Effect of the Current Density in the Electrical Resistance Sintering Technique to Fabricate a β-Ti-Nb-Sn Ternary Alloy for the Biomedical Sector

The electrical resistance sintering is a fast method to fabricate metallic samples in the eld of metallurgy and it was used to obtain the Ti-Nb-Sn alloy to be applied as biomaterial, variyng different electrical current densities (11, 12 and 13 kA). The powders were obtained by mechanical alloying, then they were compacted at pressure of 193 MPa during 700 ms. The structure and microstructure of the powders and the samples was evaluated by x-ray diffraction, by Field Emission Scanning Electron Microscopy and electron backscattered diffraction. The mechanical properties were evaluated by microhardness assay and corrosion resistance was made in Ringer Hartmann’s solution at 37ºC. The samples are formed by α, α” and phase β. The % of phase β in the samples obtained at 11, 12 and 13 kA was 96.56, 98.12 and 98.02 respectively. The peripheral zone present more presence of microporosity than the central zone. The microstructure is also formed by bcc-β grains equiaxial, and the samples obtained at 12 kA present better homogeneity of the microstructure. The grain size increased with the increase of the electrical current density. The microhardness are in the range of 389-418 HV and decreased with the increase of electrical current density. Corrosion tests proved excellent corrosion resistance of the alloys (0.24-0.45 µA/cm 2 ). The standard deviation of kinetic parameters of the samples at 11 and 13 kA were very higher, related to the lack of homogeneity of the microstructure.


Introduction
The studies about metallic alloys based on titanium for application in orthopedics is a eld in a constantly development, due to the inconveniences of commercial metallic prostheses that lead to long-term failure and to rejection by the body. The main characteristics required to the success of these biomaterials are nontoxicity, good resistance to corrosion and elastic modulus close to bone tissue [1][2][3]. It is known that metallic materials based on cobalt, chromium, aluminum, vanadium and nickel promote allergenic and carcinogenic effects, in addition to proven respiratory and neurological disorders [4,5]. Corrosion resistance is closely linked to the chemical composition of these materials, such as the content of elements stabilizing hcp (α phase) and bcc (β phase) crystalline structures and the type of processing that was employed, which will interfere with their microstructure and morphology. Regarding the elastic modulus, most of the employed materials have higher values than that of the bone tissue, such as Co-Cr, Ti-6Al-4V or Ti-CP alloys [6,7]. This difference leads to the adaptive bone resorption, which lasts for 10 to 15 years, when the patient needs surgery for repairing and replacement of the prosthesis [8]. This phenomenon arises due to the cushioning effect whereby the bone absorbs all the charge loaded on the prosthesis as its elastic modulus is larger. The bone is daily discouraged and its fragility intensi es around the device causing the prosthesis to lose stability. In the short term, these effects are not observed. However, a prosthesis that requires bone healing, which will be inserted to remain in the pacient for life, tends to fail in the face of this problem. In order to repair such inconveniences, Ti alloys based on Nb and Sn have been studied as it is known Nb (beta isomorphic type stabilizer) and Sn (neutral element) are elements that do not present toxicity and are good candidates for making Ti alloys with lower elastic modulus [9] and increased mechanical strength [10]. Based on phase diagram of the Ti-Nb binary system [11], Nb exhibits complete solubility in Ti above 882 °C, which allows to study its in uence over the entire range of Nb content in a system and to observe the modi cation of its properties by diffusion process [11,12]. At lower temperature, it is possible to notice two stable phases (α + β), but rapid cooling of the β phase results in a structure composed by metastable phases. In the case of Ti-Nb system, this sudden cooling results in the martensitic transformation of the β phase, which can result in the α' or α" phases depending on Nb content [13]. The effects of Nb content on the mechanical properties of Ti-Nb alloy and its relationship to the β phase content were evaluated by Yahaya et al. preparing these alloys by powder metallurgy (PM) [14]. It was found that up to 35% (wt%) of Nb, the alloy do not cryistallize as α phase.
However, under higher contents of Nb, such as 45% (wt%), resulted in less resistent alloy to the compression, despite of lower elastic modulus (13.46GPa). In general, the the alloy with 35 (wt%) of Nb presented the most adequate characteristics regarding the microstructure and mechanical properties. In most works on Ti-Nb-Sn alloys, they used spark plasma sintering (SPS) and conventional powder metallurgy (PM) for their preparation. There is a lack in the study on the preparation of these materials by electrical resistance based sintering technique (ERS). It is known that the rst patent for electric current sintering was registered in 1906 by Lux [15]. Few years later came Taylor's [16] (in 1933) and Lenel's studies [17] (in the 1950s), who called this technique electrical pressure resistance sintering (ERS). Later, during the 70 and 80 decades, the ERS technique received a new boost mainly by Sovietic and Japanese researchers [18,19]. Nowadays, the ERS technique is studied as a modality of the electrical eld assisted sintering technique (FAST), which is the common name for the technique based on electrical consolidation of powder metallurgy (PM). During this long period, many FAST variants were developed for use on an industrial scale, as overall ultimate goal [20][21][22]. Most popular FAST is the so-called SPS, in which low-wear (electrically conductive) graphite dies and punches are used, to combine the application of alternating current and vacuum or argon for atmosphere controll.
The ERS technique allows the use of durable electrical insulation matrices, and the process can take place in the air, as a typical cycle being completed in seconds. In addition, the required equipment can be easily adapted from the well-known resistance welding technology, a well-tested technology.
Among the advantages of ERS, compared to PM, there are the use of relatively low pressures (about 100 MPa) to reach the nal material with very high density, extraordinarily short processing times (about 1 s or less), and the possibility to operate in the air without the need for a controlled atmosphere [23,24]. The main disadvantages of the ERS technique arise from operational di culties (incomplete knowledge of parameters and how certain parameters in uence the process) and non homogeneous temperature distribution in the powder mass [25]. In order to increase knowledge about the ERS process, especially in the applicability for the consolidation of alloy with a high content of β phase stabilizing element, the in uence of the electric current on the physical, mechanical and electrochemical characteristics was evaluated in the promising Ti-34Nb-6Sn alloy for its use as medical implant.

Powders preparation
Elemental powders of Ti (99.9%), Nb (99.8%) and Sn (99.8%) were purchased from Atlantic Equipment Engineers, which presented particle sizes around 30 µm, 16 µm and 17 µm, respectively. They were weighed in order to obtain alloys with 34 (wt%) Nb and 6 (wt%) Sn in a glove box to minimize the contamination of oxygen and nitrogen.
Then, they were mixed with a process control agent (PCA), NaCl (in powder) was used, in the amount of 1.5 (wt%) in relation to the total mass of the powders. The mixture was made for 30 min in a tumbler mixer (Inversina 2L-Bioengineering AG) at 90 rpm, followed by 72 h milling process to promote mechanical alloying.
The milling process was carried out at room temperature in a high energy mill (Retsch-PM 400/2), for 72 h at 240 rpm, interspersed with 10 seconds of rest every 8 min of grinding interval. Before starting the process, the steel vessel was purged by argon for 4 times. The proportion of steel balls (5 mm of radius) used was 10:1 in relation to the powder mass.

Synthesis via ERS
The consolidation of specimen was carried out using the ERS technique in an equipment developed within the EU funded EFFIPRO project, used in collaboration with the AMES Company (Barcelona-ES). The materials were produced in a cylindrical matrix with a diameter of 2.2 cm and a thickness of 0.55 cm. In order to obtain samples with these dimensions, the density of the alloy of d Ti34Nb6Sn = 5.52 g/cm 3 was taken into account for the calculation of powder mass. The mixture powder was placed in a silicon nitride die between two tungstencopper electrodes [26] and compacted at pressure of 193 MPa. The current intensities tested for consolidation were 11, 12 and 13 kA during 700 ms of sintering time.

Phase/chemical and microstructural characterization
For the characterization and evolution of the current phases as well as the lattice parameters of the powders, was used X-ray diffraction (Bruker/D2Phaser). A Cu Kα radiation (l = 1.541 Å) was use that works at 30 kV and 10 mA. The measurement was made in the range of angles between 30 and 90 degrees, with a step of 0.02º every 10 s. The re nement of the structure and the quantitative phase analysis were carried out using the free MAUD software (version 2.94). The phases and diffraction planes were analyzed by comparing the d value of each peak of the diffraction pattern from those of the Inorganic Crystal Structure Database (ICSD). In addition, the crystallites size and induced micro strain in the titanium lattice were calculated from the peak broadening and peak positioning of the X-ray diffraction pattern of the powder milled and the sintered parts according to Williamson-Hall equation [27].
Where β = Full width at half maxima, k = shape factor generally taken as 0.9 [28], λ = wave length of used radiation in XRD, ϴ=peak diffraction angle, D=crystalline size and ε=induced strain in the lattice. A graph was plotted between 4 sin ϴ and β*cosϴ for selected peaks of powder sample, along the x and y axis respectively and a best tted linear curve is drawn for different Ti peaks.
The particle size distribution was measured using particle size analyzer by laser scattering (Mastersizer 2000-Malvern Instruments). Distilled water was used to disperse the powders. To evaluate the contamination by O and N, it was used the LECO TCH600 series equipment with an inert gas fusion analyzers.
The microstructure of the powders and the sintered was characterized by Field Emission Scanning Electron Microscopy (FE-SEM) (ZEISS-ULTRA 55) with backscattered electrons (BSE), secondary electrons (SE) and X-ray energy dispersive detectors (EDS) (Oxford Instruments Ltda). To characterize more precisely the microstructure, like the grains and the phase quanti cations was used electron back-scattered diffraction (EBSD) with a scanning electron microscope (Zeiss-ULTRA 55 operating at 20 kV) equipped with an Aztec HKL Max System (Oxford Instruments Ltda) under an acceleration voltage of 20 kV with a step size of 0.1 µm, selecting the three possible phases to be analyzed, β-Ti, α-Ti and α"-martensite. For the powders, they were embedded in resin and prepared metallographically from cross-sections in order to analyze particle size, homogeneity of the Ti, Nb and Sn elements and the evolution of the alloyed process.

Microhardness test
Vickers microhardness tester (HMV-SHIMADZU) was used to evaluate the samples consolidated by ERS. To measure the microhardness was made 12 indents on the surface of each sample applying 490.3 mN (HV 0.05) during 12 s. For each indent a distance of 50 µm was obeyed.

Corrosion resistance in Ringer Hartmann's solution
The corrosion behavior was studied by measuring the open circuit potential (OCP) for 35 min and potentiodynamic polarization assay after that time in a potentiostat/galvanostat 144 (Metrohm potentiostat-PGSTAT204) on a surface of 0.785 cm 2 , using a conventional three-electrode cell in a Ringer-Hartmann's solution (6<pH<7), (NaCl 5.7 g/L, KCl 0.358 g/L, CaCl 2 0.221 g/L, Lactate 4.956 g/L) at 37°C. The Ag/AgCl electrode was used as a reference electrode and Pt thread as auxiliary one. All potential values were expressed in relation to this reference electrode hereinafter as well as the current density was normalized by geometric area. The electrochemical tests consisted of OCP measurement (for 35 min) and polarization from OCP to 2 mV/s scan rate. The corrosion parameters as the Tafel slopes (b a and b c ) in the region of cathodic and anodic process domains, corrosion potential (E corr ) and corrosion current (I corr ) were determined using Wolfram Mathematica 12.1 software. The Electrochemical Impedance Spectroscopy (EIS) technique was followed within the 100 kHz-5 mHz frequency range with a sinu-soidal amplitude wave of 0.01 V on Eocp.
Potentiodynamic polarization curves were carried out by swiping the potential from 1 V to 3 V at the 2 mV/s scan rate. The parameters were acquired by the Zview 3.5 software by tting equivalent circuits.
Chronoamperometric tests were carried out at 1 V (passivation potential) for 5 min with a 0.01 s time interval for data acquisition in the previously prepared sample.

Results And Discussion
For the analysis of particle size (seen in table 1), it is noted that after the milling process by mechanical alloying (MA), the mixture showed a reasonable decrease in particle size, with the average obtained being similar than the initial size of the Nb and Sn, but smaller than the initial titanium particles. The maximum particle size found was higher than the initial size of the 3 elements Ti, Nb and Sn, approximately 171µm. It is certain that the MA process promotes a signi cant reduction in the size of the particles over time, and also promotes their agglomeration due to the more interaction of the surface promoted by to the reduction of sizes. Another factor that contributes to the agglomeration process is the cold welding effect and the increase in temperature during the constant impacts of the balls and the friction caused between the balls and the powders. The use of PCA tends to diminish these effects. However, even with the use of this agent, due to the high milling time, the cold welding process is di cult to be controlled. In part, the agglomerated powders, due to the high milling time, and the greater impact of the powders due to shocks with the vessel balls, can contribute to decrease the porosity, towards a higher density. In addition, longer period of milling can promote morphological changes in the particles, resulting in more spherical particles creation, which are more suitable for the compaction and sintering process.
By the XRD pattern shown in gure 1, it can be seen the 72 h of milling resulted a material structured under two phases. Being them: phase α, represented by the compact hexagonal (hcp) structure, formed by a small peak that appears around 35°, characterized by the plane {100} α and phase β, represented by the body centered cubic (bcc) structure. The milling process promoted almost 80 % of β phase formation (Table 2) Table 3. O and N content in the powders after MA process.

Milled Powder O (wt%) N (wt%)
Ti-34Nb-6Sn 1.09±0.04 0.216±0.08 It is more clear to observe two distinct regions present in the microstructure of the material after milling process in gure 3a. The homogeneity of the particles (Figure 3b) is indicated by the map, with good distribution of the three elements (Ti, Nb and Sn). These structures of suitable uniformity, resemble a plate, which consists of a ne and homogeneous distribution of the components of the solute in the titanium matrix. The other region, on the other hand, is visible the presence of zones enriched by niobium and titanium that did not react within the milling time, remaining trapped in the microstructure of the material, but are less amount present in the microstrure. From the analysis by EDS (Figure 3c), the components present the percentage reasonable close to the compositional. The second peak present in Figure 3c is represented by the carbon detected due to the resine used to embbed the samples.
In this case, tin presents 4% more content than added. This characteristic may be due to the low homogeneity of this element in the regions presented by the map (in green color). However, in general, the MA process promoted good homogeneity of the structure by the milling time used.
The XRD pro le of materials sintered via ERS with electrical current density of 11, 12 and 13 kA, are shown in gure 4 and were compared wih the XRD pro le powders after the MA at the same scale. As can be seen, the XRD patterns are formed by peaks related to the α"-martensite and β phases in all samples. The orthorhombic phase α", presents an elastic modulus smaller than the hexagonal α phase [31], being interesting for orthopedic application. In also gure 4a the typical microstructure obtained by EBSD evaluation show the slight presence of Ti-α (represented in blue) that could not be posible identi ed by the XRD.
Among the component elements of the alloy, niobium is present in greater proportion, compared to the tin.
Besides, niobium has a melting point (2468°C) higher than tin (231,97°C). Thus, the diffusion rate is lower compared to tin. Due to the low diffusion rate, homogeneity it is also smaller, creating areas rich in titanium which in turn become orthorhombic phase during the sintering process via ERS. It is noted that obtaining samples with electrical current density at 11 kA promoted more formation of phase α"-martensite, where the peaks are more intense (Figure 4a). This fact proves the di culty in diffusing niobium, due to the low electrical current used. The energy of the system was lower than 12 and 13 kA to ensure su cient diffusion and thus promoted the formation of the α"-martensite phase, as evidenced by the higher peak intensity. The increase of the electrical current intensity, the peaks referring to the α"-martensite phase decreased. The greater stabilization of phase β can cause the transformation to β' instead of α"-martensite. However, the presence of phase by its nanometric size that would require its observation in TEM because it is not posible observed it in XRD. When comparing the XRD pro le of the samples at 12kA and 13kA, no signi cant differences are observed. Visibly, a difference is observed between the patterns of samples obtained by ERS and powders after MA process. The main differences from the XRD pattern of the powders is the unique presence of phase β, being mostly Ti-β and also niobium particles that present the lattice parameters close to Ti-β, which makes di cult to distinguish them. Another signi cant difference is the widening of the Tiβ phase peaks, related to the {110}, {200} and {211} plans present in the powders' XRD pro le. The larger peaks and their electrical lower intensities are due to the milling time used, the high re nement of the niobium grains and the micro-deformation inside these grains [32]. According to Fecht, it was found that grain re nement and micro-deformation increase during the milling [33]. Besides, increasing the milling time, the particles are increasingly re ned. In the work of Zhang et al., the effect of milling nioium, titanium and silicon powders was evaluated at 2, 5, 10, 20 and 40 h. After 20 h the peaks became wider and less intense [34].
The mechanical milling in signi cantly long times increases beyond the local deformations in shear strips, with high displacement densities at the beginning of the re nement process. These displacements combine among themselves creating grain limits formation of low angle [35]. These small limits of formed grains are transformed into grain limits at a greater angle, producing more re nement particles. In addition, the microdeformation increases, during the press balls in MA, along with the particles, collide continuously and irreversible deformations increase leading to micro-deformation. In table 4 shows the parameters obtained of the linear t of powders after 72h of milling time and the sintered samples by ERS. Figure 5a-d shows the best tted linear plots between the powders milled and samples sintered with 11, 12 and 13 kA respectively. Crystallites size (D) were calculated from the intercept cut on y axis (represented by βcosθ) while lattice strain (ε) was determined by the slope of the tted straight line using Wlliamson Hall equation [27]. In order to compare the evolution of crystallite size and residual stresses after material consolidation via ERS.
The level of deformation experienced by the lattice was evaluated analyzing their corresponding crystallite size and residual micro-strains indicated in table 5. In the present study, the crystallite size increased signi cantly after consolidation process of the Ti-34Nb-6Sn alloy compared to the initial value obtained during the MA. The deformation mechanism that dominant in the ball milling process is the formation of shear bands which have high dislocations density due to the constant impact associated with the powder particles on the balls as discussed. The high milling time used, fragmentation of the sub-grains occurs from the region where unstrained shear band present in the previous material. Due to grain cracking, the degree of randomness of the sub-grains orientation increases. The grain size reduction occurs till the complete random orientation of the sub-grains obtained [37]. It is evident that the sintering temperature contributed to the signi cant increase the crystallite size in all conditions. However, in the condition of lower electrical current density (11 kA), there is evidence it promoted a greater number of defects introduced in the microstructure [38], since the crystallite size is signi cantly higher in this condition. While decreasing with increasing electrical current density. Likewise, the sintered samples at 11 and 12 kA showed a decreasing trend of the lattice strain, compared to the sample of milled powders. However, at an electrical current density of 13 kA, the value increased again, with a value close to that obtained after milling process. In this condition, there was a high level of material transfer and the diffusion process was more accelerated induced by the high electrical current density used, at the same time of consolidation than the others. In this way, the plasma creation between the particles contributed to the fast consolidation of the material along with the Joule effect and plastic deformation as a result of compaction in the ERS technique. According to Singh et al., the high concentration of defects is produced during the densi cation of the samples by the rapid heating and associated with the rapid deformation of the particles [39]. This defect formation is aided by the high electrical current used essential to generate the heating. Kim et al., and Besson & Abouaf, found that this defect induced by the deformation is responsible for the dynamic growth of the grains [40,41].
Some researchers have reported that the higher heating rate, the diffusion process of grain contour is driven by very substantial stresses (inversely proportional to the radius of curvature of the pore), since diffusion on the surface does not have enough time to "smooth" pore surfaces [42]. At the same time, due to densi cation, the pore size decreases signi cantly, which prevents them from exerting the xation effect. For grain growth to occur, the sintering time must be long enough than this critical point. But, for high heating rates, grain growth is slowed due to shorter processing time. In this present work, the heating time was constant, only the current density being varied. However, the mechanism of densi cation and diffusion are the same, since with the increase of the electrical current density, the system energy increases, consequently increasing the heating temperature. Thus, even with the high temperature provided by the increase in electrical current density, the time was not enough to promote an increase in the crystal size, but promoted a decreasing of 17%. The powders prepared by MA after 72 h, was consolidated via ERS. The rst sintering was using an electric current density of 11 kA, represented by gure 6a-c. The current used, promoted an advance in the alloy consolidation. It is noted that in the peripheral region, the material was poorly sintered, con rmed by the presence of microporosities in the microstructure (Figure 6a). In these regions there was less heat distribution, as the presence of solute (niobium particles) that has not just diffused into the titanium matrix is still noticeable (Figure 6b). By line analysis (Figure 6c), the lack of uniformity in the peripheral region of the material is con rmed. The tin element seems to be well dissolved in the titanium matrix, due to the linearity presented by its curve. The niobium, on the other hand, when entering the brightest contrast regions, increases its intensity abruptly. This abrupt transition between the elements indicates a transient diffusion of them, which some reach approximately 2.5 µm, resulting from the lack of heat obtained by the system. In table 6, the percentages of phase obtained by the EBSD analysis indicate a predominant phase β formation, being higher than 95% in all conditions. In the sintering at 11 kA, the percentage of β phase is smaller than at 12 kA and 13 kA. In this condition, has more pixels without indexing, which is evident is the decrease in the small amount of α phase determined, due to the lower amount of residual titanium non-diffusing. The β phase at 12 kA has an increase in relation to the lower electrical current intensity used of 1.6%. The β phase at 13 kA, it was higher at around 5% compared to samples obtained at 11 kA. The percentage of β phase also con rms the microstructure shown in gure 9, predominantly formed by bcc-β equiaxis grains. It is also noted that with the increase of the electrical current intensity, the α''-Ti phase decreased as well as the α-Ti phase. As studied by Gouvea et al., the same growth trend of β grains was observed, by increasing the electric current from 14 kA to 16 kA, in addition to the bimodal distribution of grains in a higher current intensity (16 kA) [43]. In our case, the process also occurred at a higher electrical current (13 kA), however, less than they used at work, to pruduce samples with 16 mm diameter compressed at 50 MPa. The possibility to increase the grain size early, may be due to the lower content of β phase stabilizer element used, and the longer milling time. In their work, were used niobium and molybdenum as β phase stabilizers. The microhardness was measured in the cross sections of the samples, in order to obtain more reliable results of the different conditions. In addition, due to the greater homogeneity of the microstructure of the . In this case, the hardness was greater in all cases compared to those found in the present study. The hardness value depends on several factors, such as the composition of the alloy, the microstructure and the surface conditions and also the type of processing used to obtain the samples. Comparing with Ti-CP and Ti-6Al-4V alloys that present hardness values around 200 HV and 340 HV [47], the values found are much higher.  This process occurs in different potential ranges due to the heterogeneity.
However, the potentiodynamic polarization curves (PPCs) (see Figure 11) show the samples obtained at 11 and 12 kA exhibit lower corrosion current density (i corr ) and higher corrosion potential (E corr ). Figure 11  . It was evident that hardness and E corr are inversely proportional parameters. This fact can be observed in the present work, since the hardness decreased when the current intensity increased from 11 kA to 12 kA, the E corr value increased. However, for sample obtained at 13 kA, the hardness value was smaller than obtained at 11 and 12 kA, but E corr decreased signi cantly. Possibly this fact can be attributed to a non-uniformity of the microstructure, con rmed by the heterogeneity of the grain sizes.
These potentials obtained (Table 8) by the polarization curves were signi cantly lower than those obtained from the OCP measurements, because the polarization test was started at a cathodic potential and in this way the passive oxide lm on the surface was partially removed due to the cathodic polarization.
The polarization resistance followed the same trend as the E corr values to inherit at 11 and 12 kA, altered a decrease in R p . In the condition at 13 kA, the R p increases signi cantly. However, this increase can be explained by the lack of microstructural homogeneity. Figure 11 shows regions characterized by an almost constant current density, starting where indicated by the dashed lines. This region indicates the formation and growth of a passive lm on the surface of the samples. In the conditions at 11 and 13 kA it is noted that the passivation process starts at a potential close to 1.7 V and for the material obtained at 12 kA, start of passivation was close to 1.8 V, demonstrating a later passivation. This can be con rmed by the corrosion resistance paremeter, Cr, obtained by the use of i corr of each sample (see Table 8).
In the work of Mavros et al., the Ti-Nb-Zr-Ta alloy ssystem was obtained for biomedical application by the SPS technique, with good resistance to corrosion. The excellent corrosion resistance of β-type alloys formed by refractory elements such as niobium, can contribute to the formation of a passive lm layer that is not released into the environment [49]. In addition, this oxide layer formed on the surface of titanium alloys, as well as its composition, affects the corrosion response of these alloys when used to manufacture orthopedic prostheses [6].
The corrosion current density found for this type of alloy is in the range of 0.4 to 0.7 μA / cm 2 . It is worth mentioning that the solution used was NaCl, which has an ionic concentration and pH different from those present in the body uid. Alloys of the Ti-Nb-Zr system obtained also by SPS showed a much higher corrosion current density value compared to the present work, being 2.42 μA / cm 2 for the Ti-13Nb-13Zr alloy. Regarding commercially pure Ti, one of the most used materials in the biomedical sector, I corr is approximately 3 μA cm 2 [50], signi cantly higher than the values found for the Ti-34Nb-6Sn alloy in all conditions in the present work.

Conclusions
Ti-Nb-Sn system with 34 wt% (Nb) and 6 wt% (Sn) was obtained by electrical resistance sintering (ERS) process at 11, 12 and 13kA of electrical current density. The effect of current in the microstructure was evaluated. The main conclusions are highlighted above: The samples are structured under α, α" and phase β. Increasing the electrical current density, the phase β increased to values around to 98%; The samples obtained at 12kA presented good homogeneity of the microstructure (less micro and nanoporosity and also less presence of particles unsoluble); The microstrucutre are formed by bcc-β grains equiaxial, with the grain size more uniforme in the samples obtained at 11 and 12 kA; The microhardness decreased with the increase in electrical current density, and the values are in the range of 389-418 HV; Corrosion tests proved excellent corrosion resistance of the alloys with low current densities, despite presence of micro and nanoporosity.