Titanium Alloys for Biomedical Applications

.elsevier.com/locate/msec

Materials Science and Engineerin

Titanium alloys for biomedical applications

H.J. Rack, J.I. Qazi *

School of Materials Science and Engineering, Clemson University, Clemson, SC 29634-0971, USA

Available online 12 October 2005

Abstract

Titanium alloys, because of their excellent mechanical, physical and biological performance, are finding ever-increasing application in

biomedical devices. This paper provides an overview of titanium alloy use for medical devices, their current status, future opportunities and

obstacles for expanded application. The article is divided into three main sections, the first discussing recent efforts focused on commercial purity

titanium. This is followed by considering effects of chemistry, grain size and a/h morphologies on mechanical properties of a+h alloys. Finally,

the third section reviews the status of metastable h alloys specifically designed for biomedical applications emphasizing their aging behavior and

its effects on mechanical properties.

D 2005 Elsevier B.V. All rights reserved.

Keywords: Biomaterials; Titanium alloys; Implants; Mechanical properties

1. Introduction

Continual aging of the US population has brought with it

an ever-increasing need for materials specifically suited for

bio-device application. For example, it is projected that

approximately 272,000 total hip replacements will annually

be performed by 2030. Additionally of the 152,000 total hip

replacements performed in 2000 approximately, 12.8%

involved revisions of previous hip replacements. The fact

that such a high percentage of hip replacements performed

every year are revision surgeries, although troubling, is not

surprising when the life expectancy of the implant versus the

ever-increasing life expectancy of the patient is considered.

Consistently, over 30% of those requiring total hip replace-

ments have been below the age of 65 and even those over

the age of 65 now have a life expectancy of 17.9 years.

Moreover female patients, who make up the majority of

those receiving total hip replacements, have a life expectancy

of 19.2 years at the age of 65. With a normal implant

longevity of 12 to 15 years, the majority of those that receive

hip implants at age of 65 will require at least one revision

surgery.

0928-4931/$ - see front matter D 2005 Elsevier B.V. All rights reserved.

doi:10.1016/j.msec.2005.08.032

* Corresponding author. Kemet Electronics Inc., 2835 KEMET Way

Simpsonville, SC 29681, USA. Tel.: +1 864 228 4442; fax: +1 864 228 4264.

E-mail address: [email protected] (J.I. Qazi).

Various metallic materials have been used for total hip

replacements as well as other joint replacement surgeries, i.e.,

knees, shoulders. Additional applications include trauma and

spinal fixation devices, cardiovascular stents, and, most

recently, replacement spinal discs. The material list includes

stainless steel, Co–Cr–Mo alloys, titanium alloys and other

more specialized alloys, e.g., Au–Pd. Of these titanium

alloys, the subject of the present article, offers several

benefits, including lower elastic modulus, excellent corrosion

resistance and enhanced biocompatibility [1]. The former is

particularly important for hard tissue replacement where stress

shielding, a phenomenon where reabsorption of natural bone

and implant loosening arises because of the difference in

elastic modulus between natural bone and hard tissue implant,

is one of the primary causes requiring revision surgery [2].

Another well-documented and related cause is bone necrosis.

This phenomenon has been associated with wear debris

generated from articulating components at a tabular cup.

Such wear debris has been shown to migrate and position

itself at the bone-implant stem interface thereby further

promoting bone cell death.

The present review builds on several excellent prior

summaries [1,3–5] and shows that this arena remains a

fruitful area for titanium research and development. It begins

by examining recent efforts focused at enhancing long used

biomedical titanium alloys, i.e., commercial purity titanium

and Ti–6Al–4V, these having been adapted from the

g C 26 (2006) 1269 – 1277

www

200 nm

400 nm

(a)

(b)

Fig. 1. Micrographs showing the ultra fine grained structure in CP– titanium

grade 2 produced by (a) ECAP and (b) ECAP+rolling [6].

200

400

600

800

1000

2 3 4 5 6 7 8

lg N

σ, M

Pa

CG

#1

#2

#3

*

*

*

+

+

+

Fig. 2. Fatigue response of ultra-fine grained grade 2 commercial purity

titanium with (CG) coarse grained, (#1) equiaxed cellular, (#2) elongated and

(#3) sub-grain microstructures [6].

H.J. Rack, J.I. Qazi / Materials Science and Engineering C 26 (2006) 1269–12771270

aerospace community, and then discusses newer metastable

beta titanium alloys specially designed for biomedical

application.

2. Commercial purity titanium alloys

Commercial purity titanium has long been used for

biomedical devices, for example cardiovascular stents, lead

wires and spinal/trauma fixation devices. When maintained at

low Fe content there is little concern about adverse interaction

between the implant and the human body. Notwithstanding

these positive attributes, the mechanical strength of commercial

purity titanium remains below the normal thresholds consid-

ered for hard tissue replacement. Indeed the desire for

enhanced strength has led to the increasing use of grade 4

commercial purity titanium for biodevices, the strength

increase above that of grade 2 being achieved through an

increase in the oxygen content.

Table 1

Microhardness, tensile mechanical properties and fatigue limit of grade 2 Cp Ti in different states [6]

State (structure type) Hv, (MPa) UTS, (MPa) YS (MPa) El. (%) RA (%) Fatigue limit (MPa

Coarse-grained 1800 460 380 26 60 238T10

UFG #1 (Equiaxed, submicron-grained) 2700 710 625 14 60 403T8

UFG #2 (Fibrous, with high dislocation density) 2821 960 725 10 45 434T5UFG #3 (subgrained with internal cells) 2850 1100 915 9 40 500T8

Ti–6Al–4V ELI (annealed) – 965 875 10–15 25–47 515

An alternative and potentially more attractive method for

enhancing the mechanical performance of commercial purity

titanium has recently been reported by Valiev et al. [6]. These

investigators have investigated the strengthening of grade 2

commercial purity titanium utilizing equal channel angular

pressing (ECAP) in combination with other deformation

processes. Procedures examined include ECAP (8 passes) at

400 -C (#1), ECAP+65% cold rolling (#2), and ECAP+rolling

followed by annealing at 300 -C, for 1 h (#3). Ultra-fine

grained (UFG) structures, Fig. 1, can range from an equiaxed

cellular microstructure to a sub-grain structure with a defined

boundary structure. In all cases the microhardness of severely

deformed commercial purity titanium was superior to that of

the original coarse-grained commercial purity titanium, Table

1. The yield and ultimate tensile strengths also exhibit this

enhancement, a 140% increase in ultimate tensile strength vis a

vis coarse-grained commercial purity titanium being observed.

Notably this increase was achieved while maintaining an

elongation to failure of 9%.

Table 1 also shows that the fatigue limit of ultra-fine grained

commercial purity titanium depends strongly on its micro-

structure state. For example, ultra-fine grained grade 2

commercial purity titanium processed via path 3 had a fatigue

limit of 500 MPa, almost 100% higher than its coarse-grained

commercial purity titanium counterpart. Additionally, ultra-fine

grained titanium exhibited a higher fatigue strength than

coarse-grained titanium in both the low and high cyclic fatigue

range, Fig. 2. A comparison between ultra-fine grained grade 2

commercial purity titanium processed via path 3 indicates that

its strength, ductility and fatigue limit are comparable to Ti–

)

0

0.5

1

1.5

2

Grain size

Cel

l adh

esio

n/ar

bitr

ary

unit

sCG UFG

Fig. 3. Enhanced osteoblast adhesion on ultra-fine grained (UFG) compared to

conventional grained (CG) commercial purity titanium, cell adhesion results

being normalized to adhesion on wrought Ti foil. Data are mean+std; n =3;

*p <0.1 compared to the titanium foil [9].

(a)

50μm

(b)

(c)

25μm

50μm

Fig. 4. Different microstructures that can be produced in Ti–6Al–4V; (a)

lamellar, (b) equiaxed and (c) bimodal [12].

H.J. Rack, J.I. Qazi / Materials Science and Engineering C 26 (2006) 1269–1277 1271

6Al–4V ELI. This suggests that substitution of ultra-fine

grained grade 2 commercial purity should be possible, thereby

eliminating any possible concern of the vanadium containing

Ti–6Al–4V [7,8]. Finally further enhancements in mechanical

performance should also be possible with other grades of

commercial purity titanium, this evaluation being presently

underway.

Recently Yao et al. [9] have shown that the aforementioned

ultra-fine grained structure produced by treatment #2 addition-

ally, influences the in vitro functionality of cells on a

commercial purity titanium implant surface. These results,

Fig. 3, provide evidence of increased osteoblast adhesion after

a 4 h exposure on ultra-fine grain commercial purity titanium

when compared to conventional grade 2 titanium. These

authors suggested that these results may be related to the

increased number of grain boundary sites in ultra-fine grained

materials or to an increase in the reactive site activity thereby

promoting cell adhesion. While the underlying mechanisms of

increased cell adhesion on ultra-fine grained titanium are still

under investigation, earlier studies of nano-grained metals

prepared by powder-metallurgy techniques do emphasis the

important role of grain boundaries, enhanced osteoblast

adhesion having been observed at grain boundaries. Further

efforts focused on in vivo investigations and including a study

of subsequent deposition of calcium-containing minerals, as

found in bone, are also currently underway.

Finally, some evidence exists that the frictional behavior of

ultra-fine grained commercial purity titanium is superior to

coarse-grained material. However it is not clear that the long

time response of tribo couples, where an effective ultra-fine

grain size is developed within the near surface regions during

wear, even in normal grain commercial purity titanium, will

lead to enhancement. This dependence may of course be

influenced by the specific tribo couple examined and the

particular methodology used to achieve the ultra-fine grain

microstructure.

3. Alpha–beta titanium alloys

The mechanical behavior of biomedical grade commercial

purity titanium is generally considered to lie below that desired

for total joint replacement. This led to the early introduction of

annealed Ti–6Al–4V, which today remains the largest single

titanium alloy used for biomedical device manufacture.

Continued concern with respect to the biological response of

vanadium containing materials has moreover led to the

development and introduction of Ti–6Al–7Nb [10], the level

of niobium substitution for vanadium being specified so that

the proportion of the alpha and beta phases during routine

processing mimics that of Ti–6Al–4V. Recent interest in

reduced modulus alpha-beta titanium alloys has resulted in the

development of Ti–13Nb–13Zr [11], its strength properties

being comparable to Ti–6Al–4V.

Present application of these materials tends to be limited to

the solution annealed condition, only slight attention being

given to enhancing the properties of Ti–6Al–4V by control of

the alpha/beta volume fraction and morphology. Fig. 4

illustrates the three distinct microstructures, lamellar, equiaxed

Table 2

Tensile properties the Ti–6Al–4V [13]

Microstructure YS

(MPa)

UTS

(MPa)

El.

(%)

RA

(%)

KIC

MPa/�m

Equiaxed (Std) 951 1020 15 35 61

Lamellar (Std) 884 949 13 23 78

Equiaxed (ELI) 830 903 17 44 91

Equiaxed (CMG) 1068 1096 15 40 54

Oxygen content: Std: 0.15–0.2%; Eli: 0.13 Max; Cmg: 0.18–0.2%.

YS: Yield Strength; UTS: Ultimate Tensile Strength; El.: Elongation; RA:

Reduction in area.

Table 3

Tensile properties the Ti–6Al–4V ELI [6]

State UTS (MPa) YS (MPa) El. (%)

1 Annealed 970 900 20

2 1+ECAP at 700 -C, E =6.5. 1160 1110 12

3 2+upset at 600 -C, E =55% 1450 1420 11

0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

1 10 100 1000 10000

Number of cycles

Dyn

amic

Fri

ctio

n co

eff UFG Ti 6Al-4V

Annealed Ti-6Al-4V

(a)


and bimodal, that can be produced in Ti–6Al–4V through

control of solution annealing temperature, cooling rate and

final aging temperature. The lamellar structure shown in Fig.

4a is typically produced following solution treatment above the

h transus, followed by air cooling, and aging between 700 and

800 -C. Solution annealing below the h transus, e.g., between

800 and 925 -C results in an equiaxed structure, Fig. 4b.

Finally, the bimodal structure, Fig. 4c, may be developed by

solution treatment below the h transus, typically between 900–

950 -C followed by air cooling and aging below 700 -C. Table2 shows that equiaxed alpha microstructures provide high

strength and ductility and relatively low fracture toughness,

whereas lamellar structure provides good fracture toughness

but with some compromise on strength and ductility. Finally,

the high cycle fatigue response of Ti–6Al–4V can be modified

through microstructure control. Fig. 5 shows that the bimodal

microstructure, previously discussed, has the highest high cycle

performance followed by the equiaxed structure, with the

lamellar microstructure having the lowest high cycle fatigue

resistance. Furthermore, within each of these microstructure

categories, finer microstructures result in higher high cycle

fatigue strength.

Preliminary efforts employing the severe plastic deforma-

tion procedures previously implemented in commercial purity

titanium have also shown promise for enhancing the mechan-

ical properties of alpha-beta titanium alloys, e.g., Ti–6Al–4V.

Table 3 shows that severe plastic deformation resulted in a

minimum 20% increase in yield and ultimate tensile strength

vis a vis annealed Ti–6Al–4V, with the tensile elongation

Fig. 5. Influence of microstructure on high cycle fatigue strength of Ti–6Al–

4V [14].

remaining above that typically required for biomedical

application (�10%). Additional enhancement can be achieved

by combining severe plastic deformation by equal channel

angular extrusion with upsetting.

Studies of the reciprocating sliding wear performance of

Ti–6Al–4V processed similarly to that described above again

suggests that this procedure may not offer marked improve-

ment. For example recent efforts show that the enhancement of

ultra-fine grained Ti–6Al–4V’s dynamic frictional coefficient

and the steady-state wear rate are marginal, with a slight

enhancement of the former at higher apparent contact stress,

Figs. 6 and 7. Certainly further effort examining and

understanding these phenomena are warranted.

Finally in vivo studies of cell functionality, in ultra-fine

grained Ti–6Al–4V have confirmed the enhancement in

osteoblast adhesion previously shown for ultra-fine grained

commercial purity grade 2 titanium [9]. The interpretation of

this data is however, further complicated when compared to

1 10 100 1000 10000

Number of cycles

0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

Dyn

amic

Fri

ctio

n co

eff

(b)

UFG Ti 6Al-4V Annealed Ti-6Al-4V

Fig. 6. 2-D friction traces (dynamic friction coefficient at maximum velocity

for Ti–6Al–4Vas a function of the number of dry reciprocating–sliding cycles

at an apparent contact stress of (a) 1.5 MPa and (b) 5 MPa [15].

)

0

0.2

0.4

0.6

0.8

1

1.2

1.5MPa 5MPa

Apparent Contact Stress

Stea

dy s

tate

Wea

r R

ate

(mm

/km

) UFG Ti-6Al-4V

Ti-6Al-4V

Fig. 7. Steady-state wear rate of Ti–6Al–4V during dry reciprocating–sliding

against a hardened steel counterpart [15].


the latter material since several levels of refinement are

possible in two phase Ti–6Al–4V, e.g., the alpha and beta

particle size, the uniformity of dispersion (primarily of the

beta phase) and the internal alpha and beta grain size. These

are not single valued functions of the processing history, Fig.

8, showing that low temperature annealing may have a

2.5

3.5

4.5

5.5

0 300 600 900

Temperature (ºC)

Temperature (ºC)

Nan

o ha

rdne

ss (

GP

a)

Alpha Beta

(a)

250

300

350

400

0 150 300 450 600 750 900

Mic

orha

rdne

s (V

HN

)

(b)

Fig. 8. Dependence of hardness on annealing temperature for nano-grained Ti–

6Al–4V; (a) shows nano-hardness of a and h phases and (b) shows overall

microhardness [16].

measureable effect on the nanohardness of the individual

alpha and beta phases without major affects on the micro-

hardness response.

These preliminary studies suggest that UFG structures have

superior mechanical and biological response in comparison to

their regular grained counterparts. Utilization of these proces-

sing routes to produce UFG structures for enhancement of the

mechanical and biological behavior of recently developed

titanium alloys containing only biocompatible alloying ele-

ments (e.g. Nb, Zr, Ta, etc.), remains an area for further

exploration.

4. Metastable beta titanium alloys

While commercial purity and alpha–beta titanium alloys

remain the primary titanium materials used for current

biomedical application the past decade has shown a substantial

increase in the synthesis of metastable beta titanium alloys

designed specifically for this field. Originally intended to

address the dual requirement of low modulus, approaching that

of bone, and enhanced biocompatibility, these systems are now

being considered for other applications (spinal, trauma, etc.)

which maintain the latter requirement while enhancing the

mechanical performance through artificial aging. Three alloys

were essentially developed simultaneously, Ti–29Nb–13Ta–

4.6Zr, Ti–12Mo–6Zr–2Fe (TMZF), Ti–35Nb–7Zr–5Ta

(TiOsteum) in Japan and the United States.

20

30

40

50

60

70

80

90

20 30 40 50 60 70 80 90 100

Grain Size (μm)

Grain Size (μm)

Elo

ngat

ion

& R

A (

%)

RA El

(a)

100

200

300

400

500

600

700

20 30 40 50 60 70 80 90 100

Stre

ngth

(M

Pa)

YS UTS

(b)

Fig. 9. Effect of grain size on (a) elongation and reduction in area, and (b) on

yield and ultimate tensile strengths of Ti–29Nb–13Ta–4.6Zr [17].

200

400

600

800

1000

1200

200 300 400 500 600 700

Aging Temperature (ºC)



Yie

ld S

tren

gth

(MP

a)

4 hrs

24 hrs

48 hrs

(a)

0

10

20

30

40

200 300 400 500 600 700

Elo

ngat

ion

(%)

4 hrs

24 hrs

48 hrs

(b)

60

65

70

75

80

85

90

95

100

105

200 300 400 500 600 700

You

ng's

mod

ulus

(G

Pa) 4 rs

24 hrs

48 hrs

(c)

Fig. 10. Effect of aging time and temperature on tensile (a) elongation, (b) yield

strength and (c) Young’s modulus of Ti–29Nb–13Ta–4.6Zr [18].

473

573

673

773

873

973

0 28.8 57.6 86.4 115.2 144 172.8 201.6

Time (ks)

Tem

pera

ture

(K

)

β

α+β

α+β+ω

β+ω

Fig. 11. Time temperature transformation diagram of Ti–29Nb–13Zr–4.6Ta.

The open symbols denote samples containing a// martensite after quenching

into ice water. The dotted, short dash and long dash lines denote 10, 30, and 50

vol.% of a phase, respectively [18]. T(-C)=T(K)�273.

Table 4

Tensile properties of Ti–35Nb–7Zr–5Ta [21]

Thermal treatment YS (MPa) UTS (MPa) El. (%) RA (%)

ST (0.06% O) 530 590 21 69

SA (0.06% O) 630 686 17 42

DA (0.06% O) 697 753 15 35

ST (0.46% O) 937 1014 19 55

SA (0.46%O) 1007 1055 12 27

DA (0.46% O) 1202 1244 8 16

ST (0.68% O) 1081 1097 28 50

SA (0.68% O) 1222 1252 9 13

DA (0.68% O) 1234 1260 7 9

SA: 260 -C 4 h, AC; DA: 260 -C 4 h, AC, 427 -C 8 h, AC; AC: Air cooled.


The former, Ti –29Nb–13Ta–4.6Zr, following water

quenching from the h-phase field displays a mixture of hphase and orthorhombic martensite (a//) and has an elastic

modulus of 65 GPa [17]. Further examination has shown that

the average h grain size and the volume fraction of martensite

has an important influence on this material’s mechanical

properties, Fig. 9. For example, solution treatment at 850 -Cfor 1 h, results in an average grain size of ¨50 Am with some hphase transforming to a//, the alloy having a yield strength of

¨250 MPa and elongation of ¨45%. Reducing the solution

treatment temperature to 750 -C and time to 0.5 h results in a

grain size reduction to 25 Am, reduces the volume fraction of

martensite considerably, increases the yield strength to 400

MPa and reduces the elongation to 30%, without any apparent

influence on elastic modulus [18].

The strength of this alloy can also be increased signifi-

cantly by aging, this increase in strength coming at the

expense of ductility and elastic modulus. Yield strengths as

high as 1100 MPa have been attained after long aging

treatments, e.g. aging at 450 -C for 48 h resulting in a yield

strength of ¨1150 MPa. However, this increase occurs at the

expense of elongation, which is reduced to less than 3%, Fig.

10, and an increase in elastic modulus to 85 GPa. Micro-

structure analysis indicated that the observed increase in the

strength after aging results from N and/or a phase precipita-

tion; aging at temperatures below 400 -C resulting in N phase

precipitation while aging at higher temperatures and/or longer

times, resulting in a mixture of N and a phases, Fig. 11.

Finally, aging at temperature above 475 -C results in only a

phase precipitation [18].

200

300

400

500

600

700

1 10 100 1000 10000

Cycles of Fatigue (x104)

Ben

ding

Str

ess

(MP

a) 0.06% O 0.46% O

Fig. 12. Stress-controlled fatigue response for Ti–35Nb–7Zr–5Ta with two

different oxygen contents [22].

(a)

(b)

4μm

4μm

0

500

1000

34 36 38 40 422θ

Inte

nsit

y (a

rbit

rary

uni

ts) β(110)

(a)

(b)

(c)

48 52 56

ααα(102 )

β(200)

ω(211 )

76 78 80 82 84

ω(212)

ω(301)

β(220)

(201)

(212)

Fig. 13. X-ray diffraction patterns of Ti–35Nb–7Zr–5Ta containing (a) 0.06,

(b) 0.46 and (c) 0.68 wt.% O aged at 427 -C for 8 h [21].


In contrast to Ti–29Nb–13Ta–4.6Zr, water quenching of

Ti–12Mo–6Zr–2Fe (TMZF) from the h phase field com-

pletely retains the h phase. In solution treated condition, TMZF

also has much higher tensile yield strength (1000 MPa),

comparable to mill annealed Ti–6Al–4V, along with 18%

elongation. In solution treated condition TMZF has an elastic

modulus of 79–84 GPa although bit lower than Ti–6Al–4V,

but much higher than Ti–29Nb–13Ta–4.6Zr. The tensile yield

strength of TMZF can also be increased by artificial aging with

a concurrent increase in the elastic modulus [19].

Finally, Ti–35Nb–7Zr–5Ta, at 0.06 O, has the lowest

elastic modulus, 55 GPa, of the more recently developed alloys

[20]. Table 4 illustrate the tensile properties achievable in this

alloy after different aging treatments for three different oxygen

contents. Yield strength can be increased at fixed oxygen

content, either by single (SA) or duplex (DA) aging, this

increase being accompanied by a slight decrease in tensile

elongation. An increase in the oxygen content from 0.06 to

0.46 wt.% O also increases the solution treated yield strength

from 530 MPa (0.06% O) to 937 MPa with a slight decrease in

elongation from 21% to 19%. It is notable however, that this

increase in the yield strength is accompanied by an increase in

the elastic modulus, 63 GPa [20]. Additionally, this increase in

the oxygen content also increases the high cycle fatigue

0

500

1000

34 36 38 40 422θ

Inte

nsit

y (a

rbit

rary

uni

ts)

β(110)(a)

(b)

(c)

48 52 56

(102 )

β

α

αα

(200)

ω(211 )

72 76 80 84

β(220)

(201)

(212)

ω(301)

ω(212)

Fig. 14. X-ray diffraction patterns of Ti–35Nb–7Zr–5Ta containing (a) 0.06,

(b) 0.46 and (c) 0.68 wt.% O aged at 260 -C 4 h/427 -C for 8 h (DA) [21].

strength of the alloy in the solution treated condition from

¨275 MPa to ¨450 MPa, Fig. 12 [22].

The oxygen content of Ti–35Nb–7Zr–5Ta strongly influ-

ences its aging behavior and hence its mechanical properties.

Solution treatment in h phase field followed by water

quenching results in complete h phase retention in alloys

containing up to 0.68 wt.% oxygen, some diffuse N being

present in low (0.06 wt.%) oxygen alloys [23]. Oxygen

addition restricts the motion of linear defects in metastable htitanium alloys thus hindering the collapse of alternating (111)

planes, the latter being required for N phase formation. Thus

4μm

(c)

Fig. 15. Scanning electron micrographs illustrating the microstructure of Ti–

35Nb–7Zr–5Ta containing (a) 0.06 (b) 0.46 and (c) 0.68 wt.% oxygen aged at

593 -C for 8 h [24].

(a)

(b)

20μm

20μm

(c)

20μm

Fig. 16. Scanning electron micrographs illustrating the tensile fracture surfaces

of Ti–35Nb–7Zr–5Ta containing (a) 0.06 (b) 0.46 and (c) 0.68 wt.% oxygen

aged at 593 -C for 8 h [24].


oxygen addition suppresses N phase formation. Being a strong

a stabilizer, it promotes a formation as well. Figs. 13 and 14

confirm that the increase in strength after both aging

treatments, at low oxygen content results from fine N phase

precipitation; duplex aging resulting in higher volume fraction

of N phase. In contrast, the increase in yield strength observed

at higher oxygen content (0.46 wt.% O) results from a mixture

of fine N and a phase precipitation [21,23]. Finally, the

increase in the yield strength observed at the highest oxygen

content (0.68 wt.%) results from a precipitation.

Formation of a phase in high oxygen alloys results from

local clustering of oxygen atoms which then acts as a

preferential nucleation site for a phase. Aging of 0.46/0.68

wt.% O alloys at 538 -C or higher results in oxygen diffusion

to grain boundaries from its surrounding areas, which in turn

leads to grain boundary (GB) a formation. This decreases the

oxygen content in the vicinity of grain boundary, hence

suppressing oxygen clustering and thereby resulting in a

denuded zones along the grain boundaries, Fig. 15. These

denuded zones act as a preferred path for crack propagation

resulting in premature failure, Fig. 16.

5. Conclusions and summary

This paper provided a snap-shot of several areas of current

exploration focusing on the synthesis and understanding

required for successful application of titanium alloys for

biomedical applications. Continued activity within this arena

will hopefully bring new materials and techniques to bear,

increasing the quality of patient care and lifestyle. Success in

this exciting endeavor will in the future require an ever

increasing cooperation of individuals with expertise in materi-

als science, biomechanics and cell biologists.

Acknowledgements

The authors would like to acknowledge the many former

and current co-workers, graduate and undergraduate students

for their contributions to our biomaterials activities. These

include Prof. T. Webster, Prof. R. Valiev, Dr. T. Ahmed, Dr. R.

Cooks, Dr. E. Fu, Dr. N. Istephaneous, Dr. M. Long, Dr. T.

Lowe, Dr. V. V. Stolyarov, Dr. V. Tsakiris, Mr. B. Marquadt,

Mr. C. Yao, Ms. M. Richards and Ms. H. C. Chandana. Of

particular note is Dr. J. Black who one day asked if it were

possible to synthesize a titanium alloy whose elastic modulus

approached that of bone. Without much thought the response

was ‘‘what is the elastic modulus of bone’’, thus beginning a

dialogue between materials scientists, biomechanics and cell

biologists that continues to this day. The financial support

during the early stages of this study by the Stryker Corporation,

under the guidance of Mr. Paul Serekian and Allegheny

Teledyne — ALLVAC, under the guidance of Mr. Howard

Freese and Dr. R. Kennedy are greatly appreciated. Finally the

continuing discussions with Prof. M. Niinomi have ‘‘kept’’ us

on our toes.

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Titanium Alloys for Biomedical Applications