2015-AB-Investigations into Ti–(Nb,Ta)–Fe alloys for biomedical applications

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Investigations into Ti–(Nb,Ta)–Fe alloys for biomedical applications

Arne Biesiekierski a ,Jixing Lin b ,c ,Yuncang Li d ,Dehai Ping e ,Yoko Yamabe-Mitarai e ,Cuie Wen d ,a ,?

a

Faculty of Science,Engineering and Technology,Swinburne University of Technology,Hawthorn,Victoria 3122,Australia b

Department of Material Engineering,Zhejiang Industry &Trade Vocational College,Wenzhou,Zhejiang 325003,China c

College of Materials Science and Engineering,Jilin University,Changchun,Jilin 130025,China d

Department of Aerospace,Mechanical and Manufacturing Engineering,RMIT University,Melbourne,Victoria 3000,Australia e

National Institute for Materials Science,Sengen 1-2-1,Tsukuba 305-0047,Japan

a r t i c l e i n f o Article history:

Received 9September 2015

Received in revised form 7November 2015Accepted 8December 2015Available online xxxx Keywords:Titanium

Orthopedic biomaterials Admissible strain Microstructure High strength

a b s t r a c t

In this study,a Ti–(Ta,Nb)–Fe system was investigated with aims toward the development of high strength,biocompatible titanium alloy suitable for the development of porous orthopedic biomaterials with minimal processing.Notable ?ndings include yield strengths of 740,1250and 1360MPa for the Ti–12Nb–5Fe,Ti–7Ta–5Fe and Ti–10Ta–4Fe alloys,respectively,with elastic moduli comparable to exist-ing Ti-alloys,yielding admissible strains of 0.9?0.3,1.2?0.2and 1.13?0.02%for the Ti–12Nb–5Fe,Ti–7Ta–5Fe and Ti–10Ta–4Fe alloys,respectively;more than twice that of human bone.Observed microstructure varied signi?cantly depending on alloy;near pure b -phase was seen in Ti–12Nb–5Fe,b with some x precipitation in Ti–10Ta–4Fe,and a duplex a +b structure was observed throughout the Ti–7Ta–5Fe.In addition to suitable mechanical parameters,all investigated alloys exhibited promising corrosion potentials on the order of à0.24V SCE,equalling that seen for a C.P.-Ti control at à0.25V SCE,and substantially more noble than that seen for Ti–6Al–4V.Electrochemical corrosion rates of 0.5–3l m/year were likewise seen to agree well with that measured for C.P.-Ti.Further,no statistically signi?cant difference could be seen between any of the alloys relative to a C.P.-Ti control regards to cell proliferation,as investigated via MTS assay and confocal microscopy.As such,the combination of high admissible strain and low corrosion indicate all investigated alloys show signi?cant promise as potential porous biomaterials while in the as-cast state,with the Ti–10Ta–4Fe alloy identi?ed as the most promis-ing composition investigated.Statement of Signi?cance

The ?ndings of this paper are of signi?cance to the ?eld of metallic biomaterials as they detail the devel-opment of alloys of satisfactory biocompatibility and electrochemical behaviour,that furthermore dis-play exceptional mechanical properties.Notably,both extremely high compressive yield strengths and admissible strains,up to 1.36GPa and 1.2%respectively,are reported,exceeding or rivalling that seen in traditional alloys such as Ti–6Al–4V,which typically displays compressive yield strengths and admis-sible strains on the order of 895MPa and 0.81%respectively,as well as modern alloys such as Gum Metal or TNZT.That this is achieved in the absence of thermomechanical processing represents a signi?cant and novel outcome of substantial bene?t for application as a porous biomaterial.

ó2015Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved.

1.Introduction

Metallic biomaterials play major roles in many branches of modern medicine,and nowhere are these more apparent than orthopedics.Due to their excellent biocompatibility and mechani-cal properties,Ti-based alloys are the material of choice for both skeletal repair and joint arthroplasties [1–4].Unfortunately,the use of conventional,high Young’s modulus (E )alloys such as Ti–6Al–4V in these load-bearing roles may result in a phenomenon known as stress shielding,caused by a mismatch in the stiffness of the implant with the surrounding bone [1].While the overall stiffness of the implant may be reduced via increasing porosity,this must simultaneously result in a drastic reduction of mechani-cal strength [5–7].

https://www.360docs.net/doc/697557872.html,/10.1016/j.actbio.2015.12.010

1742-7061/ó2015Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved.

?Corresponding author at:Department of Aerospace,Mechanical and Manufacturing Engineering,RMIT University,Melbourne,Victoria 3000,Australia.

E-mail address:cuie.wen@https://www.360docs.net/doc/697557872.html,.au (C.Wen).

As such,a value of great importance in designing new orthope-dic biomaterials is the admissible strain (d ),the ratio of yield strength (r YS )to E [8,9].A material with a high d is thus best able to maintain the required mechanical strength even as the stiffness of the implant is decreased to closer match that of the osseous tis-sue.A reasonable target for a minimum d is that of bone itself,which falls between 0.43%and 0.55%[10].However,the formation of a porous metallic structure is closely tied with the initial manu-facturing step,preventing traditional thermomechanical working (e.g.cold-rolling,equal channel angular pressing and hot forging)to improve strength.Further,given the larger surface area and thus attendant risk of oxidation of porous materials,intensive thermal treatments have their own complexities.

Hence,the issue at hand;the absence of high-d materials avail-able in an as-cast or minimally processed state.Of the conventional Ti alloys in clinical use,only Ti–6Al–4V,Ti–6Al–7Nb and austenitic Nitinol displayed d values exceeding that of bone in the absence of mechanical working,at 0.78%,055%,and 0.83%,respectively [11–13].However,as these three alloys all contain elements known to have pronounced deleterious effects on the human body (namely Al,V and Ni)[14,15],there has been a signi?cant push to develop new biomedical titanium alloys with a high d ,using highly biocom-patible alloying additions such as Ta,Nb,Zr and Sn [2,16].

One additive that has seen relatively little attention,however,is iron (Fe).Both cheap,plentiful,and a strong b -stabilizer,Fe is noted for improving the mechanical attributes of binary and ternary Ti-base alloys;namely increasing the mechanical strength and reduc-ing the E with only small additions.%[8,17].While it has found some use in older Ti-based biomedical alloys such as Ti–12Mo–6Zr–2Fe [18],its relative scarcity in modern literature is likely due to concerns over the biocompatibility of Fe;although Ti,Nb,and Ta are all considered to display excellent biological responses,Fe is genotoxic and moderately cytotoxic,while also being prone to corrosion [19,20].While including this element may thus at ?rst seem counter to the stated goal of yielding a highly biocompatible alloy,these issues are distinct from more concerning attributes such as allergenicity or carcinogenicity in that minimum safe con-centrations exist.As such,they may be mitigated via a characteris-tics of the other alloying elements;the passive oxide ?lm formed by Ti on contact with oxygen (O 2),which should reduce corrosion signi?cantly compared to pure Fe [21].Evidence for this can be seen in studies on similar alloys in the Ti–Ta–Fe and Ti–Ta–Fe–Zr systems investigated by Kuroda et al.[21,22]previously.These alloys have shown both good mechanical properties,and corrosion resistance comparable to or exceeding that of the common biomedical materials of commercially pure Ti and Ti–6Al–4V.2.Alloy design and experimental methods 2.1.Design rationale for the novel titanium alloys

The design of novel Ti alloys is greatly simpli?ed by considera-tion of the variation of alloy phase as a function of certain elec-tronic parameters,namely mean bond order (Bo ),mean molecular d -orbital energy level (and the mean electron to atom ratio (e /a )of the alloy [23–25].As these values are both

closely related to observed alloy microstructure,and may be calcu-lated from fundamental atomic data,it is possible to use these parameters along with known phase maps to calculate the desired composition for a given microstructure [22–24,26].As such,com-positions of Ti–12Nb–5Fe,Ti–10Ta–4Fe and Ti–7Ta–5Fe were determined as values which are expected to fall solidly within the b -stabilized region.While this reduces the chance of obtaining the absolute Young’s moduli (E )for the b -phase,with e /a ratios falling above the limit of b -phase stability at 4.15[24],these com-positions were expected to be suf?ciently stabilized to avoid the risk of the large scale precipitation of coarse particles of undesir-able a or x -phases.Ikeda et al.[25]found that e /a ratios in the range of 4.26–4.35yielded minimal hardness attributable to these additional phases in Ti-alloys,suggesting that these ranges would also yield reduced moduli.

Finally,it should be noted that modeling by Song et al.[8]indi-cates that a large increase in d may be achieved in a binary Ti–XFe alloy where X =7at.%,but will not be achieved with a smaller Fe content of X =4at.%.This agrees well with known experimental behavior;from application of –Md design theory,7at.%Fe repre-sents an alloy falling on the b /b +x transition,and thus a mini-mally stabilized b -phase,which is known to yield the lowest E [23,26].Knowing this,it becomes possible to target equivalent regions of the Ti-phase diagram with reduced Fe contents through the addition of other b -stabilizing elements such as Ta or Nb.

2.2.Material preparation and characterization

Samples were produced via cold crucible levitation melting (CCLM)from pure metals,and were analyzed in the as-cast state.The nominal compositions of the alloys are presented in Table 1,along with predicted electronic parameters;alloys will be referred to by these compositions hereafter.Samples were cut via electrical discharge machining into small discs with a thickness of 2.00mm and 10.00mm diameter,for assessment of microstructure and bio-compatibility,and small rods with a height of 10.00mm and 5.00mm diameter,for mechanical testing.Where necessary,sam-ples were further sectioned via diamond and water-cooled SiC saws for characterization.

X-ray Diffraction (XRD)was performed on the small disc spec-imens,ground via SiC paper and polished with colloidal SiO 2to a mirror ?nish,then run on a Bruker D8diffractometer with a copper K a source and a 2h scan range of 20–85°.To ensure consistency,diffraction was performed on at least 3separate disc specimens for each alloy.Optical Microscopy (OM)and Scanning Electron Microscopy (SEM)were then performed on the same specimens as the XRD samples,after etching with Kroll’s reagent until fea-tures could be resolved.OM was performed using a Leica MEF4M optical microscope,SEM and elemental characterization by means of Energy Dispersive Spectroscopy (EDS)were performed with a Zeiss Supra 40VP Field Emission microscope.Transmission Elec-tron Microscopy (TEM)was performed on specimens cut from small disc specimens using a diamond saw,then mechanically ground with SiC paper to a thickness of $50–90l m.Next,samples of suf?cient size for analysis were punched from the foils,and elec-tropolished until perforation using a Fischione Model 120Twinjet

Table 1

Composition and electronic parameters of Ti–(Nb,Ta)–Fe alloys produced by CCLM.Sample Ti (at.%)Nb (at.%)Ta (at.%)Fe (at.%)Bo Md (eV)e /a Ti–12Nb–5Fe 83.0012.00– 5.00 2.82 2.37 4.32Ti–7Ta–5Fe 88.00–7.00 5.00 2.81 2.38 4.27Ti–10Ta–4Fe

86.00

–

10.00

4.00

2.82

2.40

4.26

2 A.Biesiekierski et al./Acta Biomaterialia xxx (2015)

xxx–xxx

Electropolisher containing a solution of 6vol%perchloric acid,35vol%n-butanol and the remainder methanol,maintained at $à30°C through addition of liquid nitrogen.TEM analysis was then performed using a JEOL JEM 2000FX TEM.

Thermal analysis was performed under vacuum using a Mac Science DSC3200S Differential Scanning Calorimeter (DSC),with samples of mass 40–90mg;these were cut using a diamond saw,then ultrasonically cleaned in acetone and dried prior to use.A sin-gle complete analysis per sample was performed over the temper-ature range 100–500°C,with two full heating and cooling cycles utilizing a heating rate of 10°C min à1,and all experiments were duplicated to ensure consistency.Temperatures exceeding 520°C,even despite the nominally evacuated environment,induced sample oxidation suf?cient to in?uence the measured

of the cathodic Tafel slope to the measured E corr .In all cases,corro-sion rate was subsequently determined from I corr via the equivalent weight method outlined in ASTM:G102[29].The breakdown potential E B was determined via extrapolation to intersection of the linear current density slopes seen immediately prior to and succeeding the breakdown events.

Due to their osteoblastic properties [30],proliferation of human SaOS-2sarcoma cells was utilized to assess the biological response to the investigated Ti–(Nb,Ta)–Fe alloys in vitro .Small disc alloy specimens were polished to 600grit ?nish,and cells were cultured at 37 C in modi?ed minimum essential medium (MMEM)for 1and 7days,with commercially pure Ti (grade 2ASTM)plates as control.Cell proliferation was subsequently measured by using a colorimetric assay utilizing the reduction of a tetrazolium salt Fig.1.XRD Spectra of Ti–(Nb,Ta)–Fe alloys.

A.Biesiekierski et al./Acta Biomaterialia xxx (2015)xxx–xxx 3

Student’s T -test was subsequently performed to gauge respective signi?cance for each sample,with a threshold of P =0.05unless otherwise noted.

3.Results and discussion

with the behavior seen in the Ta-containing cases,where both alloys display a signi?cantly reduced b -phase lattice parameter,on the order of 3.26?;distinct from the earlier case,this agrees well with predictions from application of Vegard’s Law [33],and similarly matches behavior reported by Zaluski et al.in binary Ti–Fe alloys [34].The observed x -phase is similarly reduced,and

maintains the expected a x ????

2p a b ;c x ???3p a b relationship common to x derived from a parent b -phase [35].The same cannot be said for the a 0-phase,which shows only minor decrease relative to pure a -Ti,likely due to the poor solubility of Fe in this phase [34].Another interesting note is the absence of peaks consistent with the intermetallic phases FeTi or Fe 2Ti in any alloy;given the extre-mely poor solubility of Fe in the a -phase,some degree of Fe segre-gation and FeTi precipitation would be expected to appear,particularly in the a 0-rich Ti–7Ta–5Fe alloy [34].Their absence,then,would suggest that suf?cient b -phase stabilization was achieved by the Ta/Nb additions to retard further decomposition on the timescales seen in cooling.

Table 3

Detected phases and lattice parameters in Ti–(Nb,Ta)–Fe alloys as determined via XRD.Sample

Phases

Spacegroup

Measured lattice parameters (?)a =b

c Ti–12Nb–5Fe b

Im 3m 3.31 3.31Ti–7Ta–5Fe a

P 63=mmc 2.95 4.73b Im 3.26 3.26Ti–10Ta–4Fe

b

Im 3m 3.26 3.26x

P 6=mmm

4.62

2.83

Fig.2.Optical micrographs of etched Ti–(Nb,Ta)–Fe alloys.(A)Ti–12Nb–5Fe,(B)Ti–7Ta–5Fe and (C)Ti–10Ta–4Fe.

4 A.Biesiekierski et al./Acta Biomaterialia xxx (2015)xxx–xxx

grain boundaries,while larger pores up to1mm in size were rarely observed.Judging by their location and morphology,these are attributed primarily to dissolved gases in the melt.A similar struc-ture is seen in the Ti–10Ta–4Fe alloy,with only b-phase material being visible under OM.Minimal large-scale porosity(visible to the naked eye)is present,contrasting the Ti–12Nb–5Fe alloy, although microscale porosity could still be observed,and no den-dritic patterning was visible.Grain structure resembles the Ti–12Nb–5Fe alloy,though grains were predominantly de?ned by the crack-like secondary grain structure,rather than larger equiaxed-grain boundaries,with mean diameters of$180l m. The Ti–7Ta–5Fe alloy displayed similar porosity and large equiaxed grains,with diameter600?100l m,ingrown with acicular lathes of a0in the classical Widmanst?tten structure,with additional a0 precipitation growing from the grain boundaries in a lamellar structure.

SEM observation broadly matched that observed under optical microscopy.However,the greater magni?cation of SEM allowed for the elucidation of another structure;visible in the less heavily etched regions of both the Ti–12Nb–5Fe and Ti–10Ta–4Fe alloys,?ne bands of a secondary phase precipitation distinct from the dominant b-phase were resolvable.Images of these are provided in Fig.3(a)and(b).These bands could not be observed in the Ti–7Ta–5Fe alloy due to the decreased Ta-content,and hence b-stabilization,allowing pronounced martensitic precipitation throughout the alloy;whether this precipitation competes with, inhibits or merely masks the band structures is unknown,how-ever.These bands were comprised of a dense precipitation of?ne acicular particles threading the Ti–10Ta–4Fe alloy,which can be seen most clearly in Fig.3(b).Precipitation bands observed in the Ti–12Nb–5Fe alloy varied between this acicular structure and shorter,blunter precipitates.Their size,susceptibility to etchant attack and distribution throughout the larger b-grains suggests that these bands precede the crack-like features seen in highly-etched regions under OM.Besides the bands,similar precipitates could be found along a number of the larger b grain boundaries in the Ti–10Ta–4Fe material,as well as scattered throughout the bulk b-grains.Their acicular morphology and location is suggestive of a martensitic phase such as a0or a00,but given their relatively large size and frequency it is surprising that these phases are not observed via XRD.EDS measurements taken from the precipitation bands in each alloy revealed little;the composition varied by less than0.4%in any one species’concentration,in inconsistent direc-tions.As these swings fall below the intrinsic error of the individ-ual EDS measurements[36],variation is considered negligible. While this minimal?uctuation does not identify the phase,the absence of variation,particularly in Fe content,argues against either a=a0or FeTi intermetallic phases[34].

Building on earlier work concerning microstructure,TEM was performed across the three investigated alloys,again broadly matching earlier analyzes.Representative transmission electron micrographs for the three alloys are presented in Fig.4(a)–(d),with relevant Selected Area Electron Diffraction(SAED)patterns inset. The?rst of the three alloys,Ti–12Nb–5Fe is presented in Fig.4 (a),and displayed a b-phase dominated microstructure throughout the analyzed material,interspersed with frequent twinning bands. No resolvable diffraction spots associated with the x-phase could be observed,although faint streaks consistent with extremely small,diffuse x particles matching the known orientation relation-ship[2110]x//?011

b

,were detected[37].The Ti–7Ta–5Fe alloy

(Fig.4(b))displayed the classic Widmanst?tten a0tb structure, matching earlier results;analysis of SAED however,was hampered by the absence of favorable diffraction axes within the tilt range of the TEM sample holder,despite several attempts at sample analy-sis.The Ti–10Ta–4Fe again showed a predominantly b-phase struc-ture;however,x precipitation was pronounced enough to yield distinct x-phase diffraction maxima throughout the SAED pat-terns.This x diffraction pattern displayed classic Burgers orienta-tion relations relative to the matrix b-phase[37];in Fig.4(c),this is

visible as?0111 x

1

//[113]b,along with a twinned x variant,x2, aligned along the[0axis.Further faint maxima,for instance those falling between the(2020)re?ections of the two x patterns, can also be observed;these are the result of double diffraction between the x-twins,and so may be disregarded.The clear x diffraction maxima also allowed for the successful dark-?eld imag-ing of the phase,which reveals that the x-phase is densely precip-itated throughout the material in the form of approximately ovoid particles up to10nm in diameter.That the x-phase is both so large and frequent throughout the material agrees well with the clear x diffraction peak displayed via XRD analysis in Fig.1.

Of note,a sample of the Ti–12Nb–5Fe alloy was successfully imaged along one of the aforementioned precipitation bands;this is presented,along with the attendant SAED pattern,in Fig.4(d). While clear,strong maxima consistent with diffraction of b-Ti along the[012]b axis dominates,additional streaking is evident, e.g.extending between the(200)b and(121b)re?ections.Interpre-tation of these features is less straightforward than in earlier cases; this does not match a/a0-precipitation,nor that seen for the x-phase,which would be expected to yield re?ections falling between e.g.the(200)b and b)re?ections[35].However,mul-tiple variants of a00diffracted along the various permutations of the h111i axes could potential yield these features,given no variant being of suf?cient size/crystallinity to yield resolvable maxima [38].Given the similar nature of the precipitation in Ti–10Ta–4Fe,it is expected that analogous behavior is being observed in this alloy.Notably,while lacking a positive identi?cation,this hypoth-esis?ts the observed behavior;lack of an XRD signal is unsurpris-ing in a situation with nanoscale crystallites and lack of elemental variation under EDS precludes a=a0or FeTi precipitation.While this

Fig.3.Scanning electron micrographs of etched Ti–(Nb,Ta)–Fe alloys.(A)Ti–12Nb–5Fe and(B)Ti–10Ta–4Fe.

precipitation may potentially be x in the Ti–10Ta–4Fe,particularly given the relatively high proportion of the phase under both XRD and TEM,there is minimal sign of this phase in the equivalent Ti–12Nb–5Fe structure,despite similar morphologies.Further,a00 may form microscaled structures even in the as-cast condition [39].As a result,this precipitation is attributed to a00phase,with the precipitation bands possibly forming to accommodate stresses associated with cooling.Beyond this no other phases were observed,surprising given x-phase precipitation is known to pro-vide bountiful sites for a-phase nucleation at the b/x interface [37].

Mean EDS values for the alloys are given in Table4,and their approximate positions on the Bo–Md phase diagram are given in Fig.5.Observed values match,to within inherent accuracy of EDS [36],the nominal composition targets.Further,within each alloy, concentrations were typically uniform;the greatest variation was noted in the Ti–7Ta–5Fe case,due to the elemental segregation of b-stabilizing elements between the a0and b-phase structures. Interestingly however,this segregation was not total;although the b-phase displayed a higher mean Fe content,on the order of 4.2%,Fe contents of the a0phase averaged2.3%,well above the equilibrium concentrations for that expected in an equilibrium a-phase[34].

However,the observed structure of the alloys deviate from what might be expected from consideration of Fig.5,with exten-sive precipitation of a0in the Ti–7Ta–5Fe alloy,while both the Ti–10Ta–4Fe and Ti–12Nb–5Fe alloys showed x phase precipita-tion.Of these,the faint,diffuse contrast and plentiful twinning in the Ti–12Nb–5Fe alloy is consistent with a moderately b-stabilized material where x-phase precipitation is diffuse and transient,forming BCC twins on reversion to the b-phase[35].In the less-stabilized Ti–10Ta–4Fe,this x precipitation grows more pronounced,with both greater stability and precipitate size,lead-ing to a corresponding decrease in{112}h111i

b

twins.Coupled with the presence of a00in both of these alloys,it can be seen that all three alloys display incomplete b-stabilization.These differ-ences stem from the manufacturing process;the articles from

Transmission electron micrographs of Ti–(Nb,Ta)–Fe alloys.(A)Ti–12Nb–5Fe,(B)Ti–7Ta–5Fe,(C)Ti–10Ta–4Fe and(D)Ti–12Nb–5Fe grain boundary.

available.Please note;(C)taken under dark-?eld illumination,all other?gures are bright-?eld illumination.

Table4

Mean composition of as-cast Ti–(Nb,Ta)–Fe alloys as determined using EDS.

Sample Ti(at.%)Nb(at.%)Ta(at.%)Fe(at.%)

Ti–12Nb–5Fe82.8?0.312.6?0.3– 4.5?0.2

Ti–7Ta–5Fe89?1–7.4?0.6 3.6?0.8

Ti–10Ta–4Fe86.7?0.5–9.6?0.4 3.7?0.4

Error values represent1std.dev.

which Fig.5is recreated assume the material is in an as-quenched state following solution treatment,with literature alloys decom-posing into a duplex a tb -phase structure on ageing [22,23,26].In contrast to this,the alloys presented in this work are investi-gated in the as-cast state which,in CCLM manufacturing,is inter-mediate between solution-treated-and-quenched and aged or furnace-cooled process routes,yielding a similarly intermediate microstructure.Unfortunately,this precipitation prevents the development of truly low-E alloys in the as-cast state.Neverthe-less,this does suggest the material is highly responsive to heat-treatment;a useful property given the proscription against mechanical working entailed by a porous structure.3.2.Thermal analysis

Further investigation in this regard was performed with DSC;during the initial heating trace,distinctive peaks were observed in the Ti–12Nb–5Fe and Ti–10Ta–4Fe alloys,at 339°C and 360°C respectively,and a local maxima at 359°C was seen in the Ti–7Ta–5Fe alloy.These are associated with the b !x transition,which was recently identi?ed as exothermic [40–43].Subse-quently,all three alloys displayed an increasingly endothermic trend,as expected of Ti-alloys approaching the b -transus behavior [40];due to limits of analysis,the actual transus was only resolved in the Ti–7Ta–5Fe alloy,at 483°C.On cooling,only one feature of note was observed;a small exothermic peak falling at $502°C in the Ti–12Nb–5Fe alloy.

During the second heating cycle,pronounced exothermic peaks were no longer observed at $360°C.Instead,broad exothermic peaks were detected centered on $400°C,matching existing liter-ature describing the precipitation of a =a 0at temperatures above $400 C [40,43].Beyond this point,heating traces became more endothermic due to the approaching b -transus,but this appeared slightly shifted to higher temperatures;likely a result of segrega-tion of the b -stabilizing elements (particularly Fe)due to the effec-

tive overaging engendered by the DSC analysis causing the

remaining a 0-phase to become increasing depleted of stabilizing species,raising the b -transus temperature [12].This may also explain the small exothermic peak seen during only the ?rst cool-ing cycle of the Ti–12Nb–5Fe,as the reverse b !a =a 0transition undergoing a similar shift to higher temperatures outside the lim-its of observation,although further work is needed to con?rm.In either case,it is apparent that the alloy shows a pronounced response to heat-treatment,suggesting the potential for increased versatility in its application.

3.3.Mechanical analysis

Values for the Vickers microhardness for the three alloys are presented in Table 5.The highest mean hardness is seen in the Ti–7Ta–5Fe alloy,and the lowest overall hardness in the Ti–12Nb–5Fe alloy;variation within a given alloy was minimal.These values are consistent with the observed microstructure,with the presence of extensive a 0or x -precipitation in the Ta-containing alloys increasing hardness relative to the Ti–12Nb–5Fe.

Results from monotonic compression are presented in Table 6,and representative monotonic compression curves are shown in Fig.6.Of the three alloys,Young’s modulus (E )was highest in the Ta-containing alloys,particularly in the x -rich Ti–10Ta–4Fe.Only the Ti–12Nb–5Fe alloy displayed an E falling below the accepted range for C.P.Ti (P 103GPa)[1],at approximately 86GPa.As this still remains above the expected range for human bone [10],it is clear that no investigated alloy would be suitable solely for use as a non-porous implant material.Observed r YS val-ues,determined by intersection with the 0.2%offset slope,were greatest in the Ti–10Ta–4Fe alloy,though both Ta-containing alloys exceeded 1.2GPa.The r YS was signi?cantly less in the Ti–12Nb–5Fe alloy,at 740MPa.Despite this,the Nb-containing alloy still displays a r YS more than twice that of Grade 2C.P.Ti,compa-rable with Ti–6Al–4V in the solution-treated and aged state [1,12].Combined with the measured moduli,d in all alloys greatly exceeded that of bone [10].Compressive strength (r c )for all alloys,where observed,was in excess of a gigapascal.Due to only a single de?ned failure point being observed in the Ti–12Nb–5Fe alloy,with other test specimens displaying a malleable deformation until termination of the experiment,no exact value may be given for this composition.As might be expected given the observed r YS ,the Ta-containing alloys consistently displayed higher compressive strength than the Nb-containing alloy,although the highest r c was observed in the Ti–7Ta–5Fe,rather than Ti–10Ta–4Fe,alloy.The degree of plastic deformation noticeably varied between samples,with the Nb-containing alloy showing a far greater plastic strain at failure,to the point that no failure was observed up until the termination of the experiment in the majority of runs.This contrasts to the Ta-containing alloys,which displayed plastic strains on the order of $6%for the Ti–7Ta–5Fe alloy,and $13%for the Ti–10Ta–4Fe alloy;together,these yield total failure strains ( fail )of P 42,7?1and 14?4.While caution must be exercised in comparing to tensile results,this extent rivals many existing Ti-based biomaterials [1,17].

As with the hardness,the observed E ,r YS ;r c and fail appear to be directly attributable to microstructure;namely,the absence of signi?cant x and a 0precipitation throughout the Ti–12Nb–5Fe

for Ti–(Nb,Ta)–Fe alloys,reconstructed from information ?gure shows nominal and actual electronic parameters,Dashed ellipses denote approximate range of measured Table 5

Vickers microhardness of as-cast Ti–(Nb,Ta)–Fe alloys.Sample Hardness (HV)Std.dev.(HV)Ti–12Nb–5Fe 293?6Ti–7Ta–5Fe 430?10Ti–10Ta–4Fe

410

?10

A.Biesiekierski et al./Acta Biomaterialia xxx (2015)xxx–xxx 7

alloy,and said x and a 0phases’s presence in Ti–10Ta–4Fe and Ti–7Ta–5Fe alloys,respectively.Given both x and a 0are well known for their adverse effects on these [12],increasing their volume frac-tion would be expected to noticeably impinge on the observed properties.

3.4.Corrosion analysis

Corrosion analysis of these materials was performed via open-circuit potential (E OC )vs time (T )and potentiodynamic polarization measurements.The E OC vs time measurements indicated that the investigated alloys showed an increase in the E OC with time which rapidly ($1h)settled into close accordance with an inverse power law of form:

E OC ?

1

eC 1T tC 2T

tE OC 1

e1T

where T is elapsed time,C 1and C 2are descriptive constants,and E OC 1represents the potential approached by E OC as T !1.The characteristic values for the three analyzed alloys,along with C.P.Ti and Ti–6Al–4V controls,were derived via regression of

the

measured curves to this function;these values are given in Table 7,and are plotted in Fig.7along with the measured data.

The gradual,slowing increase of E OC over time is common in materials that show a gradual passivation behavior [44];in the case of Ti-containing alloys,this is typically attributed to the for-mation of an impermeable TiO x ?lm,and clearly proceeds even despite the Fe addition [45,46].Notable also is that the vast major-ity of variation in surface potential occurs within the initial 24h,suggesting passivation behavior is relatively rapid within the bio-logical environment.The Nb-containing alloy appears to most rapidly reach a neutral open-circuit potential;this,coupled with its higher initial E OC values would suggest an ability to better resist corrosion in the immediate aftermath of a disruption of the passive oxide layer.However,as elapsed time approaches 72h,the E OC value is equaled by both C.P.Ti and the Ta-containing alloys.Extrapolating the observed trend to in?nity,it would appear that these alloys tend toward a higher,and hence more noble,overall E OC 1value.It should be stressed at this point,however,that given that the E OC 1values are derived by regression over a ?nite period,these interpretations are at best hypothetical;still,this would appear to suggest that these Ta-containing alloys are perhaps bet-ter suited to applications minimizing abrasion of the implant surface.

More de?nitive is the information provided via the potentiody-namic polarization analysis of the samples over the anodic and cathodic range.Representative polarization traces for the three alloys and two standards are presented in Fig.8,while key param-eters determined from this are presented in Table 8.The linear polarization measurements show minimal variation among these alloys;namely there is no statistically signi?cant difference in E corr values between the three Fe-containing alloys,either among the three,or with respect to the C.P.Ti control.This uniformity speaks well toward the potential application of the Ti–(Nb,Ta)–Fe alloys in a clinical role.Further,there is a statistically signi?cant increase (P <0.05for Ti–7Ta–5Fe,P <0.01for all others)in E corr for all alloys relative to the Ti–6Al–4V control,unsurprising given reported Ti–6Al–4V E corr values in comparable literature cases [45–48].I corr values of the investigated alloys display a broader range of values,however;while the Ti–12Nb–5Fe alloy displays I corr values indistinguishable from the C.P.Ti control,a signi?cant (P <0.05)increase is seen in the Ti–10Ta–4Fe alloy relative to both controls.By contrast,the Ti–7Ta–5Fe alloy displays a far lower cor-rosion current,on the order of 0.08l A/cm 2;a value in close agree-ment with that observed in the similarly a +b duplex-structured Ti–6Al–4V control.It should be noted,however,that despite the dif-ference in I corr values between the alloys,the observed corrosion rates showed no statistically signi?cant difference to an existing control,whether either the C.P.Ti (Ti–12Nb–5Fe,Ti–10Ta–4Fe)or Ti–6Al–4V (Ti–7Ta–5Fe)controls.

Increasing the applied voltage above E corr ,the majority of alloys lacked the classical anodic branch in the Tafel region.Rather,a region extending $0.1–0.3V immediately above E corr was seen,that showed negligible increase in current density.This is best demonstrated in the Ti–6Al–4V control,but similar behavior could be observed in all alloys,and is the result of a dense initial passive layer [47].However,among the designed alloys,this behavior was not consistent;while most prevalent in the Ti–10Ta–4Fe alloy,all Fe-containing alloys showed traces where a traditional anodic slope started to form,but at $0.2–0.5V above E corr ,sharply reverted to the behavior displayed by the Ti–6Al–4V alloy,in some case yielding as much as an order of magnitude decrease in current density.This is indicative of repassivation prior to the subsequent breakdown in the transpassive region above the breakdown poten-tial (E B )[49].This repassivation can occur due to reduction of the extant passive ?lm from exposure to a polarization current

0.3–0.5V below E corr ;while in Ti–6Al–4V and C.P.Ti controls this

Table 6

Mechanical properties from compressive testing.Fig.6.Representative monotonic compression curves of Ti–(Nb,Ta)–Fe alloys.

Table 7

Characteristic parameters from open-circuit potential vs time measurements for Ti–(Nb,Ta)–Fe alloys.Sample C 1(V à1T à1)C 2(V à1)E OC 1(V)C.P.Ti

à7.592?0à5à3.4620.055Ti–6Al–4V à1.737?10à4à3.360.022Ti–12Nb–5Fe à3.598?10à4à7.400.025Ti–7Ta–5Fe à1.058?10à4à2.4150.046Ti–10Ta–4Fe

à6.266?10à5

à3.676

0.068

8

A.Biesiekierski et al./Acta Biomaterialia xxx (2015)xxx–xxx

passive layer could resist the reducing environment,the same could not be said for the Fe-containing alloys.These repassivation events were seen in80%of scans for the Ti–7Ta–5Fe alloys,40%of the Ti–10Ta–4Fe scans and25%of the Ti–12Nb–5Fe scans.As a crude metric performed in the absence of direct cathodic polariza-tion,this would appear to suggest an increasing susceptibility to cathodic erosion of the passive?lm with an increasing ratio of Fe to other alloying elements;clearly,however,further work would be necessary to con?rm this.Cathodic resistance is further consid-ered a lesser concern,given the expected environment within the body falls into the anodic regime[14].

In the transpassive region the current begins to increase,with this rate of increase rising sharply above E B(listed in Table8), yielding as much as a two order of magnitude rise in current den-sity.The Fe-containing alloys’E B values of between+0.6and0.7V SCE equal or exceed that observed for Ti-based alloys in similar conditions[47,50],indicating improved resistance to anodic corro-sion potentials.Among the alloys,the Ti–10Ta–4Fe alloy exhibits the highest E B,but given the greater variance in the measured val-ues for this alloy,ultimately no meaningful difference can be dis-cerned between any of the three.Beyond this,these materials all show similar behavior.The brief passive region occurring in the

Fig.7.Actual and LoBF E OC vs T curves for as-cast Ti–(Nb,Ta)–Fe alloys in HBSS.

Fig.8.Potentiodynamic polarization curves for the as-cast Ti–(Nb,Ta)–Fe alloys in HBSS.

Table8

Electrochemical corrosion parameters derived from potentiodynamic polarization of Ti–(Nb,Ta)–Fe alloys.

Sample E corr(V)I corr(l A/cm2)Corr.rate(l m/year)E B(V)

C.P.Tià0.25?0.040.2?0.12?10.27?0.07

Ti–6Al–4Và0.402?0.0070.07?0.030.6?0.20.4?0.1 Ti–12Nb–5Feà0.23?0.060.2?0.12?10.6?0.2 Ti–7Ta–5Feà0.25?0.090.08?0.050.5?0.40.6?0.1 Ti–10Ta–4Feà0.24?0.050.5?0.23?10.7?0.3

Note:all voltages presented relative to SCE.Listed error values represent1std.dev.

neighbourhood of 0.5–1.0V SCE matches that seen elsewhere in literature,as well as in earlier studies by this author [46,47,50];this is associated with the formation of a second,less passive ?lm.Increasing voltage further,transient current spikes occurred,heralding the likely onset of pitting corrosion [46,47].This behav-ior shows little difference between the control and investigated samples,with current densities for all investigated alloys typically

the Ti-control and all three Ti–(Nb,Ta)–Fe alloys for the 1day sam-ples;cell distribution consisted of broad swathes of little or no cell growth,interspersed with clumps with a moderate to high cell density.Cell clusters appeared marginally more diffuse in the Ti–7Ta–5Fe sample,although given the small sample size,this is not considered meaningful.Across all materials,cells predominantly displayed one of two morphologies;a spherical appearance with micrographs of SaOS-2cells cultured on Ti–(Nb,Ta)–Fe alloys:7-day incubation.(A)Ti–12Nb–5Fe,(B)Ti–7Ta–5Fe,(C)Ti–10Ta–4Fe and 10 A.Biesiekierski et al./Acta Biomaterialia xxx (2015)xxx–xxx

suitable surface chemistry and topography to adhere to,these are highly sensitive to adverse conditions.As such,their extensive presence,along with well attached,highly extended and thickly clustered cell masses is indicative of minimal to no adverse biolog-ical impact from the alloy substrates [31,51,52].

Beyond confocal microscopy,MTS assays were also performed to provide a more quantitative analysis of cell proliferation and survival ability;resulting data is presented in Fig.10,with the ver-tical scale normalized to a control well seeded with SaOS-2cells in the absence of a metal disc.Measured absorbance showed minimal variation;relative to both the empty well and C.P.Ti controls,none of the analyzed alloys showed a statistically signi?cant (P >0.05)difference in absorbance.Relaxing the level of con?dence to 90%,a slight reduction can be seen relative to the empty well control in the Ti–10Ta–4Fe alloy;however,even at this reduced level,this alloy was indistinguishable from the C.P.Ti control,suggesting all investigated alloys at least rival existing materials with respect to cellular response.As might also be expected from this data,differ-ences between the alloys were likewise minimal;to 95%con?-dence,no alloys show a statistically signi?cant variation in absorbance.

4.Summary

In conclusion,the investigated alloys all show a signi?cant degree of promise with regard to their eventual application as potential orthopedic biomaterials.Displaying elastic moduli of 90–120GPa,on par with Ti-based alloys currently in biomedical service [1],limits their value in the manufacture of high density,solid implants.Rather,the value of these implants would appear to lie instead in the creation of a porous material;all three alloys displayed 0.2%offset yield strengths at least comparable to that observed in existing high-strength Ti-based alloys such as Ti–6Al–4V or Ti–6Al–7Nb [12].In the case of the Ti–Ta–Fe alloys,yields in excess of 1.25GPa were achieved;phenomenal values,exceeding those seen even in cold-worked Gum metal or TNTZ [53,3].As a result of this,the calculated admissible strain of the alloys,at values ranging from 0.9%in the Ti–12Nb–5Fe alloy,to 1.2%in the Ti–7Ta–5Fe alloy,not only exceeds the 0.55%seen in cortical bone but furthermore rivals that seen in next generation

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