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Materials Science and Engineeri

Fatigue and wear evaluation of Ti-Al-Nb alloys for biomedical applications

C.J. Boehlert a,⁎, C.J. Cowen a, J.P. Quast a, T. Akahori b, M. Niinomi b

a Department of Chemical Engineering and Materials Science, Michigan State University, East Lansing, MI, 48824, USAb Materials Development Division, Biomaterials Science, Institute for Materials Research, Tohoku University, Japan

Available online 11 April 2007

Abstract

In this work the fatigue and wear behavior of Ti–15Al–33Nb(at.%) and Ti–21Al–29Nb(at.%) was evaluated and compared to that for othertitanium-based biomedical implant alloys, in particular Ti–6Al–4V(wt.%). Fatigue stress versus life curves were obtained for tests performed atroom temperature in air at a stress ratio of R=0.1 for maximum stresses between 75%–90% of the ultimate tensile strength. The results indicatedthat the fatigue strength and lives of the as-processed alloys are comparable to that for Ti–6Al–4V(wt.%). Heat treatment significantly increasedthe orthorhombic-phase volume fractions in the alloys and resulted in reduced fatigue strength. The wear resistance for the alloys was significantlygreater than that for Ti–6Al–4V(wt.%). Based on the current results, it is proposed that titanium–aluminum–niobium alloys will be ofconsiderable future interest for biomedical applications.© 2007 Elsevier B.V. All rights reserved.

Keywords: Biocompatibility; Titanium; Wear; Hardness; Fatigue

1. Introduction

Erosive bone diseases are a growing medical problem that canresult in the need of replacement surgery with bone-substitutematerials. Some of the important characteristics of bone-substitute materials include biocompatibility with near-bonehard and soft tissue, an elastic modulus near that for bone, andtensile and compressive strength and fracture toughness equal toor greater than that of bone. In addition, the fatigue strength andwear resistance of the material should guarantee a safe operationof the implant during the expected period of use.

Degradation of artificial implant materials due to high wearrates can lead to unfavorable biological effects on bone densityand implant fixation, resulting in shorter lifetimes. Implant wearalso leads to significant patient discomfort. For these reasons, wearresistance is an important property when considering a suitableimplant material. The wear resistance of a material is directlyrelated to the contact stresses and friction at the point of contact.

The friction between artificial materials is much higher thanthat between healthy, natural joints. This higher friction isattributed to the high rigidity of artificial materials leading tonon-recoverable wear [1]. The most common wear stems from

⁎ Corresponding author. Tel.: +1 517 353 3703.E-mail address: boehlert@egr.msu.edu (C.J. Boehlert).

0928-4931/$ - see front matter © 2007 Elsevier B.V. All rights reserved.doi:10.1016/j.msec.2007.04.003

the deterioration of ultra high molecular weight polyethylene(UHMWPE) in contact with a metal alloy, as frequently seen inartificial joints. Studies have shown that excessive wear ofUHMWPE and/or metal leads to long-term failure [2,3]. Thedebris accumulated during wear can also lead to adverse effectssuch as inflammation, release of damaging enzymes, osteolysis,infection, pain, and implant loosening [4].

The wear of UHMWPE from contact with titanium (Ti) alloyshas been attributed to the mechanical instability of the metaloxide layer [5]. At high enough shear stresses, the oxide layer isbroken down, exposing the unoxidated surface below. Thissurface either reforms an oxide layer or adhesively bonds to theUHMWPE, leading to continual deterioration of the polymerand/or metal [1]. While this is occurring, the contact surfaces arecontinually increasing in roughness, resulting in furtherbreakdown of the artificial joint materials. Free debris,composed primarily of the eroded oxide surface, is known as athird body abrasive component that also increases the wear rate.

Commercially pure titanium (CP Ti) and α/β type Ti alloysare currently widely used as structural biomaterials for thereplacement of hard tissues in devices such as artificial hipjoints and dental implants. This is due to their excellent specifictensile and fatigue strength and electrochemical corrosionresistance in addition to their exceptional biocompatibilitycharacteristics (i.e. benign biological responses) among metallic

1 Throughout the remainder of the paper these alloys will be designated asTi–15Al–33Nb and Ti–21Al–29Nb and Ti–6Al–4V(wt.) will be designated asTi–6Al–4V.

Table 1Compositions of the studied Ti–Al–Nb alloys

Alloy Ti, at.% Al, at.% Nb, at.% Fe, ppm O, ppm N, ppm

Ti–15Al–33Nb 51.4 15.3 33.3 110 1100 100Ti–21Al–29Nb 50.7 20.6 28.7 2000 790 110

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biomaterials. In particular, the specific alloy most widely usedfor implants is Ti–6Al–4V(wt.%) extra-low interstitial (ELI)because of its excellent biocompatibility and its combination ofhigh specific strength, fracture toughness, fatigue and corrosionresistance, low density, ductility, and elastic modulus, oxi-dation resistance, and conventional processability. However, itselastic modulus is significantly greater than that for boneand vanadium (V) is potentially toxic in elemental form [6].Therefore, alloying elements other than V are currently beingexamined.

β-type titanium alloys are currently being developed in orderto obtain low rigidity, which is considered effective forpromoting bone healing and remodeling. Such alloys, com-posed of non-toxic and non-allergic elements, exhibit goodfatigue resistance, and workability [7–10]. Examples of thesealloys include Ti–15Mo(wt.%) [11], Ti–35Nb–7Zr–5Ta(wt.%)[7], and Ti–3Al–2.5V(wt.%) [8] and Ti–29Nb–13Ta–4.6Zr(wt.%) [12–17]. Although the rigidity of α/β type titaniumalloys is less than that of Co–Cr type alloys and stainless steelsused for biomedical applications, it is still considerably greaterthan that of the cortical bone [9].

Several investigations have included newer alloys based onthe Ti–Al–Nb system, such as Ti–6Al–7Nb(wt.%) [Ti–10.5Al–3.6Nb(at.%)] where Nb replaces V [18–26]. Substitu-tion of Nb for V is attractive as, depending on the concentrationof Nb, this does not result in degradation of several mechanicalproperties. In the current work Ti–Al–Nb alloys containinglower Ti concentrations, in particular Ti ∼50 at.%, wereevaluated to peruse their attractiveness for biomedical implantapplications. Fatigue strength and wear resistance are twoimportant mechanical parameters which are needed in order toassess the reliability of a material in medical implantapplications, as an implant has to withstand not only one-time peak stresses but also several million load cycles which itusually experiences during its lifetime. In this regard, thefatigue and wear behavior of Ti–21Al–29Nb(at.) and Ti–15Al–33Nb(at.%) were evaluated and compared to that forother Ti-based biomedical alloys. Microstructure-propertyrelationships of these alloys are also discussed. It is notedthat these two alloys have exhibited attractive biocompatibilitycharacteristics compared to Al2O3 and CP Ti as measuredthrough cytotoxicity, T-cell, and histological experiments [27],and recent unpublished work in our laboratory has qualitativelyshown that cell adhesion for these alloys is comparable to thatfor Ti–6Al–4V(wt.%). It is also noted that others have foundsome detrimental side effects due to the presence of Al ions,and thus Al alloy additions may be a concern for biomedicalapplications [28,29].

2. Experimental procedures

2.1. Materials processing

The two titanium–aluminum–niobium alloys were doublevacuum arc remelted in a vacuum furnace. The melts were castinto 15 cm diameter ingots (weighing 27 kg), which were thenmultistep forged to approximately 14 cm by 7 cm by 58 cm.

Approximately 1.6 cm thick slabs were then cut from thetransverse sections of the ingots and hot rolled to 0.17 cm thickby 7 cm wide sheets. The hot rolling produced a reduction inarea of 90% and introduced a total effective true shear strain onthe order of two. Identical hot-working temperatures, whichwere always maintained below the BCC-phase transustemperature, were used for the forging and rolling operationsfor a given material. The Ti–15Al–33Nb(at.%) and Ti–21Al–29Nb(at.%)1 sheets were hot rolled at 899 °C and 982 °C,respectively. After the hot rolling was complete, the sheets wereannealed at their respective rolling temperatures for one hourand then air-cooled. All alloy processing was performed at RMITitanium Company, Niles, Ohio. The chemical compositions ofthe rolled sheets, listed in Table 1, were determined usingInductively Coupled Plasma Optical Emission Spectroscopyand Inert Gas Fluorescence. The samples were evaluated both inthe as-processed (AP) condition as well as after the followingheat treatment (HT) performed in flowing argon: 1005 °C/3 h/FC at 10 °C per minute to 855 °C/8 h/CC at 1 °C/min to 650 °Cthen FC to room temperature (RT). A schematic illustrating thisheat treatment schedule is given in Fig. 1.

2.2. Microstructural characterization

Characterization of the microstructures was carried out usingoptical microscopy, scanning electron microscopy (SEM), andX-ray diffraction analysis. Samples were mounted in Konduc-tomet™ then ground using successively finer grits of SiCpaper. After grinding, the samples were polished using diamondpaste, according to the following schedule: 30 μm for 5 min,15 μm for 5 min, 6 μm for 5 min, 1 μm for 20 min. Colloidalsilica with an average particle size of 0.06 μm was used for thefinal polish, with polishing times ranging up to two hours.High-contrast digitized backscattered-electron (BSE) SEMimages, obtained using both an AMRAY 1810 SEM andPhillips 515 SEM, were used for volume fraction analysis ofthe BCC, orthorhombic (O), and hexagonal close packed(HCP) α2 phases present. Approximately twenty BSE imageswere taken at different magnifications and from differentsections of each sample. These images were then analyzedusing NIH Image Analysis Software to measure the areafraction of the phases containing different contrast due todifferent chemical contents. Average equiaxed grain size wasdetermined using the ASTM standard E112 [30]. Similar to the

Fig. 1. A schematic of the heat-treatment schedule used on the as-processedalloys. Note that 1005 °C is above the BCC transus for Ti–15Al–33Nb andbelow the BCC transus for Ti–21Al–29Nb.

Table 2Average phase volume fractions and grain size

Alloy Heattreatment,°C

α2 volumefraction

O volumefraction

BCCvolumefraction

Grainsize⁎,μ m

Ti–15Al–33Nb AP 10 0 90 3Ti–15Al–33Nb 1005 °C 1 44 55 120⁎⁎

Ti–21Al–29Nb AP 4 0 96 3Ti–21Al–29Nb 1005 °C 5 78 17 12

⁎ Measured independent of phase except in the case of ⁎⁎ where the grain sizerepresented the prior-BCC grain size for this supertransus heat-treated sample.

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phase volume fraction determinations, several micrographswere taken at different magnifications from different sheetorientations for the grain size calculations.

2.3. Fatigue testing

Flat dogbone-shaped specimens were machined from the APsheets with their longitudinal directions parallel to the rollingdirection. The geometry of the fatigue test specimen, which isidentical to that used for other investigations of biomedical Tialloys [13], is depicted in Fig. 2. All the machined specimenswere polished using a sequentially finer grits of SiC paper witha final polish using 0.06 m colloidal silica to produce mirrorsurfaces.

Fatigue tests were performed at a frequency of 10 Hz with astress ratio of R=0.1 in air at 295 K using an electro–servo–hydraulic fatigue-testing machine [13]. The maximum appliedstresses ranged between 820–950 MPa (N0.86 UTS) for the APmaterial and 700–940MPa (N0.8 UTS) for the HTmaterial, and1×107 cycles were considered to be “run out” where after theunbroken sample was removed. The fracture surfaces of thefatigue specimens were evaluated using SEM.

Fig. 2. A schematic of the fatigue specimen geom

2.4. Wear and hardness testing

To determine the wear resistance of the Ti–15Al–33Nb, Ti–21Al–29Nb, and Ti–6Al–4V alloys, in vitro friction wear testswere performed. This test method utilized a 5 mm diameterzirconia ball contacting the rotating metallic alloy disk at aconstant load of 4.9 N and a constant velocity of 60 rpm,covering 31.4 mm/s. The comparisons between materials weremade based upon the amount of subsequent weight loss. Thesetests were performed in Ringer's solution at 310 K. Thespecimen dimensions were 20×20×2 mm and the rotatingradius was 10 mm. Rockwell C hardness measurements weremade on both the wear surfaces and the nonwear surfaces. Fivehardness values were performed in each case using a 150 kgload for 10 s.

3. Results and discussion

3.1. Microstructure

Table 2 lists the measured phase volume fractions and grainsizes for the microstructures evaluated. SEM images of the APTi–21Al–29Nb and Ti–15Al–33Nb microstructures are shownin Fig. 3a and b. X-ray diffraction observations on super-transussolutionized and water-quenched microstructures indicated thatthe BCC phase in Ti–21Al–29Nb was ordered (B2) [31], whileit was disordered (β) in Ti–15Al–33Nb [32]. The orderingbehavior in Al-rich Ti–Al–Nb alloys has been documentedpreviously [31–37]. The AP alloys were composed almostentirely of the BCC phase, see Table 2. SEM images of the HT

etry used. All values are given in millimeters.

Fig. 3. BSE SEM images of the as-processed (a) Ti–21Al–29Nb and (b) Ti–15Al–33Nb microstructures. The dark phase is α2. The lighter matrix region is BCC.

Fig. 4. BSE SEM images of the heat-treated (a) Ti–21Al–29Nb and (b) Ti–15Al–33Nb microstructures. The darkest phase is α2. The lightest phase is BCC. The greyphase is O.

Fig. 5. Maximum applied stress versus fatigue life for the as-processed (AP) andheat-treated (HT) Ti–21Al–29Nb and Ti–15Al–33Nb alloys. Note that run outwas 1×107 cycles.

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Ti–21Al–29Nb and Ti–15Al–33Nb microstructures are shownin Fig. 4a and b. The heat-treatments resulted in greater overallgrain sizes, greater volume fractions of the O phase, and lowervolume fractions of the BCC phase.

3.2. Fatigue behavior

Fig. 5 compares the stress-life (S–N) behavior for the AP andHT Ti–21Al–29Nb and Ti–15Al–33Nb samples. On average,the AP samples maintained higher fatigue lives than the HTsamples. This may have been related to the increased RT tensilestrength, which is particularly important for high-cycle fatigueand is in agreement with recent correlations between tensilestrength and high-cycle fatigue behavior for a Ti–22Al–27Nb(at.%) alloy [38]. The S–N behavior of other Ti alloys,including conventional biomedical Ti–6Al–4V(wt.%) ELI andTi–6Al–7Nb(wt.%), obtained from the literature [26,27,39,40],are shown in Fig. 6 in comparison to the AP Ti–15Al–33Nband Ti–21Al–29Nb samples. The Ti–15Al–33Nb and Ti–21Al–29Nb alloys exhibited greater fatigue strength than Ti–6Al–7Nb(wt.%), while the fatigue lives were comparable tothose for Ti–6Al–4V(wt.%) at a given maximum applied stress.It is noted that the actual compositions of the current alloys inweight percent was Ti–9.8Al–47.0Nb(wt.%) [Ti–21Al–29Nb]

and Ti–6.9Al–51.7Nb(wt.%) [Ti–15Al–33Nb]. Thus theadditional Al and Nb concentration at the expense of Tiimproved the fatigue life as evident in the comparison with Ti–

Fig. 6. Maximum applied stress versus fatigue life for the as-processed (AP) Ti–21Al–29Nb and Ti–15Al–33Nb alloys and Ti–6Al–4Vand Ti–6Al–7Nb. Notethat run out was 1×107 cycles.

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6Al–7Nb(wt%). This is considered to be related to the greaterRT tensile strength exhibited by the Ti–21Al–29Nb and Ti–15Al–33Nb alloys [32] compared with Ti–6Al–7Nb [25].

The good fatigue strength of the investigated microstructuresis considered to be due to the balance of strength and ductilitybrought by the O and BCC phases [32] which leads to increasedresistance to fatigue crack initiation and small fatigue crack

Fig. 7. Secondary electron SEM fractographs of as-processed Ti–21Al–29Nb samplinitiation region where the crack growth emanated from the specimen surface at theoverload regions.

Fig. 8. Secondary electron SEM fractographs of as-processed Ti–15Al–33Nb samp(b) stable crack propagation, and (c) overload regions.

propagation. It has been reported that, for the Ti alloys withrelatively fine microstructures, a large part of fatigue life isoccupied by the small fatigue crack initiation and propagationlife [41]. This may be one reason why the fine microstructuresexhibited good fatigue strength compared to the otherbiomedical Ti alloys. In every case, the fatigue cracks tendedto initiate at the specimen surface then propagate parabolicallytowards the inside of the specimen. Relatively wide striationswere observed in the stable crack growth region, and equiaxeddimples were observed in the overload or fast fracture region.Such fracture surface morphologies are observed generally forductile metallic materials. It is noted that a smaller area ofdimpled regions was observed for Ti–21Al–29Nb than for Ti–15Al–33Nb, which may have been a result of the lower εfvalues [32]. In addition, the fracture surface in the crackinitiation and stable crack growth regions was relatively flat forTi–21Al–29Nb (see Fig. 7) compared with the more tortuousTi–15Al–33Nb surfaces (see Fig. 8). Considerable facets wereobserved near the crack initiation site in both low and high cyclefatigue life regions for Ti–21Al–29Nb (see Fig. 7). In the fastfracture area, a mixed-mode fracture surface composed ofintergranular fractures, facets, and dimples were observed.Therefore, the fatigue crack initiation and propagation char-acteristics of Ti–21Al–29Nb resembled more of a brittlefracture than that for Ti–15Al–33Nb.

Surface slip traces were evident on all the AP fatiguesamples, see Fig. 9a and b. Such observations indicated that the

es tested at RT in air for a maximum stress of 930 MPa in the (a) fatigue crackupper right-hand-side corner of the image, (b) stable crack propagation, and (c)

les tested for a maximum stress of 865 MPa in the (a) fatigue crack initiation,

Fig. 9. BSE SEM observations of surface slip traces for the as-processed fatigue samples: (a) Ti–21Al–29Nb; maximum cyclic stress=890 MPa, Nf=14,963 and (b)Ti–15Al–33Nb max. cyclic stress=865 MPa, Nf=37,972cycles.

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slip was wavy within the ductile BCC phase and slip did transferfrom the BCC phase to either the O phase or α2 phase. Theductile behavior is expected to have had a beneficial effectespecially within the low-cycle fatigue regions where crackpropagation is expected to dominate fatigue lives. The slip wasmore evident within the BCC phase than either the O or α2

phases. Thus the BCC phase may play a significant role inpromoting low-cycle fatigue resistance at RT for the Ti–Al–Nballoys, as low-cycle fatigue is dominated by crack propagationfor Ti alloys [42].

3.3. Wear behavior

The in vitro friction wear experimental results, illustratedin Fig. 10, show that the Ti–Al–Nb alloys exhibitedsignificantly more wear resistance than Ti–6Al–4V. Theweight loss of Ti–6Al–4V exceeded 30 mg, while weight lossof Ti–15Al–33Nb and Ti–21Al–29Nb alloys ranged from 2–9 mg. This is significant in terms of the relative wearresistance of materials targeted for biomedical application,and indicates the potential for Ti–Al–Nb alloys. It is notedthat a previous study demonstrated that the wear resistance of

Fig. 10. Weight loss for the three alloys after wear testing.

Ti–6Al–4V was approximately equal to that of Ti–6Al–7Nb(wt.%) [43]. SEM micrographs of the wear surfaces of eachof the alloys are illustrated in Figs. 11–13. From these imagesit is evident that the Ti–6Al–4V alloy exhibited a greaterwear path. Visible inspection of the wear path also indicatedthat it was significantly deeper than that for the Ti–Al–Nballoys.

Table 3 lists the average hardness values measured withinthe wear tracks and outside the wear tracks. It is evidentthat the wear surfaces were harder than the nonwearsurfaces and although Ti–6Al–4V was harder than the Ti–Al–Nb alloys, its wear resistance was significantly lower.This anomaly, which may arise from a difference in wearmechanisms between the alloys, remains unexplained at thispoint.

4. Summary and conclusions

The ambient-temperature fatigue and tensile behavior ofTi–21Al–29Nb and Ti–15Al–33Nb alloys was investigatedand compared to that for Ti–6Al–4V and other alloys in orderto evaluate newer Ti–Al–Nb alloys (i.e. other than Ti–6Al–7Nb(wt.%)) for potential biomedical implant applications.The main advantage that the Ti–Al–Nb alloys had over theTi–6Al–4V alloy was in wear resistance.

Fig. 11. Secondary electron SEM micrograph of the wear surface for Ti–6Al–4V.

Fig. 12. Secondary electron SEM micrograph of the wear surface for (a) as-processed and (b) heat-treated Ti–21Al–29Nb.

Fig. 13. Secondary electron SEM micrograph of the wear surface for (a) as-processed and (b) heat-treated Ti–15Al–33Nb.

329C.J. Boehlert et al. / Materials Science and Engineering C 28 (2008) 323–330

The following conclusions were made in this work:

1. The AP Ti–21Al–29Nb and Ti–15Al–33Nb microstructuresexhibited a majority of BCC phase and the grain sizes wererelatively fine, containing an average diameter of 3 μm. TheHT microstructures contained a larger grain size and agreater volume fraction of the O phase.

2. The chosen heat treatment resulted in lower fatigue lives thanthe AP alloys, which exhibited fatigue lives comparable tothose for Ti–6Al–4V(wt.%) ELI.

Table 3Hardnesses of the alloys in the wear tracks and outside the wear tracks

Alloy Condition Surface Rockwell C

Ti–6Al–4V AP Wear 44.2Ti–6Al–4V AP Nonwear 38.3Ti–15Al–33Nb AP Wear 37.5Ti–15Al–33N AP Nonwear 34.7Ti–15Al–33Nb HT Wear 39.1Ti–15Al–33Nb Ht Nonwear 35.4Ti–21Al–29Nb AP Wear 40.0Ti–21Al–29Nb AP Nonwear 37.6Ti–21Al–29Nb HT Wear 36.7Ti–21Al–29Nb HT Nonwear 37.1

AP: as processed; HT: heat treated.

3. Surface crack initiation, striations within the stable fatiguecrack propagation area, and equiaxed dimples on the fastfatigue crack propagation area were observed in low andhigh cycle fatigue life regions for both alloys. Ti–15Al–33Nb exhibited a ductile fracture surface morphology whileTi–21Al–29Nb tended to exhibit more of a brittlemorphology.

4. The wear resistance of the Ti–Al–Nb alloys was more thanthree times greater than that for Ti–6Al–4V.

5. Based on the data obtained on the fatigue and wearbehavior, the AP Ti–15Al–33Nb and Ti–21Al–29Nballoys merit further evaluation for biomaterial applications.

Acknowledgments

This work was partially supported by the National ScienceFoundation through grant DMR-0533954.

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