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Effects of Sn addition on the microstructure, mechanical properties andcorrosion behavior of Ti–Nb–Sn alloys

Paulo E.L. Moraes, Rodrigo J. Contieri, Eder S.N. Lopes, Alain Robin,Rubens Caram

PII: S1044-5803(14)00255-1DOI: doi: 10.1016/j.matchar.2014.08.014Reference: MTL 7666

To appear in: Materials Characterization

Received date: 15 May 2014Revised date: 25 July 2014Accepted date: 14 August 2014

Please cite this article as: Moraes Paulo E.L., Contieri Rodrigo J., Lopes Eder S.N.,Robin Alain, Caram Rubens, Effects of Sn addition on the microstructure, mechanicalproperties and corrosion behavior of Ti–Nb–Sn alloys, Materials Characterization (2014),doi: 10.1016/j.matchar.2014.08.014

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Effects of Sn Addition on the Microstructure, Mechanical Properties and Corrosion

Behavior of Ti-Nb-Sn Alloys

Paulo E.L. Moraes1, Rodrigo J. Contieri2, Eder S.N. Lopes3, Alain Robin4, Rubens Caram5

1University of Campinas, School of Mechanical Engineering, Rua Mendeleiev, 200, Campinas, SP 13083-860, Brazil, Email: pauloeduardo.leitedemoraes@gmail.com 2University of Campinas, School of Mechanical Engineering, Rua Mendeleiev, 200, Campinas, SP 13083-860, Brazil, Email: contieri@fem.unicamp.br 3University of Campinas, School of Mechanical Engineering, Rua Mendeleiev, 200, Campinas, SP 13083-860, Brazil, Email: ederlopes@fem.unicamp.br 4University of São Paulo, School of Engineering of Lorena, Polo Urbo-Industrial Gleba AI-6, Lorena, SP, 12600-00, Brazil, Email: alain@demar.eel.usp.br 5(Corresponding author) University of Campinas, School of Mechanical Engineering, Rua Mendeleiev, 200, Campinas, SP 13083-860, Brazil, Phone: +55-19-35213314 / Fax: +55-19-32893722 Email: caram@fem.unicamp.br

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Abstract

Ti and Ti alloys are widely used in restorative surgery because of their good

biocompatibility, enhanced mechanical behavior and high corrosion resistance in

physiological media. The corrosion resistance of Ti-based materials is due to the

spontaneous formation of the TiO2 oxide film on their surface, which exhibits elevated

stability in biological fluids. Ti-Nb alloys, depending on the composition and the processing

routes to which the alloys are subjected, have high mechanical strength combined with low

elastic modulus. The addition of Sn to Ti-Nb alloys allows the phase transformations to be

controlled, particularly the precipitation of ω phase. The aim of this study is to discuss the

microstructure, mechanical properties and corrosion behavior of cast Ti-Nb alloys to which

Sn has been added. Samples were centrifugally cast in a copper mold, and the

microstructure was characterized using optical microscopy, scanning electron microscopy

and X-ray diffractometry. The mechanical behavior evaluation was performed using

Berkovich nanoindentation, Vickers hardness and compression tests. The corrosion

behavior was evaluated in Ringer’s solution at room temperature using electrochemical

techniques. The results obtained suggested that the physical, mechanical and chemical

behaviors of the Ti-Nb-Sn alloys are directly dependent on the Sn content.

Keywords: Ti alloys, X-ray diffraction, nanoindentation, microstructure, mechanical

behavior, corrosion

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1. Introduction

A crucial challenge for the development of metallic orthopedic implants is to obtain a

material that combines high mechanical strength, unique fatigue behavior, enhanced

biocompatibility and high biocorrosion resistance [1]. In addition to enhanced strength,

artificial biomaterials employed in orthopedic implants for repairing hip and knee joints

must exhibit reduced elastic modulus to prevent insufficient loading of bone adjacent to the

implant, also defined as bone stress-shielding [2]. Bone stress-shielding is due to the

difference in elastic moduli of the implant material and the bone. This phenomenon can

cause bone mass loss and occasionally, bone failure [3]. This set of material properties

can be observed for some specific titanium alloys [4].

Recently, interest in the development of metallic biomaterials employed in orthopedic

devices has been concentrated on the metastable Ti alloys that can be produced using

biocompatible elements, especially Nb, as the alloying element [1,2]. The microstructure

and mechanical behavior of metastable Ti alloys result from the amount and type of

alloying elements added to titanium and especially from the processing pathway to which

the alloy is subjected [5]. Metastable Ti alloys are found in the Ti-Nb system when the

amount of stabilizer is sufficient to avoid martensite formation during cooling from

temperatures in the phase field [6]. When a metastable Ti-Nb alloy is subjected to β

solution heat treatment followed by rapid cooling to room temperature, a microstructure

formed by β phase will be obtained. β phase in Ti-Nb alloys is characterized by presenting

reduced mechanical strength and low elastic modulus [7]. Metastable Ti alloys with

elastic modulus as low as 50 GPa have been reported in the literature [8].

In Ti alloys, the type of phase and its volume fraction are directly related to the

mechanical behavior. As a result, improvements in the mechanical behavior of metastable

Ti alloys are achieved by forcing the retained β phase to decompose into more stable

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phases as a result of aging heat treatments at intermediate temperatures. This procedure

results in phase precipitation, which is stronger and more rigid than the parent phase.

Stabilization of phase at room temperature is frequently followed by precipitation of the

detrimental phase [9], which causes the severe brittleness of Ti alloys. While the

phase is assumed to be a transition phase between α and β phases, it is hard, has high

elastic modulus and its precipitates have nanometric characteristics. Additionally, phase

is supposed to act as the nucleation substrate in the phase matrix during phase

nucleation [10]. Thus, to improve the mechanical behavior of metastable Ti alloys, it is

essential for phase precipitation to be well controlled.

An effective means of preventing excessive phase precipitation in metastable β Ti

alloys is the use of Sn as an alloying element [11]. In terms of metallic materials for

orthopedic applications, another imperative question is the corrosion phenomenon

originated by the reaction of the implanted device with body physiological fluids.

Frequently, the corrosion processes in implanted orthopedic devices can lead to ion

release, which is harmful to the human body. In the literature, there are very few studies

on the effect of Sn addition on the corrosion behavior of Ti-Nb alloys [12-13]. An effective

method of investigating the corrosion of metallic alloys is the use of electrochemical

techniques that can describe corrosion evolution and oxide layer growth.

Therefore, the aim of this work was to investigate the microstructure and the

mechanical properties of Ti-Nb-Sn alloys as a function of the Sn content and to discuss the

alloys’ corrosion behavior in Ringer’s solution at room temperature using electrochemical

techniques.

2. Experimental Details

The samples investigated were prepared in an arc furnace with a non-consumable

tungsten electrode and water-cooled copper crucible under argon atmosphere. Ingots of

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30 to 50 g were prepared using high chemical purity Ti, Nb and Sn. Samples were

homogenized by inverting and re-melting them at least five times. Six compositions in the

Ti-Nb-Sn system were chosen: Ti-30Nb, Ti-30Nb-2Sn, Ti-30Nb-4Sn, Ti-30Nb-6Sn, Ti-

30Nb-8Sn and Ti-30Nb-10Sn (wt %). Afterward, these samples were arc-melted and cast

in a permanent copper mold at room temperature using centrifugal casting equipment.

Samples were subjected to a maximum rotation of 1,000 rpm. The size of the cast

samples was 10 x 10 x 4 mm.

The chemical composition of the arc-cast samples was determined using X-ray

fluorescence spectroscopy (Rigaku RIX3100). Interstitial content was measured using

LECO 400 equipment. Sample microstructure features were evaluated by employing

standard metallographic procedures. After grinding and polishing, the samples were

etched using a solution of 5 ml HF, 30 ml HNO3 and 65 ml H2O. Evaluation of

microstructures was carried out using optical microscopy (Olympus BX60 M), scanning

electron microscopy (Zeiss Evo 15) and X-ray diffractometry (Panalytical X'pert PRO)

using CuK radiation, 40 kV and 30 mA.

Vickers hardness was measured using Buehler equipment with a load of 200 gf

applied for 15 s. The values were determined based on 10 different measurements. Elastic

constants including elastic moduli were measured using a pulse-echo acoustic emission

technique, as described by the ASTM E494 standard (2005), using a Panametrics-NDT

5072PR pulser-receiver equipped with transducers operating at a frequency of 5 MHz. In

addition, nanoindentation was applied to measure hardness and elastic modulus, which

were obtained from load-penetration depth curves using a Nano Hardness Tester (NHT-

CSM Instruments) equipped with a three-sided Berkovich diamond indenter and applying a

maximum load of 500 mN. Compression samples 4 mm high and 2 mm in diameter were

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machined by the electrical discharge machining . Compression tests were carried out

using an EMIC DL2000 testing machine at a strain rate of 8 x 10-3 s-1.

For the electrochemical tests, square-shaped specimens (10 mm x 10 mm) of Ti-Nb-

Sn alloys were used. Their surfaces were mechanically ground with emery-paper up to

1200 grit, rinsed with distilled water and dried. The simulated body fluid was Ringer’s

solution with the following composition: 8.6 g NaCl + 0.3 g KCl + 0.33 g CaCl2-2H20 in

1000 mL deionized water. The pH was approximately 7.5. The solution was naturally

aerated and the experiments were conducted without stirring. The experiments were

performed at room temperature. The counter electrode was a square-shaped platinum

sheet of 18 cm2 area. All potentials were referred to the saturated calomel electrode (SCE)

potential.

Open-circuit potential, electrochemical impedance (EIS) and polarization

measurements were performed using the Electrochemical Interface Solartron mod. 1287A

and the Frequency Response Analyzer Solartron mod. 1260 A, controlled by the

Ecorr/Zplot Solartron mod. 125587S software. Prior to polarization experiments, the

working electrodes were immersed in the Ringer’s solution for 3 hours, taking the moment

of immersion as the zero time point. Then, impedance measurements at open-circuit

potential (OCP) were made using a sinusoidal signal of 10 mV amplitude and frequencies

in the 0.01 Hz - 100 kHz range. Potentiodynamic polarization was then carried out using a

1 mV s-1 sweep rate from -0.5 V/OCP to +2.8 V/SCE. After each run, the samples were

reground with emery papers up to a 1200 grit finish to remove any product formed on the

metal surface which could affect the following tests, and the samples were rinsed with

distilled water and dried. All experiments were performed in triplicate and good

reproducibility was observed.

3. Results and Discussion

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3.1. Microstructure Features

This investigation included the processing and characterization of the Ti-30Nb alloy

(wt %) with additions of Sn. The chemical compositions of the samples that were arc-

melted and cast in copper mold are shown in Table 1. These results suggest that the

measured compositions are very close to the nominal compositions. The interstitial

(oxygen and nitrogen) contents are in good agreement with the ASTM-B-364-83 standard.

The oxygen content is in the range of 0.093 ± 0.03 wt % to 0.166 ± 0.05 wt %.

Taking into account that the solidification in copper mold led to rapid solidification and

cooling, it is expected that final microstructure would consist of metastable phases. It

should be mentioned that during the solidification, the latent heat release decreases the

cooling rate and it can affect the solid/solid phase transformations. According to literature

[14,15], solution and rapidly cooled β-Ti alloys can present martensitic transformations and

depending on the content of β stabilizing elements, the β phase can be retained at room

temperature. The addition of Sn to the Ti-30Nb alloy can also change the solidification

pathway by changing the liquidus and solidus temperatures and hence, the final

microstructure.

Figure 1 depicts optical images of the microstructure of the experimental Ti alloys. It

is clear that all the samples presented dendritic solidification. It is apparent that the

addition of Sn led to an increase in the dendritic arms, which could be related to changes

in the solidification range. Additionally, the dendritic growth produces local changes in

composition, which can be confirmed by the X-ray images (SEM) displayed in Figure 2 that

shows Nb, Sn and Ti distribution in the Ti-30Nb-6Sn alloy. These images suggest that the

Nb is segregated in the core of the dendritic arms whereas Sn accumulated between the

dendritic arms. Whereas Sn dissolved in Ti has a partition coefficient below the unit level,

Nb in Ti has a partition coefficient higher than one, which is a result of the high melting

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temperature of Nb. Figure 2.a also shows retained β grains and their respective grain

boundaries.

It is well known that the microstructure and hence, the mechanical behavior of Ti

alloys change with the alloy composition and the processing routes to which the samples

are subjected. Solution heat-treatment of Ti alloys at temperatures in the β phase field

followed by rapid cooling produces microstructures formed by metastable phases [16,17].

Samples cast in the copper mold are subjected to very high cooling rates, which can result

in martensitic transformation. In the case of high β stabilizing content, orthorhombic

matensite is formed (α” phase), whereas low β stabilizing content leads to hexagonal

matensite (α’ phase).

Figure 3 presents X-ray diffraction (XRD) patterns of the Ti-30Nb-XSn (X=0, 2, 4, 6,

8 and 10) wt %. The diffraction pattern of the Ti-30Nb alloy shows peaks related to the β

phase and the α” phase. In addition, evidence of phase precipitation is detected using

slow scan speed (Figure 4). The athermal ω phase could precipitate during quenching

from high temperatures. This occurs due to instabilities of the β phase and is a result of the

collapse of planes {111} producing a hexagonal crystal structure that is brittle and hard.

The diffraction pattern of the Ti-30Nb-2Sn alloy shows that the addition of Sn decreases

the amount of ω phase in the microstructure. As the amount of Sn increases, the

suppression of ω phase precipitation becomes more evident. It is apparent that the

addition of 4 wt % Sn reduced the ω phase precipitation to a value undetected by X-ray

diffraction, as observed in Figure 4, which shows overlapping peaks related to the {301}ω

and {112}ω planes. The addition of Sn also affects the martensite formation temperature.

As the Sn content reached 4 wt %, the volume fraction of martensite exhibited some

reduction, and when it reached 10 wt %, the full β stabilization occurred. Such information

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supports the conclusion that Sn combined with Nb acts as a β stabilizing element in Ti

alloys. Additionally, it was found that Sn can lower the martensitic start temperature.

The X-ray diffraction patterns obtained by the use of slow scanning speed also show

that the diffraction peaks of the β phase shift to smaller angles with Sn addition, which

indicates that the lattice parameter of the β phase increases with Sn addition.

3.2. Mechanical Behavior

The mechanical properties of Ti alloy are affected by the morphology and volume

fraction of phases in the microstructure. To evaluate the effect of Sn addition on the

mechanical behavior, elastic modulus, Vickers hardness and compression tests of the cast

sample were carried out.

Figure 5 shows the Vickers hardness of the alloys measured using conventional and

nanoindentation techniques as a function of the Sn content. These results show that there

is a strong relationship between the hardness and the Sn content. As the Sn was added to

the Ti-30Nb alloy, there was a clear decrease in the Vickers hardness, which could be

related to the suppression of ω phase precipitation and also due to the increase in the

lattice parameter of the phase. As the amount of Sn reached a value between 6 and 8 wt

%, the hardness values stabilized.

Figure 6 depicts the effect of the Sn content on the elastic modulus of the cast Ti-Nb-

Sn samples. Similar to the Vickers hardness behavior, the elastic modulus depends on the

interatomic forces. An increase in the interatomic distance results in a drop in the elastic

modulus. It is apparent that the addition of 6 wt % Sn reduces the elastic modulus value to

a minimum value.

Figure 7 shows stress versus strain curves obtained by compression tests. Again, the

lattice parameter changes because Sn addition influences the elastic/plastic transition. The

obtained results suggest that the Ti-30Nb alloy and the Ti-30Nb-10Sn alloy exhibit the

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highest yield strength values. This behavior is in good agreement with the hardness and

elastic modulus measurements. The Ti-30Nb alloy likely exhibits a higher strength due to

the phase precipitation. Conversely, the high strength of the Ti-30Nb-10Sn alloy is due

to a saturation of Sn. The X-ray diffraction analyses have not indicated the formation of

any intermetallic phase. Under compression, the Ti-30Nb and Ti-30Nb-10Sn alloys

exhibited maximum strength near 1.750 MPa. However, compared with the other alloys

(X=2, 4, 6 and 8) wt %, they presented brittle behavior and very low ductility. Figure 7 also

presents fractographic details of the Ti-30Nb and Ti-30Nb-6Sn alloys after the

compression tests. While the Ti-30Nb-6Sn alloy shows evidence of significant reduction in

diameter, the Ti-30Nb alloy exhibits 45o fractures that are typical of brittle materials. The

fracture mode of the Ti-30Nb alloy was transgranular, whereas the Ti-30Nb-6Sn alloy

exhibits dimples.

3.3. Electrochemical Evaluation

Figure 8 presents the variation of open-circuit potential (OCP) for the Ti-Nb-Sn alloys

as a function of exposure time in Ringer’s solution. All OCPs shift in the more noble

direction with time, which is indicative of the formation and growth of passivating films [18-

19]. After 3 h-immersion, the OCPs of all alloys have reached a near steady state and are

in very close proximity to each other. The OCP values are in the -0.361 to -0.308 V/SCE

range, which is in accordance with the values obtained for the Ti-16Nb-XSn (X=4.0, 4.5

and 5.0 wt %) alloys in NaCl 0.9% and Hank’s solution of pH 7.4 [12]. Nevertheless,

Rosalbino [13] observed higher values for Ti-16Nb-5Sn and Ti-18Nb-4Sn in Ringer’s

solution: +0.130 and +0.250 V/SCE, respectively. No clear correlation between stabilized

corrosion potential and Sn content of the alloys can be established (Table 2). The

corrosion potentials obtained in Ringer’s solution at pH 7.5 are observed in the stability

regions of TiO2, Nb2O5 and SnO2 oxides of the Ti-H2O, Nb-H2O and Sn-H2O Pourbaix

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diagrams [20], respectively, which confirm the passive behavior of the alloys. Zheng [12]

showed through XPS analysis that the oxide passivating films formed on Ti-16Nb-xSn

alloys are mainly comprised of TiO2, Nb2O5 and SnO2.

The Nyquist and Bode diagrams obtained for the Ti-Nb-Sn alloys at OCP in Ringer’s

solution are presented in Figures 9 and 10, respectively. The Nyquist diagrams (Figure 9)

exhibit depressed and incomplete semi-circles for all alloys, and the Bode diagrams

present a linear relation (slope near -1) between log Impedance and log (Frequency) and

a phase angle near -90o from the low to the intermediate frequencies (Figure 10). These

results show a predominantly capacitive behavior of the metal/solution interface, which is

characteristic of passive materials. Rosalbino [13] observed the same behavior for Ti-

16Nb-5Sn and Ti-18Nb-4Sn alloys in Ringer’s solution, but the values of impedance for

both the alloys were higher than the values obtained for the Ti-30Nb-xSn alloys of the

present work.

Different equivalent circuit models were used to represent the Ti alloys/aqueous

solution interface. The existence of a compact oxide film was represented by an equivalent

circuit with one time constant [21-22], whereas for a duplex oxide film (inner compact and

outer porous layers), a circuit with two time constants was employed [23-25]. This latter

circuit was sometimes modified to take into account the sealing of pores by corrosion

products [22] or the contribution of the space charge layer [26].

The duplex oxide film model fit the experimental data better. The corresponding

equivalent circuit is shown in Figure 11 and consists of the following parameters: R

resistance of the electrolyte, C1 and R1 capacitance and resistance of the outer oxide film

(porous layer), C2 and R2 capacitance and resistance of the inner oxide film (compact

layer). Instead of capacitances, constant phase elements (CPE) were employed in the

fitting routine. The impedance of a CPE is given by ZCPE = [(i)n.C]-1, where C is the

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capacitance, is the angular frequency and n is related to non-uniform current distribution

due to surface roughness or inhomogeneity.

Table 3 summarizes the circuit element values for Ti-Nb-Sn alloys obtained by fitting of

impedance experimental data. The good agreement between the experimental and the

simulated data is confirmed by the 2 values on the order of 10-4. From the impedance

results (Table 3), the following observations can be made:

- The resistance of the inner layer (421 to 776 k cm2) is high and is 4 to 5 orders of

magnitude higher than the resistance of the outer layer (few cm2). This indicates

that the corrosion resistance of the materials is provided by the inner compact oxide

layer.

- Considering the variation of the resistance of the barrier oxide layer as a function of

Sn content, the corrosion resistance of the Ti-Nb-Sn alloys increases from 0 wt %

Sn to 6 wt % Sn and then decreases as the Sn content increases.

- The order of increasing corrosion resistance is as follows:

Ti-30Nb-10Sn < Ti-30Nb < Ti-30Nb-8Sn < Ti-30Nb-2Sn < Ti-30Nb-4Sn < Ti-30Nb-

6Sn.

- The values of capacitance (on the order of 10-5 F cm-2 for the inner and outer layers)

show that the passivating film is thin. The slightly higher values of C2 could indicate that

the inner barrier layer is thinner than the outer porous layer.

The polarization curves of the Ti-Nb-Sn alloys in Ringer’s solution are shown in

Figure 12. The polarization curves have the same shape, which indicates that the cathodic

and anodic reactions occurring on all surfaces are the same. Polarization curves with the

same feature were obtained for Ti-16Nb-xSn alloys in Hank’s solution of pH 7.4 [12].

On the cathodic branch, the main reaction is H2O or H+ reduction. On the anodic

branches, the current density increases first with increasing potential from null-current

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potential to 0.1 V/SCE and then remains constant up to 1.3 V/SCE with values between 5

and 9 A cm-2, which is characteristic of passive behavior. According to Assis [27], the

initial increase of current density could be related to the replacement of the spontaneous

oxide film by a less protective oxide layer, which only becomes stable above 1.3 V/SCE, or

it could be related to the oxidation of TiO or Ti2O3 to TiO2. Between 1.3 and 2.1 V/SCE, the

polarization curves exhibit a peak in anodic current density, and above 2.1 V/SCE, a

secondary passive region with current density stabilizing between 10 and 26 A cm-2 is

observed. The primary and secondary passive current densities are close to the values

obtained by Wang [28] for Ti-16Nb, by Assis [29] for Ti-13Nb-13Zr and by Zheng [12] for

Ti-16Nb-xSn alloys in Hank’s solution. The higher current density in the secondary

passivation region could indicate that the oxide film formed on the surface is more

defective than that formed in the primary passivation region. The anodic peak has been

attributable to O2 evolution [25] but also to possible phase transformation in the passive

oxide film [30].

The corrosion current density (icorr) was determined by the Tafel extrapolation method

of the cathodic and anodic linear branches, and the primary and secondary passive

current densities (ipass1 and ipass2) and the anodic peak current density (ipa) were measured

according to the schematic representation shown in Figure 13. The respective values are

reported in Table 4 for the five alloys.

It is noted (Table 4 and Figure 14) that the corrosion current density decreases from

0 wt % Sn to 6 wt % Sn and afterwards increases as the Sn content increases.

Nevertheless, the very low values of corrosion current density (on the order of 10-7 A cm-2)

depict the high corrosion resistance of all materials in Ringer’s solution. Values of

corrosion current density of the same order of magnitude were obtained for Ti-16Nb-xSn

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alloys in Hank’s and NaCl 0.9% solutions [12]. The increasing order of corrosion

resistance is as follows:

Ti-30Nb-10Sn < Ti-30Nb < Ti-30Nb-8Sn < Ti-30Nb-2Sn < Ti-30Nb-4Sn < Ti-30Nb-6Sn,

which is in agreement with the EIS results (R2 values of the resistance of the barrier film)

(Figure 14).

The improvement in corrosion resistance with increasing Sn content from 0 to 6 wt

% could be related to the decrease of and ” phase content in the alloys. If the

distribution of alloying elements is not homogeneous in all phases, this can lead to the

formation of a less stable oxide film. The improvement of corrosion resistance with the

suppression of ” phase and the increase of alloying element content was also observed

for Ti-22Nb-XHf (X = 2, 4 and 6 wt %) alloys [31]

The subsequent increase in corrosion rate for the alloys with higher Sn

concentrations (> 6 wt %), for which the microstructure consists of phase, can be

attributed to the higher amount of SnO2 oxide in the passive film, which is less stable than

TiO2 and Nb2O5. Zheng [12] also showed that increasing Sn content from 4 to 5 wt % in Ti-

16Nb-xSn alloys has a detrimental effect on the passive film stability in 0.9 wt % NaCl

solution, but no explanation was presented.

From Table 4, the values of the primary, secondary passive current density and

anodic peak current density are very close for the Ti-30Nb, Ti-30Nb-2Sn, Ti-30Nb-4Sn and

Ti-30Nb-6Sn alloys, but are lower than the values measured for Ti-30Nb-8Sn and Ti-30Nb-

10Sn. This depicts a more protective character of the oxide film formed on the former

alloys.

4. Conclusions

Samples of Ti-30Nb-XSn (X=0, 2, 4, 6, 8 and 10) wt % alloys were centrifugally cast

in copper mold, and the resulting microstructures indicated that dendritic solidification

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occurred. The microstructural characterization of samples suggests that the Sn addition

causes reduction of the ω phase precipitation and an increase in the β phase lattice

parameter. For high Sn contents, the β phase lattice value appears to stabilize. It was

observed that the addition of Sn significantly affects the mechanical behavior of the Ti-Nb-

Sn alloys. The minimum values of Vickers hardness and elastic modulus were determined

to occur when the Sn content was close to 6 wt % Sn. Compression tests revealed that

addition of 6 wt % Sn resulted in the maximum value of ductility and minimum value of

compression mechanical strength. Further increase in Sn content resulted in loss of

ductility.

All Ti-30Nb-XSn (X=0, 2, 4, 6, 8 and 10) wt % alloys are passive in Ringer’s solution

at room temperature. The passivating films present a duplex character and consist in an

inner compact layer and an outer porous layer. The inner layer which presents high

resistance provides the corrosion resistance of the alloys. The highest corrosion resistance

was observed for 6 wt % Sn content.

Acknowledgments

The authors gratefully acknowledge the Brazilian research funding agencies

FAPESP (State of São Paulo Research Foundation) Grant # 2011/23942-6, CNPq

(National Council for Scientific and Technological Development) and CAPES (Federal

Agency for the Support and Evaluation of Graduate Education) for their financial support of

this work.

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Table captions

Table 1. Nominal and measured compositions of the experimental alloys.

Table 2. Open-circuit potential (OCP) of the Ti-Nb-Sn alloys measured after 3 h immersion

in Ringer’s solution.

Table 3. R, C1, n1, R1, C2,, n2 and R2 values obtained by fitting the equivalent circuit

model of Figure 11 to the experimental impedance data for the Ti-Nb-Sn alloys in Ringer’s

solution at OCP.

Table 4. Values of corrosion current density (icorr), primary and secondary passive current

densities (ipass1 and ipass2) and anodic peak current density (ipa) determined from the

polarization curves of Figure 12.

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Figure captions

Figure 1. Microstructures of the Ti-Nb-Sn alloys arc-melted and copper-cast molds: (a) Ti-

30Nb, (b) Ti-30Nb-2Sn, (c) Ti-30Nb-4Sn, (d) Ti-30Nb-6Sn, (e) Ti-30Nb-8Sn and (f) Ti-

30Nb-10Sn.

Figure 2. Elemental distribution in the Ti-30Nb-6Sn alloy microstructure after arc-melting

and copper mold casting: (a) Backscattered (SEM) image; (b) Sn distribution, (c) Nb

distribution and (d) Ti distribution.

Figure 3. X-ray diffraction patterns of the Ti-30Nb-XSn (X=0, 2, 4, 6, 8 and 10) wt % after

arc-melting and copper mold casting.

Figure 4. Effect of Sn addition on the intensity of X-ray diffraction peaks related to {301}ω

and {112}ω obtained by slow scanning speed.

Figure 5. Effects of Sn addition to the Ti-30Nb alloy on the Vickers hardness.

Figure 6. Effects of Sn addition to the Ti-30Nb alloy on the elastic modulus.

Figure 7. (a) Stress versus Strain curves obtained in compression tests of the Ti-30Nb-

XSn (X=0, 2, 4, 6, 8 and 10) wt% alloys. Fractographic detail after the compression tests

of the (b) Ti-30Nb and the (c) Ti-30Nb-6Sn alloys.

Figure 8. Variation of open-circuit potential (OCP) for the Ti-Nb-Sn alloys as a function of

exposure time in Ringer’s solution.

Figure 9. Nyquist diagram obtained for the Ti-Nb-Sn alloys at OCP in Ringer’s solution.

Figure 10. Bode diagram obtained for the Ti-Nb-Sn alloys at OCP in Ringer’s solution.

Figure 11. Equivalent circuit model used to represent the Ti-Nb-Sn alloy/Ringer solution

interface.

Figure 12. Polarization curves of Ti-Nb-Sn alloys in Ringer’s solution.

Figure 13. Determination of corrosion current density (icorr), primary and secondary passive

current densities (ipass1 and ipass2) and anodic peak current density (ipa).

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Figure 14. Variation of the corrosion current density icorr and barrier film resistance R2 of

the Ti-Nb-Sn alloys in Ringer’s solution as a function of Sn content.

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Figure 1

(a)

(b)

(c)

(d)

(e)

(f)

Figure 1. Microstructures of the Ti-Nb-Sn alloys arc-melted and cast in copper molds: (a)

Ti-30Nb, (b) Ti-30Nb-2Sn, (c) Ti-30Nb-4Sn, (d) Ti-30Nb-6Sn, (e) Ti-30Nb-8Sn and (f) Ti-

30Nb-10Sn.

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Figure 2

(a)

(b)

(c)

(d)

Figure 2. Elemental distribution in the Ti-30Nb-6Sn alloy microstructure after arc-melting

and cast in copper mold: (a) Backscattered (SEM) image; (b) Sn distribution, (c) Nb

distribution and (d) Ti distribution.

100 µm

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Figure 3

30 40 50 60 70 80 90-0.5

0.0

0.5

1.0

1.5

2.0

2.5

3.0

3.5

4.0

4.5

5.0

5.5

6.0

6.5

Ti-30Nb-10Sn

Ti-30Nb-8Sn

Ti-30Nb-6Sn

Ti-30Nb-4Sn

Ti-30Nb-2Sn

"

"

"

"

"

"

""

"

"

N

orm

ali

ze

d I

nte

ns

ity

(a

.u.)

Angle 2

"

Ti-30Nb

Figura 3. X-ray diffraction patterns of the Ti-30Nb-XSn (X=0, 2, 4, 6, 8 and 10) wt% after

arc-melting and cast in copper mold.

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Figure 4

75 80 85 90

Inte

nsity (

a.u

.)

2

Ti-30Nb

75 80 85 90

Inte

nsity (

a.u

.)

2

Ti-30Nb-2Sn

75 80 85 90

Inte

nsity (

a.u

.)

2

Ti-30Nb-4Sn

75 80 85 90

Inte

nsity (

u.a

.)

2

Ti-30Nb-10Sn

Figura 4. Effect of Sn addition on the intensity of X-ray diffraction peaks related to {301}ω

and {112}ω obtained by the use of slow scanning speed.

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Figure 5

0 2 4 6 8 10

200

225

250

275

300

325

350

375

400

Vickers hasdness test

Nanoindentation test

Vic

ke

rs H

ard

ne

ss

(H

V)

Sn content (wt%)

Figure 5. Effects of Sn addition to the Ti-30Nb alloy on the Vickers hardness.

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Figure 6

0 2 4 6 8 10

40

50

60

70

80

90

100

110

120 Ultrasound test

Nanoindentation test

Ela

sti

c m

od

ulu

s (

GP

a)

Sn content (wt%)

Figure 6. Effects of Sn addition to the Ti-30Nb alloy on the elastic modulus.

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Figure 7

0 5 10 15 200

400

800

1200

1600

2000

Ti-30N

b-1

0S

n

Ti-30N

b-8

Sn

Ti-30N

b-6

Sn

Ti-30N

b-4

Sn

Ti-30N

b-2

Sn

Str

es

s (

MP

a)

Strain (%)

Ti-30N

b

Figure 7. (a) Stress versus Strain curves obtained in compression tests of the Ti-30Nb-

XSn (X=0, 2, 4, 6, 8 and 10) wt% alloys. Fractographic detail after the compression tests

of the (b) Ti-30Nb and the (c) Ti-30Nb-6Sn alloys.

(b)

(c)

(a)

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Figure 8

0 3000 6000 9000 12000

-0.75

-0.50

-0.25O

pen-c

ircuit p

ote

ntial, V

/SC

E

Time, s

Ti-30Nb

Ti-30Nb-2Sn

Ti-30Nb-4Sn

Ti-30Nb-6Sn

Ti-30Nb-8Sn

Ti-30Nb-10Sn

Figure 8. Variation of open-circuit potential (OCP) for the Ti-Nb-Sn alloys as a function of

exposure time in Ringer’s solution.

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Figure 9

0 50000 100000 150000 200000 250000 300000

0

-50000

-100000

-150000

-200000

-250000

-300000

Imagin

ary

im

pedance,

cm

2

Real impedance, cm2

Ti-30Nb

Ti-30Nb-2Sn

Ti-30Nb-4Sn

Ti-30Nb-6Sn

Ti-30Nb-8Sn

Ti-30Nb-10Sn

Figure 9. Nyquist diagram obtained for the Ti-Nb-Sn alloys at OCP in Ringer’s solution.

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Figure 10

10-2

10-1

100

101

102

103

104

105

0

-20

-40

-60

-80

-100

Phase a

ngle

, degre

e

10-2

10-1

100

101

102

103

104

105

101

102

103

104

105

106

Impedance m

odulu

s,

cm

2

Frequency, Hz

Ti-30Nb

Ti-30Nb-2Sn

Ti-30Nb-4Sn

Ti-30Nb-6Sn

Ti-30Nb-8Sn

Ti-30Nb-10Sn

Figure 10. Bode diagram obtained for the Ti-Nb-Sn alloys at OCP in Ringer’s solution.

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Figure 11

Figure 11. Equivalent circuit model used to represent the Ti-Nb-Sn alloy/Ringer solution

interface.

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Figure 12

10-9

10-8

10-7

10-6

10-5

10-4

-1.0

-0.5

0.0

0.5

1.0

1.5

2.0

2.5

3.0

Po

ten

tia

l, V

/SC

E

Current density, A cm-2

Ti-30Nb

Ti-30Nb-2Sn

Ti-30Nb-4Sn

Ti-30Nb-6Sn

Ti-30Nb-8Sn

Ti-30Nb-10Sn

Figure 12. Polarization curves of Ti-Nb-Sn alloys in Ringer´s solution.

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Figure 13

10-9

10-8

10-7

10-6

10-5

10-4

-1.0

-0.5

0.0

0.5

1.0

1.5

2.0

2.5

3.0

ipass2

ipass1

ipa

icorrP

ote

ntia

l, V

/SC

E

Current density, A cm-2

Figure 13. Determination of corrosion current density (icorr), primary and secondary passive

current densities (ipass1 and ipass2) and anodic peak current density (ipa).

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Figure 14

0 2 4 6 8 105.0x10

-8

1.0x10-7

1.5x10-7

2.0x10-7

4x105

5x105

6x105

7x105

8x105

i corr,

A c

m-2

R

2,

cm

2

Sn concentration / wt%

Figure 14. Variation of the corrosion current density icorr and barrier film resistance R2 of

the Ti-Nb-Sn alloys in Ringer’s solution as a function of Sn content.

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Table 1

Table 1. Nominal and measured compositions of the experimental alloys.

Nominal

Composition

(wt %)

Measured Composition (wt%)

Ti Nb Sn O N

Ti-30Nb Balance 29.3 ± 1.5 0.0 ± 0.0 0.159 ± 0.05 0.010 ± 0.003

Ti-30Nb-2Sn Balance 28.8 ± 1.4 2.3 ± 0.3 0.162 ± 0.01 0.007 ± 0.002

Ti-30Nb-4Sn Balance 27.9 ± 1.3 4.3 ± 0.6 0.145 ± 0.03 0.007 ± 0.003

Ti-30Nb-6Sn Balance 28.8 ± 1.4 6.3 ± 0.8 0.093 ± 0.03 0.011 ± 0.005

Ti-30Nb-8Sn Balance 28.5 ± 1.4 8.1 ± 1.1 0.166 ± 0.05 0.009 ± 0.004

Ti-30Nb-10Sn Balance 28.9 ± 1.5 10.8 ± 1.5 0.153 ± 0.02 0.010 ± 0.002

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Table 2

Table 2. Open-circuit potential (OCP) of the Ti-Nb-Sn alloys measured after 3h-immersion

in Ringer’s solution.

%Sn (wt%) 0 2 4 6 8 10

OCP / V/SCE -0.356 -0.333 -0.308 -0.331 -0.361 -0.339

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Table 3

Table 3. R , C1, n1, R1, C2,, n2 and R2 values obtained by fitting the equivalent circuit

model of Figure 11 to the experimental impedance data for the Ti-Nb-Sn alloys in Ringer’s

solution at OCP.

Alloy R

( cm2)

C1

(F cm-2)

n1

R1

( cm2)

C2

(F cm-2)

n2

R2

( cm2)

2

Ti-30Nb

Ti-30Nb-2Sn

Ti-30Nb-4Sn

Ti-30Nb-6Sn

Ti-30Nb-8Sn

Ti-30Nb-10Sn

15

18

17

17

30

19

1.28x10-5

1.76x10-5

1.23x10-5

2.05x10-5

1.89x10-5

1.81x10-5

0.99

0.97

1.00

0.96

0.97

0.99

26

19

14

27

10

14

2.62x10-5

1.98x10-5

2.33x10-5

2.12x10-5

4.02x10-5

3.89x10-5

0.86

0.91

0.90

0.91

0.91

0.88

499620

573880

770180

776820

518130

421910

7.5x10-4

1.8x10-4

1.8x10-4

3.7x10-4

1.7x10-4

5.0x10-4

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Table 4

Table 4. Values of corrosion current density (icorr), primary and secondary passive current

densities (ipass1 and ipass2) and anodic peak current density (ipa) determined from the

polarization curves of Figure 12.

Alloy icorr

(A cm-2)

ipass1

(A cm-2)

ipa

(A cm-2)

ipass2

(A cm-2)

Ti-30Nb

Ti-30Nb-2Sn

Ti-30Nb-4Sn

Ti-30Nb-6Sn

Ti-30Nb-8Sn

Ti-30Nb-10Sn

13.9x10-8

11.2x10-8

10.3x10-8

7.5x10-8

12.3x10-8

19.7x10-8

5.8x10-6

6.4x10-6

5.9x10-6

6.8x10-6

9.2x10-6

8.3x10-6

1.0x10-5

2.0x10-5

1.9x10-5

2.6x10-5

5.6x10-5

3.9x10-5

1.0x10-5

1.1x10-5

1.4x10-5

1.6x10-5

2.6x10-5

2.2x10-5

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EFFECTS OF Sn ADDITION TO THE Ti-30Nb ALLOY

ON THE ELASTIC MODULUS

0 2 4 6 8 10

40

50

60

70

80

90

100

110

120

Ultrasound test

Nanoindentation test

E

las

tic

mo

du

lus

(G

Pa

)

Sn content (wt%)

Graphical abstract

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Highlights

Sn addition causes reduction of the ω phase precipitation

Minimum Vickers hardness and elastic modulus occurred for 6 wt % Sn content

Addition of 6 wt% Sn resulted in maximum ductility and minimum compression strength

All Ti-30Nb-XSn (X=0, 2, 4, 6, 8 and 10%) alloys are passive in Ringer’s solution

Highest corrosion resistance was observed for 6 wt % Sn content