ISIJ International. Vol. 31 (1991 ), No. 10, pp. 1139-1 146

Review

Creepof ce2+ Titanium Aluminide Alloys

Anthony W.THOMPSONand Tresa M. POLLOCKDepartmentof Metailurgical Engineering and Materials Science. Carnegie Mellon University. Pittsburgh. PA1521 3, U. S. A.

(Received on March 1. 199l, accepted in final form on May24. 199l)

The interest in titanium aluminide alloys includes elevated temperature applications, for which creepresistanc~ is a primary property. Tests have been madebetween 650 and 870'C on a variety of mi-crostructures of Ti-24Al-1 1Nb and Ti-25Al-10Nb-3M0-1V (ato/o) alloys. It has been found that micro-structure plays an important role in creep of these materials, so that thermal and mechanical history is

important. Stress exponents for power-law creep, and apparent creep activation energies, have beendetermined for these alloys. As is usually found in structural alloys microstructural characteristics whichincrease ductility and toughness at low temperature tend to accelerate creep considerably, particularly thepresence of pphase, and most notably whenarranged as locally-continuous Pfilms between plates ofthe ce. phase. Solution treatment in the pphase provided optimumcreep resistance, but cooling rate effects

were different in the two alloys considered. Comparison to near-c( titanium alloys developed for creepresistance, showsthat the aluminide alloys have better performance, especially above 700'C.

KEYWORDS:titanium aluminides; creep; microstructure; temperature dependence; stress dependence.

l. Introduction

There has been considerable recent interest in thedevelopment of advanced titanium aluminide alloysbasedon Ti3Al for aircraft engine componentsand otherhigh temperature applications. It has been found thatTi3A1 retains adequate creep strength to muchhighertemperatures than do conventional titanium alloys,

but it exhibits low ductility at room temperature.1'2)Therefore, development of aluminide alloys has beendirected to combinations of room temperature ductility

and elevated temperature creep strength. Onealloyingapproach has been the addition of p-phase stabilizing

elements such as niobium to Ti3Al. A110ying of this kindhas proven successful in providing some propertyimpro,vements, in alloys such as Ti-24A1-1 INb (values

are ato/o or a/o), and more recently, the alloy Ti-25A1-10Nb-3V-lMowas developed ~vith further ad-ditions of pphase stabilizers.1~5) These alloys are re-ferred to below as Ti-21~ll and Ti-25-10-3-1 re-spectively.

For conventional titanium alloys, most studies oncreep behavior in recent years have been performed onnear-c( alloys, such as Ti-6242S6) and IMI 685,7) andhaveshownthat the creep resistance of such alloys is stronglydependenton microstructure. It wasfound, in these andother studies,8.9) that Widmanst~tten microstructures,obtained by solution treating above the p transus,produced creep resistance at high temperatures superiorto that provided by equiaxed microstructures obtainedby processing in the o( +Pphase field. It wasalso foundthat creep resistance is maximizedat intermediate coolingrates from the psolution treatment temperatures.

1139

For alloys basedonTi3A1, on the other hand, relatively

few data on creep behavior have been published. Thepioneering work of Mendiratta and Lipsittlo) was per-formed to investigate the steady-state creep behaviorof Ti3A1 and Ti3Al + 10 w/o Nb(about 5a/o Nb) in the

temperature range of 550 to 825'C. The Ti3A1 phase,called oc2, mayin alloys be accompaniedby the bcc pphase. In that study,10) it wasshownthat at temperaturesabove 700'C, the stress exponent n of the power-lawcreep equation indicated a transition in mechanism,andin the high stress and temperature regime addition of Nbincreased the apparent activation energy for creep de-formation. Subsequent work,11•12) together with con-siderable information in the form of contract researchreports, has been reviewed.13 - 15)

In two recent investigations,16,17) room temperaturetensile and elevated temperature creep studies have beenperformed on various microstructures of Ti-24-1 1andTi-25-10-3-1

.These studies evidenced particular inter-

est in the creep behavior of these alloys at the tempera-tures between 650 and 870'C, since the use of conven-tional titanium alloys has generally been restricted5) to

temperatures below 600'C. In the present overview, the

primary focus is the influence of microstructure, since it

is evidentl4) that this factor is very important in creepof these alloys; identification of stress and temperatureeffects on creep behavior is also included.

2. Experiments

Composition and processing of the Ti21~11 andTi-25-l0-3-1 alloy materials has been presented previ-ously.16,17) A few tests were performed on as-received

C 1991 ISIJ

ISIJ Internationa!, Vol. 31

material, which had been processed in the cc2 +pphasefield. Preparation of other microstructures was accom-plished by resolution treating in either the o(2+p orthe pphase field and controliing the cooling rate fromthe solution treatment ternperature. Transformed pmi-

crostructures were produced by a I h/P solution treat-

ment followed by one of the following cooling schemes:

slow or furnace cooling, SC; controlled cooling, CC;air cooling. AC. The SC, CCand ACtreatments pro-ducedaverage cooling rates of O, I ,

0.5, and 10 K/sec, re-

spectively, whenmeasuredbetween the solution treat-

ment temperature and 760 'C. The oc2+P microstruc-

tures were produced by solution treating within the oc2 +pphase field an(} the samecooling schemesas above.(The ptransus temperature for Ti-2d~l I is about 1125'Cand for Ti25-l0-3-1 is about I075'C.) For the Ti2d~1l, two solution treating times were used in the oe2+pheat treatment to vary the size of o(2 plates.

Conventional creep testing of Ti-21~11 and Ti-25-10-3-1 alloys was conducted, using constant-loadtechniques in air. Specirnens had a 6.4 mmdiameter and

a 37mmeffective gage length, with sufficient metalremoval during machining to ensure that contaminatedsurfaces were removed. Test procedures followed ASTME139-79. Creep strain was measured using a SLVCtransducer which allowed a strain resolution of 5x IO6.

For tests which were continued to fracture, the fracture

surfaces were examined in a scanning electron micro-

scope; thin foils were examinedby transmission electron

microscopy (TEM) for various creep strains.

(1991 ), No. 10

3. Results and Discussion

3.1. Microstructures

The microstructures obtained from the various heat

treatments were as follows. Solutionizing in the pphasefield and air cooling produced fine acicular oe2 plates

arranged in a Widmanstatten basketweavemorphology,while slower cooling rates resulted in coarser oe2, Iess

retained pphase, and a transition from basketweave to

aligned colony o(2 structures. Solutionizing in the oc2•h pphase field produced elongated primary c(2 plates in aWidmanstatten basketweave morphology, surroundedby a "transformed p" matrix similar to that of psolution

treated microstructures. This transformed pconstituent

comprised secondary acicular c(2 plates as well as someretained pphase. Whennucleation of individual plates

becomesslow "sympathetic" nucleationl8) gives rise to

parallel o(2 plates of commonorientation, which are called

colonies.

Examplesof these microstructural types are containedin Fig. 1. The relatively fine Widmanstatten oc2 plates

formed on cooling from the pphase are shownin Fig.

l(a), indicating that the transformation is dominatedby easy nucleation. Cooling from the oc2 +ptwo-phaseregion typically gives coarser o(2 plates, although still of

Widmanstatten character, indicating that growth of the

plates is moresignificant, Fig. 1(b); note also in this figure

that sorne primary oc2 has formed at prior pgrain

boundaries. Figure l(c) shows that with moreworkingin the c(2 +ptwo-phase region, Iarger and moreglobular

primary cc2 forms at grain boundaries,18) while oc2

(d)

Fig. l.

Examples of microstructural range observed in e(2 +ptitanium aluminide alloys. (a) Widmanstatten orbasketweave morphology of c(2 plates, formed in

Ti-25l(~3-1 on cooling from the pphase at 0.5 K/sec.

(b) Coarser Widmanstzitten c(2 plates, formed in

Ti-25-l0-3-1 on cooling from the cc2 +p two-phaseregion at 0.5 K/sec. (c) Primary o(2 and "packet" e(2 plates

in Ti-25-lO31 cooled from the ec2 +p two-phase re-

gion. (d) Colonies of ce2 plates in Ti21~l I cooled from

1OOO'C(low in the ec2 +ptwo-phase region) at 5K/sec.

C 1991 ISIJ 1140

ISIJ International, Vol. 31 (1991), No. 10

plates showsomegrouping into " packets" of parallel

plates in most prior pgrains. This is a commonas-hot-worked microstructure. Finally, colonies contain-ing a numberof oc2 plates with a commonorientation

can be formed by cooling from low in the oe2+ptwo-phase region, as in Fig. l(d). More detailed light

and electron micrographs of these structures have beenpresented elsewhere. 16 ~ 18)

It should be mentioned that the retained pphase in

Ti3A1 alloys maybe ordered or disordered. The details

of retained pphase were not investigated via TEMin

the studies discussed here, but it is believed from earlierreportsi9 ~ 23) that retained pphase in Ti-25-l0-3-1 has

a fully transforme~drordered structure (B2. CsCl type). It

wasshownin those phpers that addition of pstabilizers

such as Mo, Nb, or V tended to result in the orderingof the pphase to a B2structure. In the work of Strychoret al.,22,23) the nature of pphase ordering even in

water-quenched specimens of Ti-27.8A1-11.7Nb (a/o)

was confirmed by the observation of superlattice re-flections characteristic of a CsCl B2 structure. This or-dering has been observed in Ti3A1+Nb alloys whoseNb contents lie between about 15 and 35 w/o, eventhough the extent of this field at the high Nbend is still

undetermined in detail.22,23) On the other hand, TEMevidence to date in Ti-21~ll has typically shown thatthe pphase mayor maynot be ordered. Evidently the

Ti24-1 1composition lies near the threshold composi-tion23) for ordering to take place under the heat treatingconditions typically studied.

3.2. Creep Properties

A principal goal of the two studies emphasizedherel6,17) was identification of microstructure effects in

creep. The role of microstructure can be illustrated bystrain-time curves for Ti-25-lC~3-1

, as shownin Fig. 2.

Here, two temperatures and stress levels both showthat

creep is dramatically affected by the microstructure ofthe material. Most of the creep curves have a classical

appearancewith a primary or transient period followedby a steady state regime. It is useful in Fig. 2to interpret

creep behavior in terms of two processing parameters:(1) solution treatment temperature (p/SC vs. oe2 +fi/SC),

and (2) cooling rate from the pregion (p/AC. CC, andSC). At both temperatures, the oc2 +pprocessed micro-structure (c(2+pISC) which had elongated primary ae2

grains shows lower creep resistance, whencomparedto the psolution treated microstructure (p/SC), althoughthe magnitude of the effect of solution treating processwasmuchgreater at 760'C, Fig. 2(b).

Creep resistance, however, cannot be characterized bythe fi solution treatment alone. Theeffect of cooling ratefrom the pphase region on microstructure and its rolein creep performance needs to be taken into account.Creep resistance increased with decreasing cooling rateat 760'C, Fig. 2(b), and thus creep resistance of theslow-cooled colony microstructure (O.1K/sec) wassignificantly improved. However, at 650'C, as shownin

Fig. 2(a), creep resistance increases with increasingcooling rate through a maximumat an intermediatecooling rate of 0.5 K/sec which produced a basketweave

5(a)

~4Z

3~(,)~UJ 2UJa:O

1

o

ce2+~!SC

pIAC

6500c, 414MPa

P/SC As•RecGlved

p/CC

o

10

10 20

TIME(HRS.)

(b)

30 40

8

Z:~ 6u)I~LU 4LL,

a:

O2

o

760'C, 207MPapl AC

oc2+~/SCAS-RECEIVEO

INE WIDMAHSTA~EN

pl CCBASXETWEAVE

pl SCCOLONY

1o 20 3011ME(HRS.)

Fig. 2. Creep strain as a function of time, for Ti-25-l0-3l.Each curve is labeled with the microstructure as(solution treatment)/(cooling rate), (a) Test at 650'C,414MPa(60ksi) stress. (b) Tests at 760'C, 207MPa(30 ksi) stress. Microstructural descriptions added to

somecurves. FromRef. 16).

microstructure. The effect of cooling rate is discussedfurther below.

3.3. Stress DependenceThe creep tests on both alloys were analyzed to de-

termine the stress and temperature dependenceof creepdeformation in the temperature range 650 to 870'C.Measurementsof apparent stress exponent n were madeby decreasing the load during testing in order to provide

a step change in stress and thus strain rate.24) At each

newstress level, the test continued until a newsteady-state creep rate could be determined, thus minimizingstructural changes for each successive decrease in load.

Most tests were terminated during steady-state creep.For materials which undergopower-law creep, accordingto the usual equation,

~* =A(T"e ~ Q'/R T (I)where i* is the steady-state creep rate at a tensile stress

cr, n and A are constants, Q, is the apparent activation

energy. R the gas constant, and T the absolute tem-perature. Therefore, plotting log of steady-state creeprate vs. Iog a will result in a straight line with a slopeequal to n, the power-1awexponent.25 ~27)

The results for the role of microstructure for bothTi-24-11 and Ti-25-10-3-1 can be illustrated with the

two examples in Fig. 3. In each case, a single test

temperature is used to comparevarious starting micro-structures, with specimen codes identified in the figure

caption. Similarly, the effects of temperature on any onemicrostructure can be compared as in the examples

1141 C 1991 ISIJ

ISIJ International. Vol, 31 (1991), No. 10

oco,,)Q)

101

100

10~ 1

10~

lcr 3

10~ 4

10~

10~

La)

L~~lE:i200~1ieAic'.hlil

e loeoAcFP+ TeAl (hl&L)

3J1

377

434

T=650'C

1 10 1oo 1ooO

~eZ:

~(1)Cl)

10'2

stress (MPa)

Lb)

6~O~ocCREEPE9 AS'RECEiVED,n=4.6

. plsc, n=5.8

a ,x2+0/sc, n=4.9e p,cc, n~4.2

lcr 1

10~

(a)

lcr3,

10~

lcr5

lcr

lcr7.

1cr8..

::::

,t,::

u

h,,,>,:]uJ

e,,

10'3 n:4.2

x 870'c

n~.O

10'4

10~5

1O'6

n=5

Fig. 3.

/2'1

4'4

// 4~3

/

1200AC

10~2

10o 1O1

Stress (MPa)

(b)

102 103

1oo IoooSTPEss (tAPa)

Stress dependenceof creep in two aluminide alloys. (a)

Creep rate vs, stress for Ti21~ll at 650'C; "I200AC" refers to pIAC treatment, "IOOOAC" to

c(2 +pIAC. Also shown is result for stoichiometric

Ti3Al from Ref, lO). Data from Ref, 17). (b) Creeprate vs. stress in Ti25lO3-1 for several micro-

structures (indiv{dual stress exponents, n, shown).

FromRef, 16).

'( I0'3

EC

Oh 4IO~

,FO~hU)

I0~5

FINEW~DMANSTAITEN

n=3 5

8700c

n=2.8

76'oocl40c

n=4'o

////n

::ooc

n=4 7

of Fig. 4. It is evident in Figs. 3and 4that two patterns

of stress dependencewere observed, either a constantvalue of n or a varying n value. In the latter cases, the

data appeared to be adequately described by a two-partlinear fit, i,e. two n values.

Increasing creep test temperature generally led to

slightly decreasedvalues of n. At lower temperatures andlower stresses, the abrupt change in stress exponent in

some cases may be an indication of a transition in

rate-controlling mechanism.Higher values of n, in the

range of 4 to 5for single-phase materials and up to as

muchas 11in two-phase or multi-phase materials, areusually accepted as indicating diffusion-controlled dis-

location creep processes.28~30). On the other hand, avalue of n near or equal to unity indicates diffusional

creep by a stress-directed flow of atoms with grain

boundaries acting as sources and sinks, with this processoccurring at low to intermediate stresses, cr=(10~5 to

l0~4)E, whereEis the Young's modulus.30,31) The lowstress slopes of the present aluminide materials, Figs. 3and 4, approximate the slope of the curves at high

temperatures, where diffusion processes are dominant.Elevation of the low-stress creep rates (relative to anextrapolation of behavior at high stress) mayoccur dueto somecontribution of grain (or phase) boundarysliding

or diffusional creep to the overall creep rate.31~34) Avalue of n=2has been taken to indicate a contribution

to the creep rate from grain (or phase) boundary sliding

mechanisms.

32)

The gradual reduction in stress exponent with

C 1991 ISIJ

10

Fig. 4.

1oo Iooo

STREss (h4P8)

Temperature effect on stress dependenceof creep in

the samealloys as Fig. 3, shownin form of creep rate

vs. stress curves. In each case, note change of slopein some tests. (a) Example for Ti-21~ll, the pIACmicrostructure, from Ref. 17). (b) Example for

Ti-25-10-3-1, for fine Widmanstzitten structure, Ref.16).

increasing temperature, Fig. 4, mayalso have occurredif boundary sliding were to emerge as an important

process and thus the exponent, n, would decrease, since

grain boundary sliding tends to decrease n, relative to amaterial without grain boundary sliding.32) Becauseboundary sliding is not an independent mechanismproducing a steady-state creep strain and must beintegrally connected with intra-crystalline creep deforma-tion,25,30,33) it is possible that two mechanismsareoperative at the highest test temperature used, 870'C.

3.4. Temperature DependenceIn addition to the stress exponent n, the apparent ac-

tivation energy Q, is considered helpful29,35) in estab-

lishing the controlling mechanismof creep deformation.

In Fig. 5, plots of steady-state creep rate vs. temper-ature are shown. Again, assuming power-1aw creep,

Eq. (1), whenthe log of creep rate is plotted against llT,

a straight line results, with a slope equal to Q.. Therehavebeenseveral determinations of Q, for Ti-21~1 1.

Theactivation energy in the low stress regime for Ti-21~l 1with a 100o/. transformed pmicrostructure was found

to be 120kJ36) under conditions of n= I ,while an c(2 +p

microstructure exhibited a value of 259kJ/mol,12) jn the

stress range of 103 to 138MPa. In the work described

here, 17) Q. values ranged from I IOkJ/mol for as-received

material, to the 142kJ/mol for the c(2 +pmicrostructure,

1142

~e~(,(lDC')

ISIJ International, Vol. 31

100 Ca)

100MPa

lcr 1Q=106*kJ

lcr1~

10r3 135 kJ

10*

c AR134 kJ

10~5,El lOOOAC-15+ 1200AC

10*6,

o. o008 o. o009 0.001 o o. Oll/T (1/K

1o'2

l/T (1/K)

(b~

ool 1

a: A~~RECEIVED207 MPa:::

Hcc

l0~3

~ Q=305kJlmola:otL'

h I04hu)

>o~h"' Io~;

9eoe-4 9.80o-4 1,00e'3 1,02e'3 104e-3 1oee-3 1.08e 3lrT (K)

Fig. 5. Creep activation energy determinations, from creeprate vs, l/T plots. (a) Data for Ti2d~1 1at stress of100MPa(average slope values in text are averages ofdeterminations at several stresses), from Ref. 17). (b)

Data for fine Widmanstzitten microstructure ofTi25-lO3-1, at 207MPa,from Ref, 16).

and 134kJlmol for the Psolution treated structure. All

values were measured in the low stress regime (belowlOOMPa), and are in fairly good agreement with theactivation energy for diffusional creep (n=1) of121 kJ/mol in fine-grained c( titanium at stresses below2MPa.37) This latter value was attributed to grain

boundary diffusion. Values for self-diffusion, pipe dif-

fusion, or graln boundary diffusion in c(2 alloys are notavailable, either for the oc2 phase or for the Pphase.

Thepphase data maybe of particular interest, since, asdiscussed below, the plasticity of the pphase plays animportant role in deformation of o(2 aluminide alloys at

both low and high temperatures.

A creep activation energy of 206kJ/mol for stoichio-

metric Ti3Al at all stresses and temperatures has beenmeasured,10) while addition of lOw/o Nb to thestoichiometric o(2 increased Q, to 285kJ/mol.io) TheTi25-l0-3-1 alloy, with considerably morepstabilizer

content, was found to have a Q, of 305kJ/moll6) in theclimb-controlled creep regime. This value of Q., althoughlarge compared to the roughly 125kJlmol values for

Ti-21~ll, mayreflect a different activation energy for

diffusion, or a changing rate-controlling diffusion path

or process. It has been concluded that the apparentactivation energies for high temperature creep in metalsshould be independent of creep stress and strain,38) butdifferent mechanismsseemto be dominant in different

stress regimes. It has been suggested that the activation

(1991 ), No. 10

energy for both crystallographic slip and boundarysliding are the sameand that s]ip maycontrol sliding,39)

since, as mentioned, sliding is not an independent processin creep.25,30,33) More work appears to be needed toclarify the source of high Q. values in c(2 alloys, includingthe role of pstabilizing elements in both diffusivity valuesand in ordering of the pphase, microstructure, anddiffuslon paths.10,16,17) Ordered or disordered pphase,varying interphase boundary composition, and relativeplasticity of the oc2 and pphases, could all be relevantto establishing the creep mechanism(s) in these alloys.

3.5. Microstructural Effects

The role of coollng rate in Fig. 2can apparently beinterpreted in the following way. A phenomenologicalmodel6.40) was recently proposed to interpret creep be-havior of near-oc titanium alloys in terms of micro-structural factors. Observation of an optimumcoolingrate for maximumcreep resistance was explained inthis model in terms of the trade-off between two factors:

an increasing amountof retained pphase (detrimentalin creep resistance) and decreasing c( plate size (beneficialin creep resistance). The present observations for Ti-25-lO3-1 on the interaction between microstructureandcreep deformation can be rationalized using the sametwo factors for near-oc titanium alloys.

Oneimportant factor is that the effective slip lengthin elevated temperature creep deformation is evidentlycontrolled by primary oe2 size in o(2 +pprocessed micro-structures and by Widmanstzitten oc2 in pprocessed mi-crostructures. Even though the Burgers relationship is

expected to exist between o(2 and ordered pphase, as it

does betweeno( andpphase,4i'42) it is not surprising thato(2/B2 boundaries are effective barriers to slip and act toconfine active dislocations to individual oc2 platelets,

becauseof the relatively low applied stresses necessitatedby the high creep temperatures. Therefore, inferior creepresistance of the c(2 +fi/SC microstructure is attributedto a larger slip length than that of the p/SC micro-structure.

However, oc2 plate size alone is not the sole controllingfactor in creep deformation. If it were, it would beexpected that the microstructure with fine Widmanstatteno(2 platelets would exhibit superior creep resistance. Themaximumin creep resistances that occursl6) in the slowcooled microstructure (p/SC) at 760 and 650'C and lowstresses suggests that another microstructural factor,

competingwith oe2 plate dimension, is present. That factoris the amount of retained ordered pphase (B2)' Fastcooling increases the amountof B2 phase,22,23) believedto be present in Ti25-l0-3-1, thus resulting in poorcreep resistance, becauseof its ductility and intrinsic highdiffusion coefficient. It is, therefore, required to haveoptimumc(2 plate size and optimumamountof retainedB2 phase for superior creep resistance at elevated

temperatures. If the pphase is disordered, these factors

maybe different. At 650'C and high stresses, however,

a somewhatcomplicated role of microstructure wasobserved.16) The low creep resistance in p/SC micro-structure in this regime suggests that a different structuralunit beyondc(2 plate size waseffective, similar to that of

1143 C 1991 ISIJ

ISIJ International, Vol. 31 (1991), No, 10

room temperature tensile behavior,5,13) or that the

thermally activated deformation process in retained B2phase wasnot significant in controlling creep deforma-tion at this temperature. Thuscreep at 650'C. Fig. 2(a),

can be regarded as a transition betweenpatterns of lower

temperature deformation and those of high temperature

creep deformation.These interpretations are supported by dislocation

structure and fractographic results.16) The predominantdislocation type had a Burgers vectors, i.e. ,predominantly on prism planes, at both room and el-

evated temperatures, as has also been shownby several

other investigators. 5, 13.43 ,44) Extensive cross-sli pof these

dislocations to.basal and pyramidal planes at the higher

temperatures was observed,16) again in agreement withprior work. Other dislocations, evidently c+a type,

were also observed.13,16.44) Fractographic examinationshowedclearly that the retained por transformed pconstituent acted in a relatively ductile manner, com-pared to the cleavage-1ike fracture of the oe2 phase,

both at room and elevated temperature.16,18,45) Thisresult is also in accord with deformation studies.44) At750'C and above, someintergranular fracture wasalso

noted, in accord with prior work.43)

Turning from the Ti-25-l0-3-1 results of Fig. 2to the pattern of results for Ti-21~l l, similar conceptsshould be applicable. Recrystallization and grain growthduring creep, oc plate width, pvolume fraction, and ccloe

vs. p/p boundaries are all factors that are thought to

influence steady-state creep rate in conventional titaniumalloys.6) If only cc2 plate size were considered to control

steady-state creep rate in titanium aluminides, a conceptthat wasshownabove to be inadequate to describe creepin Ti-25-10-3l, then we would expect the micro-

structure with the smallest plates, that is, the p/ACandoc2+p/AC-15 structures ("I5" designates the solution

treatment time in minutes), to have the best creepresistance, due to their short slip length. It is seen,however, that c(2 +p/AC-30 (30 min solution treatment),

with coarser oc2 plates, has superior creep resistance upto 870'C. Therefore, slip distance and correspondingdislocation motion within plates is not the sole con-trolling factor in creep of Ti-21~1 1in the temperatureand stress range studied. Rather, since the oe2 +p/AC-30plates are coarser than the plates present in the other

microstructures, then it is true that there is less interphase

boundary per unit area in this stucture. The ratio egb/et'

the strain due to grain boundarysliding to the total creepstrain, was found to increase with increasing tempera-tures39) or decreasing stresses.46) Also, the oe2+fi/ACstructures should contain smaller amountsof pphase, afactor which is comparable to that invoked in discussing

Ti-25-10-3-1 above. This phase has a higher diffusion

rate and thus, inferior creep resistance. This wouldexplain the fact that oe2+p/ACstructures creep moreslowly than the p/AC structure. Examination of micro-

structures in crept specimensl7) demonstrated that re-

crystallization and grain growth did not occur appre-ciably during creep. The presence of crystallographic

texture, e.g. an alignment in basal planes in processing,

could help explain the high steady-state creep rate of as-

C 1991 ISIJ

received material.

3.6. Comparisonto Other Alloys

The creep behavior of several o(2-based alloys is

comparedat 650'C in Fig. 6. Three results for Ti-24-1 1are included,12,i7,36) as well as a stoichiometric Ti3A1resultlo) and one set of data for Ti-25l0-3-1 .1 6) A Iine

for the p/ACmicrostructure in Ti-24l I is included, as

a fairly typical result amongthe data shown. It is evident

that the Ti3Al and Ti25-l0-3l have somewhatbetter

creep resistance (about one order of magnitude lower

creep rate) than the Ti-21~l I data shown.In Fig. 7, creep behavior of Ti-25-10-3-1 is compared

at 650 and 760'C with those of two near-c( titaniumalloys, Ti-6242S6) and Ti-1 100,16) the latter of which is

1144

~O~

~Cl)C,)

10~

lcr3,

lcr4,

1(T

lcr6,

lcr

Fig. 6.

i?;~

UJ~~,:CLUJuJC:OUJHC~>C:

UJ~CO

c:~:

uJh~(

,:

~gUJh~

~,,,>OUJ~u,

7

,,

+DXI

lOOOAC(Hayes)

BHT(B anerjee)

Ti3Al (M&L)25- l0-3 - I(Cho)

1200AC(AIbcrt)

+ +

n=4.4

,(

+

,(

,(

+

Dx

nD

+

X

X

T=650'C

1O1 102 103Stress (MPa)

Comparison of creep in several titanium aluminidealloys. Data for Ti-24-1 1of Hayes,12) Albertl7) andBanerjee3s); Ti-251(~3-1 from Ref, 16); and Ti3Al

from Ref. lO).

10~2

10~3

10~4

10~5

10~6

1012

1o' 3

10~4

Fig.

10~5

10~

7.

10

~,p

65O'CCREEP

Ti~242STi-1 1OOTi•2~1 O•~1

76rCCREEP

o.9

1.7

a)

n=6.3

1oo

STREss("p8)

Lb)

n=9.1

6.

4.8

I

t]

lp

l

5.7

Ti ~242STi*1 1OOTi-25-10-~1

1OOo

10 Ioo looo

STRESS(MPa)

Comparisonof creep in several titanium andaluminide

alloys. Data for Ti-6242S6) and Ti-1 10016) compared

to Ti25-l0-3-1 16) at two temperatures.

ISIJ International. Vol.

a newly-developed alloy which achieves improved creepresistance with 0.450/0 Si, known24) to provide creepresistance, and smaller amounts of pphase stabilizers.

For this comparison, the optimummicrostructure (colo-ny) was used for all alloys. Generally, creep resistanceof Ti25-l0-3l is improved by one order of magni-tude over Ti-1 100 and by two orders of magnitude overTi-6242S. Unlike near-o( alloys, a transition in mecha-nism in Ti-2510-3-1 was not found at 760'C in thesimilar stress range studied. At the lower temperature,650'C, deformation mechanismsappear to be similar tothose observed at lower temperature,5,13,44,47) as notedabove. As mentioned above, the possibly varying roles

of pphase in the~lluminide alloys would makeit mostvaluable to test alloys of varylng pstabilizer content,particularly if the degree of order of the pvaries.

4. Conclusions

(1 ) Microstructural effects on creep behavior of twotitanium aluminide alloys based on Ti3Al. Ti-21~lland Ti-25-l0-31, has been examined in somedetail

as a function of stress and temperature (650-870'C).(2) It is believed that 650'C creep showsa transition

deformation behavior between room temperature ande]evated temperature behavior. At temperatures above650'C, p solution treated microstructures exhibitedsuperior creep resistance comparedto oc2 +pprocessedmicrostructure (c(2+p/SC vs. pISC) unless a:2 plates

were coarsened by extending solution treatment in the

c(2 +pfield.

(3) A pronounced influence of cooling rate fromthe p solution treatment temperature of creep wasobserved in Ti-25l0-3-1. At 760'C, p/SC (colony)

wasproven to be an optimummicrostructure for creepresistance at all stresses studied. However, at 650'C,possibly representing transition deformation behav-ior, a somewhatcomplicated effect that depends onstress level was found. At this temperature, p/SC showssuperior creep performance only at relatively lowstresses. Thedegree of order, and the relative plasticity,

of the pphase maybe important in these effects, anddeserves further study.

(4) The stress exponent n tended to decrease withincreasing temperature. At higher temperatures, somemicrostructures also exhibited an apparent transitionin creep mechanism, depending on stress level. Theobtained n values are generally within the range ofvalues predicted by theoretical models of dislocation-

controlled creep. Acreep activation energy of 305kJlmol

was measured for Ti-25-lO3l and values of about125kJ/mol were obtained for Ti-24l I .

Reasonsfor thelarge difference in values are not known, and mayreflect

major differences in behavior of pphaseor of interphaseboundaries, or possibly a major of role of stress level.

(5) Creep resistance of Ti-25-l0-3-1 wascomparedto Ti21~l I and to two creep-resistant near-c( titaniumalloys, and was found to be improved by one or twoorders of magnitude over each, in the test regimesstudied, when the optimum microstructure (colony) is

used for all alloys. The Ti-25-lO-31 alloy is also

1145

31 (1991). No. 10

similar to or superior to stoichiometric Ti3Al in the

creep tests discussed here.

Acknowledgments

Weappreciate helpful discussions with J. C. Williamsand W. Cho, and provision of unpublished results byD. E. Albert. Support of the experimental work describ-ed here was provided by the U.S. Air Force Office ofScientific Research, both under contract F49620-87-C-O017, as part of the University Research Initiative pro-gramon High TemperatureMetal Matrix CompositesatCarnegie Mellon University, and under grant AFOSR90-0033.

REFERENCESl) H. A. Lipsitt: High-Temperature Ordered Intermetallic Alloys,

MRSSymp., Vol. 39, ed, by C. C. Koch, C. T. Liu and N. S.

Stoloff, MRS.Pittsburgh, (1985), 351.2) R. L. Fleischer, D. M. Dimiduk and H. A. Lipsitt: Ann. Rev.

Mate,'. Sci., 19 (1989), 231.3) M. J. Blackburn, D. L. Ruckle and C. E. Bevan: Research to

Conductan Analytical Investigation of Alloys, Technical ReportAFML-TR-78-18, Air Force Mater. Lab., Wright-PattersonAFB, Ohio, (1978).

4) M. J. Blackburn and M. P. Smith: R&Don Composition andProcessing of Titanium Aluminide Alloys for Turbine Engines,Technical Report. AFWAL-TR-82-4086.Air Force WrightAeronautical Lab., Wright-Patterson AFB. Ohio, (1982).

5) Y.-W. Kimand F. H. Froes: High Temperature Aluminides andIntermetallics, ed. by S. H. Whang,C. T. Liu, D. P. PopeandJ. O. Stiegler, TMS-AIME,Warrendale, (1990). 465.

6) W.Cho. J. W.Jones, J. E. Allison andW.T. Donlon: Proc. SixthInt. Conf. on Titanium, Vol. l, ed, by P. Lacombe,R. Tricot andG. Beranger, Les Editions de Physique. Paris, (1989), 187.

7) P A. Blenkinsop, D. F. Neal and R. E. Goosey: Titanium andTitanium Alloys, Proc. Third Int. Conf. on Titanium, Vol. 3, ed.

by J. C. Williams and A. F. Belov, Plenum, NewYork, (1982),

2003.8) C.C. ChenandJ.E. Coyne: Titanium '80Scienceand Technology,

Proc. Fourth Int. Conf. on Titanium. Vol. 2, ed by H. Kimuraand O. Izumi. TMS-AIME,Warrendale, (1980), 1197.

9) P. A. Blenkinsop: Titanium. Science andTechnology, Proc. Fifth

Int. Conf, on Titanium, ed by G. Ltitjering, U. Zwicker and W.Bunk. DGM.Oberursel, (1985) 2323.

lO) M. G. Mendiratta and H. A. Lipsitt; J, Mate,'. Sci., 15 (1980),

l l)

l2)

l3)

l4)

l5)

l6)

l7)

18)

2985.C. G. Rhodes: Proc Sixth Int. Conf, on Titanium, Vol, l, ed. byP. Lacombe,R. Tricot andG. Beranger. Les Editions de Physique,Paris, (1989), 199.

R. W. Hayes: Sc,'. Meta!l., 23 (1989), 1931.

D. A. Koss. D. Banerjee. D. A. Lukasakand A. K. Gogia: HighTemperatureAluminides and Intermetallics, ed. by S. H. Whang,C. T. Liu, D P. Popeand J. O. Stiegler, TMS-AIME,Warrendale,(1990). 175.

R. G. Rowe: High temperature Aluminides and Intermetallics,

ed. by S. H. Whang,C. T. Liu, D. P. Popeand J, O. Stiegler,

TMS-AIME,Warrendale, (1990), 375.J. M. Larsen, K. A. Williams, S. J. Balsone and M. A. Stucke:

High Temperature Aluminides and Intermetallics, ed. by S, H.Whang,C. T. Liu, D. P. Popeand J. O. Stiegler, TMS-AIME,Warrendale, (1990), 521.

W. Cho, A. W.Thompsonand J. C. Williams: Metal/. Trans. A,21A (1990), 641

.

D. E. Albert and A. W.Thompson:Creep Behavior ofTi-24Al-l INb, Microstructure/Property Relationships in Titanium Al-loys and Titanium Aluminides, ed. by Y.-W. Kim, R. R. Boyerand J. A. Hall, TMS-AIME.Warrendale, in press.C. H. Ward, J. C. Williams. A. W. Thompson,D. G. Rosenthaland F. H. Froes: Proc. Sixth Int. Conf. on Titanium, Vol. 2, ed.

by P, Lacombe. R. Tricot and G. Beranger. Les Editions de

C 1991 ISIJ

ISIJ International, Vol. 31 (1991), No. 10

l9)

20)

21)

22)

23)

24)

25)

26)

27)

28)

29)

30)

31)

32)

33)

34)

Physique, Paris, (1989), I103.

T. Hamajima.G. Ltitjering and S. Weissmann:Meta[1. Trans, 4(1973), 847.

D, Banerjee, T. K. Nandi and A. K. Gogia: Scr. Metall., 21

(1987), 597.

D. Banerjee, A. K. Gogia, T. K. Nandi and V. A. Joshi: ActaMeta!!., 36 (1988), 871.

R. Strychor, J. C. Williams and W. A. Soffa: Metall. Trans. A,

19A (1988), 225.

R. Strychor: pPhase Decomposition in Ti-A1-Nb Alloys, Ph.

D. Thesis. Carnegie Mellon University, Pittsburgh, (1990),

N. E Paton and M. W. Mahoney:Meta!l. Tt'ans. A, 7A (1976),

1685.S. Takeuchi and A. S. Argon: J. Mater. Sci,, Il (1976), 1542.

A. M. Brownand MF. Ashby: Sc,'. Metall., 14 (1980), 1297.

B. Walser and.P. D. Sherby: Scr. Metal!., 16 (1982), 213.J. Weertman: Trans. Am. Soc. Met., 61 (1968), 681.

W.D. Nix and B. Ilschner: Strength of Metals and Alloys, Proc.

ICSMA5, ed. by P. Haasen, V. Gerold and G. Kostorz,

Pergamon,NewYork, (1980), 1507.

H. J. Frost and M, F. Ashby: Deformation MechanismMaps:The Plasticity and Creep of Metals and Ceramics, PergamonPress, Oxford, (1982), 11.

R. Raj and M. F. Ashby: Metal!, T,'ans., 2(1972), II13.

H, Ltithy, R. A. White and O. D. Sherby: Mater. Sci. Eng., 39(1979). 21 1.

F. W.CrossmanandM.F. Ashby: Acta Metall., 23(1975), 425.

H. Riedel: Fracture at High Temperature, Springer-Verlag, New

35)

36)

37)

38)

39)

40)

41)

42)

43)

44)

45)

46)

47)

York, (1987).

R. Lagneborg: Inl. Metall. Rev., 17 (1972), 130.

R. S. Mishra andD. Banerjee: Mate,', Sci. Eng., Al30 (1 990), 151.

G. Malakondaiah and R. RamaRao: Acta Meta!!., 29 (1981),

1263.

A. K. Mukherjee, J. E, Bird and J. E. Dorn: Trans. Am. Soc.

Met., 62 (1969), 155.

D. McLeanand M. H. Farmer: J. Inst. Met., 85 (1957), 41.

W. Cho: Creep Behavior of Near-cc Titanium Alloys, Ph. D.Thesis. The University of Michigan, AnnArbor, MI, (1987),

J. C. Williams: Titanium Science and Technology, Proc. SecondInt. Conf. on Ti, Vol. 3, ed, by R. I. Jaffe and H. M. Burte, eds.,

PlenumPress, NewYork, (1973), 1433.

J. C. Williams and G. Ltitjering: Titanium '80. Science andTechnology, Proc. Fourth Int, Conf, on Titanium, Vol. l, ed. byH, KimuraandO. Izumi, TMS-AIME,Warrendaie, (1980), 671.

H. A. Lipsitt, D. Shechtmanand R. E. Schafrik: Metal[. Trans.

A, 11A (1980), 1369.

B. J. Marquardt, G. K. Scarr, J. C. Chesnutt. C. G. RhodesandH. L. Fraser: Proc. Sixth Int. Conf. on Titanium, Vol. 2, ed, byP. Lacombe,R. Tricot andG. Beranger, Les Editions de Physique,Paris, (1989), 955.

K. S. Chan: Metall. Trans. A, 21A (1990), 2687.

D. McLean:J. Inst. Met., 81 (1952-53), 293.

S. J. Gittis and D. A. Koss: High-Temperature OrderedIntermetallic Alloys 111, MRSSymp. 133, ed. by C. T. Liu, A.I. Taub, N. S. Stoloff and C. C. Koch, MRS,Pittsburgh, (1989),

323.

C 1991 ISIJ 1146