinetics of the enzyme activity associated with cells of the inflammatory response. Vanadium contained in this alloy has been associated with potential cytotoxic effects and adverse tissue reactions [2]. The titanium-based alloy with aluminum and iron appears to be more suitable for implant applications. The Ti–5Al–2.5Fe alloy possesses similar corrosion resistance and mechanical properties to those of the Ti–6Al–4V alloy. More importantly, the iron-containing alloy has no apparent toxicity. In soft tissue, aluminum and iron show the reaction
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of “sequestration”, namely the formation of connective tissue membrane around the implant [9–12]. The surgical implant alloys are used for a long time (∼20 years). Therefore, the hydrolysis of different compounds of the passive film is possible. The toxicity of dissolved metal ions may cause pain and restrict action in the tissues adjoining the implants. In order to avoid such situation, the corrosion resistance of bioimplants needs to be high. The aims of the present work are: developing a new technique for discovery of the state-of-the-art materials; making combinatorial libraries of the Ti–Al–Fe alloys of different composition by direct metal deposition (DMD) and comparing corrosion resistance of the alloys in Hank’s solution with that of titanium fabricated by DMD. Fig. 1. Geometry of seven layers of coupon of Ti–xAl–yFe alloy (0 < x < 30 wt.%, 0 < y < 30 wt.%).
2. Experimental 2.1. Materials The Ti–xAl–yFe alloys (0 ≤ x ≤ 27 wt.%, 0 ≤ y ≤ 25 wt.%) were made by DMD technique. DMD 3000 machine (The POM Group, Auburn Hills) at the University of Louisville fabricated 3D-coupons by layer-by-layer deposition utilizing a 3000 W CO2 laser to melt metal powders injected by nozzle and laying down clad tracks on a substrate via a CNC position system. The powder feedstock used for depositing the alloys consisted of blends of pure elemental powders of titanium (Crucible Research, blend B344), aluminum (Praxair, Al-104) and the iron alloy H131 (Intern Mold Steel Inc., the primary source of iron). The Ti–6Al–4V alloy was used as a substrate (Principal Metals Corporation, manufactured by RMI Titanium Company, ASTM B 265 standard, grade 5). The laser forming process and corresponding coordinate system are illustrated in Fig. 1. Three sets of laser formed test coupons were nominally 12 mm wide (x), 14 mm tall (z) and 24 mm long (y). Each coupon was composed of seven layers of the Ti–xAl–yFe alloys (0 ≤ x ≤ 27 wt.%, 0 ≤ y ≤ 25 wt.%). The thickness of each layer was about 2 mm. The composition of each Ti-based alloy is shown in Table 1. Multilayer materials named combinatorial libraries were fabricated for the discovery of the new Ti–xAl–yFe alloys with high resistance to electrochemical corrosion in biofluids. DMD technique allows fabricating seven different compositions in one coupon. Applying combinatorial approach allows making the fabrication of the Ti–xAl–yFe alloys more time and costing effective. For metallographic investigation samples were cut along the x–z plane from the coupon to examine the as-deposited macro- and microstructures. The x–z samples were sectioned in the x–y plane. Sectioning was performed using an Isomet series low speed saw. 1 International Mold Steel’s Premium H-13 is a Cr-based, hot work die and plastic mold steel with the composition: 0.4 wt.% C, 0.4 wt.% Mg, 1 wt.% Si, 5.25 wt.% Cr, 1.35 wt.% Mo, 1 wt.% V and the rest Fe.
The electrochemical investigation was performed for each composition. When the coupon was cut layer-by-layer each slice was mounted in an epoxy resin. The samples were polished sequentially with 120-, 240-, 400- and 600-grit silicon carbide abrasive papers and then with 6 m diamond paste. To prevent oxidation of the surface and to degrease the samples, they were washed in 10% HCl solution and then in ethanol. The composition of each layer was homogeneous. The concentration was different on the boundary of the layers because of diffusion at remelting during the fabrication. Remelting allows making the uniform composition through each layer. The data of energy dispersive X-ray fluorescence analyzer (EDAX® ) linescan (Fig. 2) shows the 0.4 mm interlayer zone between the layers. The composition of interlayer zone is different from the regular composition of each layer. During pre-treatment this layer was removed and only regular composition material was used in the electrochemical experiments.
Fig. 2. EDAX® linescan data of the seven layers coupon (set 3).
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Table 1 Electrochemical parameters of Ti–Al–Fe alloys of different composition No.
Composition
Ecorr (V) vs. Ag/AgCl
icorr (A/m2 )
rcorr (mm/year)
ipas (A/m2 )
Epit (V)
Eprot (V)
10-1 8-1 10-6 10-7 8-2 6-2 10-3 10-2 6-7 10-4 8-4 6-6 8-6 8-7 8-5 6-3 8-3 10-5 6-4 6-5
Pure Ti Ti–3Al Ti–4Al Ti–7Al Ti–18Al Ti–27Al Ti–9Fe Ti–14Fe Ti–25Fe Ti–5Al–8Fe Ti–10Al–8Fe Ti–12Al–8Fe Ti–21Al–5Fe Ti–24Al–5Fe Ti–18Al–1Fe Ti–18Al–7Fe Ti–8Al–10Fe Ti–7Al–4Fe Ti–17Al–3Fe Ti–15Al–6Fe
−0.50 −0.10 −0.30 −0.40 −0.40 −0.42 −0.35 −0.30 −0.62 −0.30 −0.50 −0.24 −0.55 −0.62 −0.54 −0.54 −0.36 −0.38 −0.11 0.05
2.5E−06 5.0E−03 8.5E−03 2.0E−02 4.6E−02 5.5E−02 2.2E−05 2.0E−03 2.0E−02 1.6E−03 1.8E−02 3.0E−02 4.5E−03 5.0E−02 5.0E−02 8.0E−03 1.2E−03 7E−04 1.5E−02 1.0E−02
2.9E−06 5.9E−03 1.0E−03 2.4E−02 4.8E−2 6.5E−02 2.4E−05 2.1E−03 2.0E−02 1.8E−03 2.1E−02 3.7E−02 5.5E−03 6.3E−02 6.0E−02 1.0E−02 1.4E−03 7.7E−04 1.9E−02 1.2E−02
2.1E−05 3.0E−01 – – – – 6.3E−05 4.0E−02 3.0E−01 2.0E−02 – 2.0E−01 7.0E−01 – 3.0E+00 5.0E−01 7.5E+00 6.3E−03 2.1E+00 6.5E+00
2.21 0.50 1.80 1.50 −0.38 −0.40 1.15 0.80 0.70 0.60 0.35 0.50 −0.50 −0.55 0.20 0.40 0.62 1.64 0.62 0.75
2.08 0.40 1.60 0.90 – – 0.80 0.65 0.40 0.42 0.20 0.35 – – 0.00 0.10 0.45 1.42 0.26 0.50
2.2. Optical and electron microscopies The microstructure was investigated by means of optical and scanning electron microscopy. The specimens were prepared following the metallographic techniques used for titanium and its alloys [13]. These consist of grinding up to 600-grit SiC and polishing with 6 m diamond paste to mirror finish. The specimens were cleaned with deionized water and etched in Kroll’s reagent (10 ml of HF, 5 ml of HNO3 and 85 ml of water) for microscopic analysis. Etching was performed by swabbing Kroll’s reagent to the polished surface. The etching time was limited to 10 s to avoid overetching. Microstructure was observed on Leco PME-3 inverted microscope. The image viewing accessories included an Olympus camera and Sony Trinitron viewing system. Microstructural and compositional investigations were performed by using the Jeol JSM-5310 Scanning Electron Microscope (SEM) with EDAX® equipped with sapphire Si(Li) liquid nitrogen cooled EDS detector. The minimum detectable concentration has been roughly estimated to be approximately 1 at.%.
Research 273A potentiostat at a scan rate of 5 mV/s. The potential range of the polarization scans was between −0.6 and 2.5 V2 . Open circuit potential (OCP) measurements were carried out for each alloy in Hank’s solution. The potential was measured for 30 min prior to cyclic polarization experiments. The corrosion current densities determined by linear polarization icorr = k(di/dE)Ecorr were calculated from Stern formula assuming k = 0.026 V [14]. These currents were used in order to calculate the corrosion rates. All measurements were performed in Hank’s solution (g/l): 0.185 CaCl2 , 0.4 KCl, 0.06 KH2 PO4 , 0.1 MgCl2 ·6H2 O, 0.1 MgSO4 ·7H2 O, 8.0 NaCl, 0.35 NaHCO3 , 0.48 Na2 HPO4 and 1.00 d-glucose. The temperature of Hanks solution was kept at 25 ± 1 ◦ C. Nitrogen gas was purged into the solution between the measurements to eliminate the influence of dissolved oxygen on the passive film kinetics. During measurements the solution was blanketed with nitrogen atmosphere.
3. Results and discussion 2.3. Electrochemical experiments A conventional three-electrode cell with a Pt wire gauze with a large surface area as a counter electrode and saturated silver/silver chloride as a reference electrode was used to conduct electrochemical experiments. The geometrical area of the working electrode was 0.8 cm2 . The distance between the Luggin capillary and the working electrode was 3 mm. Prior to the polarization scans the samples were exposed to test solutions for 12 h to attain equilibrium condition. The anodic potentiodynamic polarization and cyclic voltammetry measurements were performed on an EG&G Princeton Applied
3.1. Ti–xAl alloys Table 1 shows composition and electrochemical parameters of all investigated alloys. The micrographs of two alpha Ti–xAl alloys (x = 4, 7 wt.%) are shown in Fig. 3(a and b). 2 The highest value of polarization potential was different for studied alloys and depended on the corrosion resistance of the alloy. For example, the Ti–27Al alloy did not have the passivation region on the anodic polarization curve. The pure Ti had the highest corrosion resistance. Only at anodic potentials higher than 1.2 V the transpassive behavior appeared. For such high corrosion resistant alloy the highest potential was 2.5 V.
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Fig. 3. Scanning electron micrographs of Ti–xAl–yFe alloys: (a) ␣-Ti–4Al, (b) ␣-Ti–7Al, (c) -Ti–9Fe, (d) -Ti–14Fe, (e) ␣/ Ti–5Al–8Fe and (f) ␣/ Ti–7Al–4Fe.
Air-cooling of coupon layers after deposition results in a fine needle-like alpha-phase referred to as fine lamellar or acicular microstructure. All binary Ti–xAl alloys have dominant alpha-phase. When the concentration of aluminum increases the structure might form the brittle Ti3 Al phases in the Tibased matrix. Because DMD technique induces a very high cooling rate (1000–10,000 ◦ C/s) [15], each layer (except the top one) experienced remelting when next layer was fabricated. Sun et al. [16] studied the microstructural changes in pure titanium after laser remelting and established that single ␣-phase was changed to acicular martensitic product. In our case optical microscopy observations indicate dramatic microstructural changes in the melted zone. Fig. 4 shows microstructures of previous layer, interlayer heat affected zone, and the next melted zone. According to EDAX® linescan data the interlayer zone is approximately 0.4 mm. Inside of the interlayer zone the metal concentrations were different from calcu-
lated ones. Figs. 3(a and b) and 4 show the prime (␣ ) martensite formed in the melted zone by the rapid cooling during fabrication. This fine needle-like acicular formation exhibits a hexagonal close-packed crystal structure and demonstrates high hardness; on the other hand, it has low ductility and toughness [15,16]. As a result cracks may appear during the sample fabrication by DMD technique. However, for electrochemical investigations only samples without cracks were used. Titanium fabricated by DMD exhibited hardness (458 ± 10 HV) and lower corrosion rate (2.5 × 10−6 mm/year) comparing to titanium made by traditional melting (from 120 to 240 HV, 2.5 × 10—6 mm/year [2,13]). The results obtained from OCP–time measurements and cyclic polarization studies for alpha Ti–xAl alloys are shown in Table 1 and in Figs. 5 and 6 and Fig. 9. The data indicate that high aluminum content weakens corrosion resistance of alloys. As seen in Fig. 5, the OCP of the Ti–3Al alloy shifts
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Fig. 4. Scanning electron micrograph of two layers and interlayer heat affected zone.
towards noble direction over time (−0.7 to −0.18 V) more than OCP of the Ti–27Al alloy does (−0.62 to −0.42 V). The anodic polarization curves of the Ti–xAl alloys support that conclusion. The shape of the curve c in Fig. 6 indicates that Ti–27Al alloy does not form a stable passive layer. Having highest Al content of all studied alloys, it dissolves in physiological solution showing rapid increase in the current density near 0 V. Such behavior is expected given high activity of Al and large amount of it in the alloy. Fig. 6 also shows a polarization curve of pure titanium fabricated by DMD technique (Fig. 6(a)). Regions of active corrosion, passivity and transpassivation are easily identified. The mechanism of electrochemical corrosion of pure titanium
is well studied [2,17–21]. Our results indicate that overall electrochemical behavior of pure titanium made by DMD technique is similar to that of CP titanium made by traditional melting. However, the passive current density, corrosion rate and all major potentials (Table 1) show that pure titanium specimen made by DMD has higher corrosion resistance than that of Ti made by traditional melting. Based on obtained data (microstructural investigations, electrochemical testing in Hanks solution, etc.) we believed that the mechanisms of the pure titanium corrosion should be related to the martensitic phase transformation because of the rapid solidification and surface remelting during DMD fabrication. However, elucidating the difference in corrosion mechanisms between Ti made traditionally and DMD fabricated still needs further investigation.
Fig. 5. Open circuit potential versus time for titanium and its alloys in Hank’s solution: (a) pure Ti, (b) Ti–3Al, (c) Ti–27Al, (d) Ti–9Fe, (e) Ti–25Fe and (f) Ti–7Al–4Fe.
Fig. 6. Polarization curves of alloys and Ti in Hank’s solution: (a) pure Ti, (b) Ti–3Al, (c) Ti–27Al, (d) Ti–9Fe, (e) Ti–25Fe and (f) Ti–7Al–4Fe. The peaks of active dissolution: 1–3.
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3.2. Ti–yFe alloys
3.3. Ti–xAl–yFe alloys
The microstructure of the Ti–yFe alloys (y = 9, 14 wt.%) is shown in Fig. 3(c and d). The Ti–yFe specimens possess a basket-weave type of microstructure. The OCP and passive current density values obtained for samples of the Ti–yFe alloys are listed in Table 1. The OCP–time curves of the Ti–yFe alloys (y = 9, 25 wt.%) are presented in Fig. 5(d and e). Iron is considered to be a betastabilizer [13]. Therefore, increase in iron content from 9 to 25 wt.% shifts the OCP values from −0.35 to −0.62 V. The Ti–25Fe alloy shows different OCP decay curve compared to other alloys (Fig. 5), which can be explained by large amount of iron on the surface of the alloy. The passive film of the Ti–25Fe alloy includes iron hydroxide complexes in addition to titanium-based oxides. In the beginning of the experiment large amount of iron dissolves into electrolyte forming the passive film according to electrochemical process described by Kolotirkin et al. [25]. As a result, the OCP value of the Ti–25Al alloy decreases over time. Thus, the shape of the OCP–time curve is determined by the electroreduction of iron(III) oxide. As seen from Fig. 3, the Ti–yFe alloys exhibit -type of structure, and therefore, show superior corrosion resistance comparing to the Ti–xAl and Ti–xAl–yFe alloys which possess ␣ and ␣ +  types of structure. However, corrosion resistance of Ti–yFe alloys also depends on several factors, such as composition, environment and microstructure [22]. Tomashov reported [23] that addition of alloying elements, such as Al, V, Mo and Fe to titanium increases the steady dissolution currents in the passive state. That happens because alloying atoms entering in the titanium lattice change the defectness degree of passive films. Small amounts of iron additive in titanium-based alloy improve its corrosion resistance. All Ti–yFe alloys have the passive region. The Ti–9Fe alloy has the minimum passive current density (6.3 × 10−5 A/m2 ). Its electrochemical parameters are very close to that of pure titanium. On the other hand, increase in the iron concentration up to 25 wt.% tremendously decreases the corrosion resistance due to formation of new phases of Ti–Fe compositions which destroy the passivating titanium dioxide surface film. Both OCP (−0.62 V) and passive current density (0.3 A/m2 ) of the Ti–25Fe alloy characterize the alloy as less resistant to corrosion than the Ti–9Al alloy. Fig. 6(e) shows the typical potentiodynamic polarization curve of the Ti–25Fe alloy in Hanks solution. Three anodic current peaks at −0.3 V (peak 1), at −0.1 V (peak 2) and at 0.25 V (peak 3) exist in the active zone. The results can be attributed to the effect of the iron in the alloy. First peak can be assigned to the electroformation of a pre-passive Fe(OH)2 layer and consecutive dissolution of iron through this layer. The Fe(OH)2 layer is later electrooxidized to FeOOH in the potential range of peak 2. Similar phenomena have been reported previously [26]. Finally, peak 3 corresponds to the electroformation of the hydrous FeOOH/Fe(OH)2 films.
The microstructure of the Ti–xAl–yFe alloys (x, y: 5, 8 and 7, 4, respectively) is shown in Fig. 3(e and f). Three different phases, ␣,  and hexagonal martensite ␣ , are present. The electrochemical behavior of the Ti–xAl–yFe alloys (0 ≤ x ≤ 27 wt.%, 0 ≤ y ≤ 25 wt.%) was studied. Figs. 5(f) and 6(f) show the typical OCP–time and potentiodynamic polarization curves obtained for the Ti–7Al–4Fe alloy in Hank’s solution. The shape of curves was typical for all Ti–xAl–yFe alloys studied in this work. Fig. 7 shows the effect of composition on the corrosion rate in Hank’s solution. The anodic polarization curve reveals that the Ti–7Al–4Fe alloy is self-passivated and the passive films are stable at up to +1.8 V. The anodic current density slightly increases because of oxidation reaction of the soluble ions Ti4+ to Ti6+ [2]. The second stable passive condition was found for the Ti–xAl–yFe alloys with low concentration of additives at the anodic potential higher than 2.2 V. The second active dissolution of passivating films started at potentials higher than 2.6 V. When total concentration of additives (both Ti and Al) was less then 15 wt.% the OCP was shifted towards the noble direction and alloy was spontaneously passive in Hank’s solution. The oxide layer of this specimen exhibits the highest OCP value compared to all other Ti–xAl–yFe samples with higher amounts of additives. The presence of the three phases, ␣, ␣ and , with even distribution of alloying elements, apparently improves the corrosion behavior. The alloying elements are distributed uniformly; therefore, oxide clusters also get distributed uniformly in the matrix and form a stable passive layer. On the other hand, the high amount of Fe and Al causes formation of Ti3 Al, TiAl and Tix Fey that destroy the TiO2 passivating film. All of the Ti–xAl–yFe alloys exhibited active-to-passive transitions in Hank’s solution. The minimum value of
Fig. 7. The composition effect on the corrosion rate in Hank’s solution.
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passive current density (6.3 × 10−3 A/m2 ) was found for the Ti–7Al–4Fe alloy. The value is 263-fold higher than that of pure titanium. The corrosion rate of Ti–7Al–4Fe alloy was 10 folds slower comparing to materials made by traditional melting: the pure titanium [2,13], Ti–5Al–4V [2] and Ti–6Al–4Fe [2]. Our results support the explanation of titanium corrosion in aqueous salts offered by Popa et al. [2]. The physiological medium contains Cl− ion as well as many others. Popa et al. [2] studied the effect of different salts on the electrochemical behavior of Ti–5Al–4V and Ti–6Al–4Fe in bioliquids and proposed the following explanation. Dissolution at passivation includes two steps: the hydrolysis of the titanium
dioxide and the transport of dissolution products to the bulk electrolyte. The second step is rate determining one. The dissolution products are neutral species, such as Ti(OH)2 or hydroxocomplex, such as TiO(OH)2 [24]. This is consistent with a decrease of corrosion resistance in Hank’s solution, which contains many salts comparing to a monoanionic NaCl solution. Resistance to pitting corrosion was studied as well. After corrosion test, the surfaces of the samples were observed using both optical and scanning electron microscopes. Pitting corrosion was found in several alloys: Ti–25Fe, Ti–18Al–7Fe, Ti–24Al–5Fe, Ti–21Al–5Fe, Ti–17Al–3Fe and Ti–27Al. The typical pits are shown in Fig. 8. The cyclic voltammetry was performed on all fabricated Ti-based samples (Fig. 9). The pitting potential and protection values are presented in Table 1. The CP Ti and Ti–3Al alloy show the highest values of the protection potentials. The Ti-based alloys with high amounts of additives have low values of pitting and protection potentials because of weak and heterogeneous passive films. Pitting corrosion is a complex process on a passivated metal surface, which leads to a localized attack. Rapid dissolution of small discrete areas may occur in four steps. These steps generally are: breakdown of passivity, initiation, propagation and repassivation. Pitting corrosion is dependent on the metal or the alloy composition as well as on microstructure. Sun et al. [16] studied the pitting process on the surface of titanium alloys after laser remelting. They showed that the probability of pitting initiation and propagation on remelted titanium sheet is very low in spite of the presence of aggressive ions, such as Cl− , that usually are essential to breakdown of the passive film and initiation of localized corrosion. Remelting of the alloys during fabrication results
Fig. 8. Scanning electron micrographs of pits on the surface of Ti–18Al–7Fe (a and b), Ti–25 Fe (c) alloys after 10 h exposure to Hank’s solution.
Fig. 9. Cyclic voltammograms for (a) Ti, (b) Ti–3Al, (c) Ti–9Fe and (d) Ti–7Al–4Fe in Hank’s solution.
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in increased resistance to pitting corrosion. No pits have been found on the surface of the Ti–xAl–yFe alloys with low amount of additives. When the concentration of either Al or Fe increases significantly pitting corrosion starts to be a major corrosion mechanism. This may be attributed to the presence of large number of beta grains and uneven distribution of the alloying elements in various phases.
4. Conclusions Considering all the data shown the following conclusions can be made: 1. The electrochemical and corrosion behavior of the Ti–xAl–yFe alloys fabricated by DMD method was studied in Hank’s solution. The Ti–xAl–yFe alloys exhibited spontaneous passivity in the artificial physiological media. 2. The microstructure of each fabricated alloy was studied using SEM method. The fraction of each phase is determined by the amount of the corresponding stabilizer, either Al or Fe and by the fabrication regime. 3. The electrochemical investigation shows that small amount of iron in titanium-based alloy improves the corrosion resistance of the alloy in Hank’s solution. Noble OCP and low passive current density occur due to the equixed microstructure. However, the passive current density of the Ti–yFe alloy (x > 15 wt.%) increased due to the formation of Ti–Fe components. 4. The ␣ + Ti–xAl–yFe alloys with low content of additives exhibited noble OCP and low passive current density comparing to the samples with high amount of alloying metals. The presence of three phases ␣, ␣ and , with even distribution of alloying elements, apparently improves the corrosion behavior of the Ti–xAl–yFe alloys with low amount of additives. High concentrations of Al or Fe caused uneven distribution of alloying elements in the phases, which resulted in poor corrosion resistance in biofluids. Acknowledgements The authors gratefully acknowledge Mr. Joseph W. Vickers, Laser Development Engineer, Rapid Prototyping Center, University of Louisville, KY, for the support in fabrication of the Ti-based alloys.
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Electrochemical and corrosion behavior of Ti–xAl–yFe alloys prepared by direct metal deposition method Natalia V. Pimenova, Thomas L. Starr ∗ Chemical Engineering Department, J.B. Speed School of Engineering, University of Louisville, Louisville, KY 40292, USA Received 29 November 2004; received in revised form 7 July 2005; accepted 7 July 2005 Available online 25 August 2005
Abstract The electrochemical and corrosion behavior of Ti-based alloys was investigated. The direct metal deposition technique was used to fabricate 21 alloys with different ratio of metals (0 ≤ Al ≤ 27 wt.%, 0 ≤ Fe ≤ 25 wt.%). Corrosion resistance of each alloy was evaluated both qualitatively and quantitatively by voltammetric measurements in the simulated human body fluid conditions (Hank’s solution). The corrosion rates of the materials were compared in Hank’s solution using Tafel extrapolation method. Among the Ti–xAl–yFe alloys the Ti–7Al–4Fe alloy exhibited the slowest corrosion rate of 7.7 × 10−4 mm/year and the least value of passive current density (6.3 × 10−3 A/m2 ). The alloy is resistant to pitting corrosion as well. © 2005 Elsevier Ltd. All rights reserved. Keywords: Titanium-based alloy; Polarization; Passivity; Pitting corrosion; Direct metal disposition
1. Introduction Titanium demonstrates an excellent corrosion resistance for long periods in the physiological fluids. The titanium dioxide film (TiO2 ) formed on its surface instantly and spontaneously in the presence of oxygen and moisture is extremely stable and protective. A damaged oxide film can usually recuperate itself relatively easily. Due to its excellent biocompatibility, mechanical properties and corrosion resistance in aggressive media titanium is the most important candidate for use as human body implant material. Although pure titanium offers better corrosion resistance and tissue tolerance, than all other implant materials [1–3], its comparatively low strength restricts its usage to certain application, such as pacemaker cases and reconstruction devices [4,5]. When aluminum and vanadium are added to titanium in small quantities, the strength of the alloy is much increased compared with that of titanium [1]. For example, the Ti–6Al–4V alloy demonstrates higher strength and lower ∗
Corresponding author. Tel.: +1 502 852 1073; fax: +1 502 852 1577. E-mail address: [email protected] (T.L. Starr).
0013-4686/$ – see front matter © 2005 Elsevier Ltd. All rights reserved. doi:10.1016/j.electacta.2005.07.024
elastic modulus than that of pure titanium. The alloy is being used as implant material in high stress-bearing situations, such as hip and knee prostheses, as well as trauma fixation devices for the last two decades. However, since the early 1980’s the use of the alloy has been questioned because of black debris containing high levels of titanium, vanadium and aluminum that are found in surrounding tissue under conditions of high wear, such as in knee and hip implants [4]. Although no toxic effects have been linked to this black debris, safety concerns about the vanadium and aluminum have been raised in literature [2,6–8]. Aluminum [2,6] is poorly absorbed within the gastrointestinal tract. Very little amount gets into the blood stream, and hence it appears nontoxic. Vanadium [7,8] can alter the kinetics of the enzyme activity associated with cells of the inflammatory response. Vanadium contained in this alloy has been associated with potential cytotoxic effects and adverse tissue reactions [2]. The titanium-based alloy with aluminum and iron appears to be more suitable for implant applications. The Ti–5Al–2.5Fe alloy possesses similar corrosion resistance and mechanical properties to those of the Ti–6Al–4V alloy. More importantly, the iron-containing alloy has no apparent toxicity. In soft tissue, aluminum and iron show the reaction
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of “sequestration”, namely the formation of connective tissue membrane around the implant [9–12]. The surgical implant alloys are used for a long time (∼20 years). Therefore, the hydrolysis of different compounds of the passive film is possible. The toxicity of dissolved metal ions may cause pain and restrict action in the tissues adjoining the implants. In order to avoid such situation, the corrosion resistance of bioimplants needs to be high. The aims of the present work are: developing a new technique for discovery of the state-of-the-art materials; making combinatorial libraries of the Ti–Al–Fe alloys of different composition by direct metal deposition (DMD) and comparing corrosion resistance of the alloys in Hank’s solution with that of titanium fabricated by DMD. Fig. 1. Geometry of seven layers of coupon of Ti–xAl–yFe alloy (0 < x < 30 wt.%, 0 < y < 30 wt.%).
2. Experimental 2.1. Materials The Ti–xAl–yFe alloys (0 ≤ x ≤ 27 wt.%, 0 ≤ y ≤ 25 wt.%) were made by DMD technique. DMD 3000 machine (The POM Group, Auburn Hills) at the University of Louisville fabricated 3D-coupons by layer-by-layer deposition utilizing a 3000 W CO2 laser to melt metal powders injected by nozzle and laying down clad tracks on a substrate via a CNC position system. The powder feedstock used for depositing the alloys consisted of blends of pure elemental powders of titanium (Crucible Research, blend B344), aluminum (Praxair, Al-104) and the iron alloy H131 (Intern Mold Steel Inc., the primary source of iron). The Ti–6Al–4V alloy was used as a substrate (Principal Metals Corporation, manufactured by RMI Titanium Company, ASTM B 265 standard, grade 5). The laser forming process and corresponding coordinate system are illustrated in Fig. 1. Three sets of laser formed test coupons were nominally 12 mm wide (x), 14 mm tall (z) and 24 mm long (y). Each coupon was composed of seven layers of the Ti–xAl–yFe alloys (0 ≤ x ≤ 27 wt.%, 0 ≤ y ≤ 25 wt.%). The thickness of each layer was about 2 mm. The composition of each Ti-based alloy is shown in Table 1. Multilayer materials named combinatorial libraries were fabricated for the discovery of the new Ti–xAl–yFe alloys with high resistance to electrochemical corrosion in biofluids. DMD technique allows fabricating seven different compositions in one coupon. Applying combinatorial approach allows making the fabrication of the Ti–xAl–yFe alloys more time and costing effective. For metallographic investigation samples were cut along the x–z plane from the coupon to examine the as-deposited macro- and microstructures. The x–z samples were sectioned in the x–y plane. Sectioning was performed using an Isomet series low speed saw. 1 International Mold Steel’s Premium H-13 is a Cr-based, hot work die and plastic mold steel with the composition: 0.4 wt.% C, 0.4 wt.% Mg, 1 wt.% Si, 5.25 wt.% Cr, 1.35 wt.% Mo, 1 wt.% V and the rest Fe.
The electrochemical investigation was performed for each composition. When the coupon was cut layer-by-layer each slice was mounted in an epoxy resin. The samples were polished sequentially with 120-, 240-, 400- and 600-grit silicon carbide abrasive papers and then with 6 m diamond paste. To prevent oxidation of the surface and to degrease the samples, they were washed in 10% HCl solution and then in ethanol. The composition of each layer was homogeneous. The concentration was different on the boundary of the layers because of diffusion at remelting during the fabrication. Remelting allows making the uniform composition through each layer. The data of energy dispersive X-ray fluorescence analyzer (EDAX® ) linescan (Fig. 2) shows the 0.4 mm interlayer zone between the layers. The composition of interlayer zone is different from the regular composition of each layer. During pre-treatment this layer was removed and only regular composition material was used in the electrochemical experiments.
Fig. 2. EDAX® linescan data of the seven layers coupon (set 3).
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Table 1 Electrochemical parameters of Ti–Al–Fe alloys of different composition No.
Composition
Ecorr (V) vs. Ag/AgCl
icorr (A/m2 )
rcorr (mm/year)
ipas (A/m2 )
Epit (V)
Eprot (V)
10-1 8-1 10-6 10-7 8-2 6-2 10-3 10-2 6-7 10-4 8-4 6-6 8-6 8-7 8-5 6-3 8-3 10-5 6-4 6-5
Pure Ti Ti–3Al Ti–4Al Ti–7Al Ti–18Al Ti–27Al Ti–9Fe Ti–14Fe Ti–25Fe Ti–5Al–8Fe Ti–10Al–8Fe Ti–12Al–8Fe Ti–21Al–5Fe Ti–24Al–5Fe Ti–18Al–1Fe Ti–18Al–7Fe Ti–8Al–10Fe Ti–7Al–4Fe Ti–17Al–3Fe Ti–15Al–6Fe
−0.50 −0.10 −0.30 −0.40 −0.40 −0.42 −0.35 −0.30 −0.62 −0.30 −0.50 −0.24 −0.55 −0.62 −0.54 −0.54 −0.36 −0.38 −0.11 0.05
2.5E−06 5.0E−03 8.5E−03 2.0E−02 4.6E−02 5.5E−02 2.2E−05 2.0E−03 2.0E−02 1.6E−03 1.8E−02 3.0E−02 4.5E−03 5.0E−02 5.0E−02 8.0E−03 1.2E−03 7E−04 1.5E−02 1.0E−02
2.9E−06 5.9E−03 1.0E−03 2.4E−02 4.8E−2 6.5E−02 2.4E−05 2.1E−03 2.0E−02 1.8E−03 2.1E−02 3.7E−02 5.5E−03 6.3E−02 6.0E−02 1.0E−02 1.4E−03 7.7E−04 1.9E−02 1.2E−02
2.1E−05 3.0E−01 – – – – 6.3E−05 4.0E−02 3.0E−01 2.0E−02 – 2.0E−01 7.0E−01 – 3.0E+00 5.0E−01 7.5E+00 6.3E−03 2.1E+00 6.5E+00
2.21 0.50 1.80 1.50 −0.38 −0.40 1.15 0.80 0.70 0.60 0.35 0.50 −0.50 −0.55 0.20 0.40 0.62 1.64 0.62 0.75
2.08 0.40 1.60 0.90 – – 0.80 0.65 0.40 0.42 0.20 0.35 – – 0.00 0.10 0.45 1.42 0.26 0.50
2.2. Optical and electron microscopies The microstructure was investigated by means of optical and scanning electron microscopy. The specimens were prepared following the metallographic techniques used for titanium and its alloys [13]. These consist of grinding up to 600-grit SiC and polishing with 6 m diamond paste to mirror finish. The specimens were cleaned with deionized water and etched in Kroll’s reagent (10 ml of HF, 5 ml of HNO3 and 85 ml of water) for microscopic analysis. Etching was performed by swabbing Kroll’s reagent to the polished surface. The etching time was limited to 10 s to avoid overetching. Microstructure was observed on Leco PME-3 inverted microscope. The image viewing accessories included an Olympus camera and Sony Trinitron viewing system. Microstructural and compositional investigations were performed by using the Jeol JSM-5310 Scanning Electron Microscope (SEM) with EDAX® equipped with sapphire Si(Li) liquid nitrogen cooled EDS detector. The minimum detectable concentration has been roughly estimated to be approximately 1 at.%.
Research 273A potentiostat at a scan rate of 5 mV/s. The potential range of the polarization scans was between −0.6 and 2.5 V2 . Open circuit potential (OCP) measurements were carried out for each alloy in Hank’s solution. The potential was measured for 30 min prior to cyclic polarization experiments. The corrosion current densities determined by linear polarization icorr = k(di/dE)Ecorr were calculated from Stern formula assuming k = 0.026 V [14]. These currents were used in order to calculate the corrosion rates. All measurements were performed in Hank’s solution (g/l): 0.185 CaCl2 , 0.4 KCl, 0.06 KH2 PO4 , 0.1 MgCl2 ·6H2 O, 0.1 MgSO4 ·7H2 O, 8.0 NaCl, 0.35 NaHCO3 , 0.48 Na2 HPO4 and 1.00 d-glucose. The temperature of Hanks solution was kept at 25 ± 1 ◦ C. Nitrogen gas was purged into the solution between the measurements to eliminate the influence of dissolved oxygen on the passive film kinetics. During measurements the solution was blanketed with nitrogen atmosphere.
3. Results and discussion 2.3. Electrochemical experiments A conventional three-electrode cell with a Pt wire gauze with a large surface area as a counter electrode and saturated silver/silver chloride as a reference electrode was used to conduct electrochemical experiments. The geometrical area of the working electrode was 0.8 cm2 . The distance between the Luggin capillary and the working electrode was 3 mm. Prior to the polarization scans the samples were exposed to test solutions for 12 h to attain equilibrium condition. The anodic potentiodynamic polarization and cyclic voltammetry measurements were performed on an EG&G Princeton Applied
3.1. Ti–xAl alloys Table 1 shows composition and electrochemical parameters of all investigated alloys. The micrographs of two alpha Ti–xAl alloys (x = 4, 7 wt.%) are shown in Fig. 3(a and b). 2 The highest value of polarization potential was different for studied alloys and depended on the corrosion resistance of the alloy. For example, the Ti–27Al alloy did not have the passivation region on the anodic polarization curve. The pure Ti had the highest corrosion resistance. Only at anodic potentials higher than 1.2 V the transpassive behavior appeared. For such high corrosion resistant alloy the highest potential was 2.5 V.
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Fig. 3. Scanning electron micrographs of Ti–xAl–yFe alloys: (a) ␣-Ti–4Al, (b) ␣-Ti–7Al, (c) -Ti–9Fe, (d) -Ti–14Fe, (e) ␣/ Ti–5Al–8Fe and (f) ␣/ Ti–7Al–4Fe.
Air-cooling of coupon layers after deposition results in a fine needle-like alpha-phase referred to as fine lamellar or acicular microstructure. All binary Ti–xAl alloys have dominant alpha-phase. When the concentration of aluminum increases the structure might form the brittle Ti3 Al phases in the Tibased matrix. Because DMD technique induces a very high cooling rate (1000–10,000 ◦ C/s) [15], each layer (except the top one) experienced remelting when next layer was fabricated. Sun et al. [16] studied the microstructural changes in pure titanium after laser remelting and established that single ␣-phase was changed to acicular martensitic product. In our case optical microscopy observations indicate dramatic microstructural changes in the melted zone. Fig. 4 shows microstructures of previous layer, interlayer heat affected zone, and the next melted zone. According to EDAX® linescan data the interlayer zone is approximately 0.4 mm. Inside of the interlayer zone the metal concentrations were different from calcu-
lated ones. Figs. 3(a and b) and 4 show the prime (␣ ) martensite formed in the melted zone by the rapid cooling during fabrication. This fine needle-like acicular formation exhibits a hexagonal close-packed crystal structure and demonstrates high hardness; on the other hand, it has low ductility and toughness [15,16]. As a result cracks may appear during the sample fabrication by DMD technique. However, for electrochemical investigations only samples without cracks were used. Titanium fabricated by DMD exhibited hardness (458 ± 10 HV) and lower corrosion rate (2.5 × 10−6 mm/year) comparing to titanium made by traditional melting (from 120 to 240 HV, 2.5 × 10—6 mm/year [2,13]). The results obtained from OCP–time measurements and cyclic polarization studies for alpha Ti–xAl alloys are shown in Table 1 and in Figs. 5 and 6 and Fig. 9. The data indicate that high aluminum content weakens corrosion resistance of alloys. As seen in Fig. 5, the OCP of the Ti–3Al alloy shifts
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Fig. 4. Scanning electron micrograph of two layers and interlayer heat affected zone.
towards noble direction over time (−0.7 to −0.18 V) more than OCP of the Ti–27Al alloy does (−0.62 to −0.42 V). The anodic polarization curves of the Ti–xAl alloys support that conclusion. The shape of the curve c in Fig. 6 indicates that Ti–27Al alloy does not form a stable passive layer. Having highest Al content of all studied alloys, it dissolves in physiological solution showing rapid increase in the current density near 0 V. Such behavior is expected given high activity of Al and large amount of it in the alloy. Fig. 6 also shows a polarization curve of pure titanium fabricated by DMD technique (Fig. 6(a)). Regions of active corrosion, passivity and transpassivation are easily identified. The mechanism of electrochemical corrosion of pure titanium
is well studied [2,17–21]. Our results indicate that overall electrochemical behavior of pure titanium made by DMD technique is similar to that of CP titanium made by traditional melting. However, the passive current density, corrosion rate and all major potentials (Table 1) show that pure titanium specimen made by DMD has higher corrosion resistance than that of Ti made by traditional melting. Based on obtained data (microstructural investigations, electrochemical testing in Hanks solution, etc.) we believed that the mechanisms of the pure titanium corrosion should be related to the martensitic phase transformation because of the rapid solidification and surface remelting during DMD fabrication. However, elucidating the difference in corrosion mechanisms between Ti made traditionally and DMD fabricated still needs further investigation.
Fig. 5. Open circuit potential versus time for titanium and its alloys in Hank’s solution: (a) pure Ti, (b) Ti–3Al, (c) Ti–27Al, (d) Ti–9Fe, (e) Ti–25Fe and (f) Ti–7Al–4Fe.
Fig. 6. Polarization curves of alloys and Ti in Hank’s solution: (a) pure Ti, (b) Ti–3Al, (c) Ti–27Al, (d) Ti–9Fe, (e) Ti–25Fe and (f) Ti–7Al–4Fe. The peaks of active dissolution: 1–3.
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3.2. Ti–yFe alloys
3.3. Ti–xAl–yFe alloys
The microstructure of the Ti–yFe alloys (y = 9, 14 wt.%) is shown in Fig. 3(c and d). The Ti–yFe specimens possess a basket-weave type of microstructure. The OCP and passive current density values obtained for samples of the Ti–yFe alloys are listed in Table 1. The OCP–time curves of the Ti–yFe alloys (y = 9, 25 wt.%) are presented in Fig. 5(d and e). Iron is considered to be a betastabilizer [13]. Therefore, increase in iron content from 9 to 25 wt.% shifts the OCP values from −0.35 to −0.62 V. The Ti–25Fe alloy shows different OCP decay curve compared to other alloys (Fig. 5), which can be explained by large amount of iron on the surface of the alloy. The passive film of the Ti–25Fe alloy includes iron hydroxide complexes in addition to titanium-based oxides. In the beginning of the experiment large amount of iron dissolves into electrolyte forming the passive film according to electrochemical process described by Kolotirkin et al. [25]. As a result, the OCP value of the Ti–25Al alloy decreases over time. Thus, the shape of the OCP–time curve is determined by the electroreduction of iron(III) oxide. As seen from Fig. 3, the Ti–yFe alloys exhibit -type of structure, and therefore, show superior corrosion resistance comparing to the Ti–xAl and Ti–xAl–yFe alloys which possess ␣ and ␣ +  types of structure. However, corrosion resistance of Ti–yFe alloys also depends on several factors, such as composition, environment and microstructure [22]. Tomashov reported [23] that addition of alloying elements, such as Al, V, Mo and Fe to titanium increases the steady dissolution currents in the passive state. That happens because alloying atoms entering in the titanium lattice change the defectness degree of passive films. Small amounts of iron additive in titanium-based alloy improve its corrosion resistance. All Ti–yFe alloys have the passive region. The Ti–9Fe alloy has the minimum passive current density (6.3 × 10−5 A/m2 ). Its electrochemical parameters are very close to that of pure titanium. On the other hand, increase in the iron concentration up to 25 wt.% tremendously decreases the corrosion resistance due to formation of new phases of Ti–Fe compositions which destroy the passivating titanium dioxide surface film. Both OCP (−0.62 V) and passive current density (0.3 A/m2 ) of the Ti–25Fe alloy characterize the alloy as less resistant to corrosion than the Ti–9Al alloy. Fig. 6(e) shows the typical potentiodynamic polarization curve of the Ti–25Fe alloy in Hanks solution. Three anodic current peaks at −0.3 V (peak 1), at −0.1 V (peak 2) and at 0.25 V (peak 3) exist in the active zone. The results can be attributed to the effect of the iron in the alloy. First peak can be assigned to the electroformation of a pre-passive Fe(OH)2 layer and consecutive dissolution of iron through this layer. The Fe(OH)2 layer is later electrooxidized to FeOOH in the potential range of peak 2. Similar phenomena have been reported previously [26]. Finally, peak 3 corresponds to the electroformation of the hydrous FeOOH/Fe(OH)2 films.
The microstructure of the Ti–xAl–yFe alloys (x, y: 5, 8 and 7, 4, respectively) is shown in Fig. 3(e and f). Three different phases, ␣,  and hexagonal martensite ␣ , are present. The electrochemical behavior of the Ti–xAl–yFe alloys (0 ≤ x ≤ 27 wt.%, 0 ≤ y ≤ 25 wt.%) was studied. Figs. 5(f) and 6(f) show the typical OCP–time and potentiodynamic polarization curves obtained for the Ti–7Al–4Fe alloy in Hank’s solution. The shape of curves was typical for all Ti–xAl–yFe alloys studied in this work. Fig. 7 shows the effect of composition on the corrosion rate in Hank’s solution. The anodic polarization curve reveals that the Ti–7Al–4Fe alloy is self-passivated and the passive films are stable at up to +1.8 V. The anodic current density slightly increases because of oxidation reaction of the soluble ions Ti4+ to Ti6+ [2]. The second stable passive condition was found for the Ti–xAl–yFe alloys with low concentration of additives at the anodic potential higher than 2.2 V. The second active dissolution of passivating films started at potentials higher than 2.6 V. When total concentration of additives (both Ti and Al) was less then 15 wt.% the OCP was shifted towards the noble direction and alloy was spontaneously passive in Hank’s solution. The oxide layer of this specimen exhibits the highest OCP value compared to all other Ti–xAl–yFe samples with higher amounts of additives. The presence of the three phases, ␣, ␣ and , with even distribution of alloying elements, apparently improves the corrosion behavior. The alloying elements are distributed uniformly; therefore, oxide clusters also get distributed uniformly in the matrix and form a stable passive layer. On the other hand, the high amount of Fe and Al causes formation of Ti3 Al, TiAl and Tix Fey that destroy the TiO2 passivating film. All of the Ti–xAl–yFe alloys exhibited active-to-passive transitions in Hank’s solution. The minimum value of
Fig. 7. The composition effect on the corrosion rate in Hank’s solution.
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passive current density (6.3 × 10−3 A/m2 ) was found for the Ti–7Al–4Fe alloy. The value is 263-fold higher than that of pure titanium. The corrosion rate of Ti–7Al–4Fe alloy was 10 folds slower comparing to materials made by traditional melting: the pure titanium [2,13], Ti–5Al–4V [2] and Ti–6Al–4Fe [2]. Our results support the explanation of titanium corrosion in aqueous salts offered by Popa et al. [2]. The physiological medium contains Cl− ion as well as many others. Popa et al. [2] studied the effect of different salts on the electrochemical behavior of Ti–5Al–4V and Ti–6Al–4Fe in bioliquids and proposed the following explanation. Dissolution at passivation includes two steps: the hydrolysis of the titanium
dioxide and the transport of dissolution products to the bulk electrolyte. The second step is rate determining one. The dissolution products are neutral species, such as Ti(OH)2 or hydroxocomplex, such as TiO(OH)2 [24]. This is consistent with a decrease of corrosion resistance in Hank’s solution, which contains many salts comparing to a monoanionic NaCl solution. Resistance to pitting corrosion was studied as well. After corrosion test, the surfaces of the samples were observed using both optical and scanning electron microscopes. Pitting corrosion was found in several alloys: Ti–25Fe, Ti–18Al–7Fe, Ti–24Al–5Fe, Ti–21Al–5Fe, Ti–17Al–3Fe and Ti–27Al. The typical pits are shown in Fig. 8. The cyclic voltammetry was performed on all fabricated Ti-based samples (Fig. 9). The pitting potential and protection values are presented in Table 1. The CP Ti and Ti–3Al alloy show the highest values of the protection potentials. The Ti-based alloys with high amounts of additives have low values of pitting and protection potentials because of weak and heterogeneous passive films. Pitting corrosion is a complex process on a passivated metal surface, which leads to a localized attack. Rapid dissolution of small discrete areas may occur in four steps. These steps generally are: breakdown of passivity, initiation, propagation and repassivation. Pitting corrosion is dependent on the metal or the alloy composition as well as on microstructure. Sun et al. [16] studied the pitting process on the surface of titanium alloys after laser remelting. They showed that the probability of pitting initiation and propagation on remelted titanium sheet is very low in spite of the presence of aggressive ions, such as Cl− , that usually are essential to breakdown of the passive film and initiation of localized corrosion. Remelting of the alloys during fabrication results
Fig. 8. Scanning electron micrographs of pits on the surface of Ti–18Al–7Fe (a and b), Ti–25 Fe (c) alloys after 10 h exposure to Hank’s solution.
Fig. 9. Cyclic voltammograms for (a) Ti, (b) Ti–3Al, (c) Ti–9Fe and (d) Ti–7Al–4Fe in Hank’s solution.
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in increased resistance to pitting corrosion. No pits have been found on the surface of the Ti–xAl–yFe alloys with low amount of additives. When the concentration of either Al or Fe increases significantly pitting corrosion starts to be a major corrosion mechanism. This may be attributed to the presence of large number of beta grains and uneven distribution of the alloying elements in various phases.
4. Conclusions Considering all the data shown the following conclusions can be made: 1. The electrochemical and corrosion behavior of the Ti–xAl–yFe alloys fabricated by DMD method was studied in Hank’s solution. The Ti–xAl–yFe alloys exhibited spontaneous passivity in the artificial physiological media. 2. The microstructure of each fabricated alloy was studied using SEM method. The fraction of each phase is determined by the amount of the corresponding stabilizer, either Al or Fe and by the fabrication regime. 3. The electrochemical investigation shows that small amount of iron in titanium-based alloy improves the corrosion resistance of the alloy in Hank’s solution. Noble OCP and low passive current density occur due to the equixed microstructure. However, the passive current density of the Ti–yFe alloy (x > 15 wt.%) increased due to the formation of Ti–Fe components. 4. The ␣ + Ti–xAl–yFe alloys with low content of additives exhibited noble OCP and low passive current density comparing to the samples with high amount of alloying metals. The presence of three phases ␣, ␣ and , with even distribution of alloying elements, apparently improves the corrosion behavior of the Ti–xAl–yFe alloys with low amount of additives. High concentrations of Al or Fe caused uneven distribution of alloying elements in the phases, which resulted in poor corrosion resistance in biofluids. Acknowledgements The authors gratefully acknowledge Mr. Joseph W. Vickers, Laser Development Engineer, Rapid Prototyping Center, University of Louisville, KY, for the support in fabrication of the Ti-based alloys.
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