E L S E V I E R Materials Science and Engineering A207 (1996) 170-180

M A T E R I A L S SCIENCE &

ENGINEERING

A

Microstructure and superplasticity of a T i - 2 5 A l - 1 0 N b - 3 V - 1 M o intermetallic alloy

D. Jobart, J.J. Blandin* Institut National Polyteehnique de Grenoble, Gknie Physique et MOcanique des Matkriaux, Unitk de Recherche Associ~e au C.N.R.S. n°793,

E.N.S. de Physique de Grenoble, BP 46, 38402 Saint-Martin d'Hkres Cedex, France

Received 24 July 1995

Abstract

Superplastic properties in tension of a Ti 25A1-10Nb-3V-1Mo (atY/o) alloy were studied in the temperature range 850 1050 °C under inert gas. The as-received material exhibits a fine elongated duplex structure with a significant rolling texture. Optimum superplastic properties are obtained for temperatures in the interval 930 ~< T~< 980 °C. At T = 930 °C, an equiaxed structure is produced in the first steps of deformation, and dynamic phase growth remains limited. A decrease in the initial texture is associated with the superplastic deformation. The values of the apparent activation energy in the superplastic regime were determined. The respective contributions of solute effects and phase volume fractions to the experimentally obtained activation energies are discussed.

Keywords: Superplasticity; Microstructure; Intermetallic alloy

I. Introduction

Titanium aluminides are promising replacements for conventional titanium alloys for structural applications in aeronautics because of their lower density, their higher stiffness as well as their better creep and oxida- tion resistance. Since the formability of these materials is very limited, superplastic forming (SPF) appears to be a promising route for fabricating complex shapes. The superplastic properties of titanium aluminides have been investigated increasingly in the past few years. Most of the studies dealt with Ti3Al-based alloys, which can be produced in sheet form through appropriate thermomechanical routes. Superplastic properties were firstly reported for the Ti 2 4 A l - l l N b (at.%) alloy [1], designated as ~2 alloy, and elongations up to 600% were obtained at 980 °C. Another Ti3Al-based alloy has been investigated more r e c e n t l ~ ' ( T i - 2 5 A 1 - 1 0 N b - 3 V - 1Mo, at.%), designated as super-~2 alloy, since this alloy exhibits more attractive mechanical properties at both room and elevated temperatures than the a2 alloy. Superplastic properties have also been reported for this

*Corresponding author.

0921-5093/96/$15.00 © 1996 - - Elsevier Science S.A. All rights reserved SSDI 0921-5093(95)10037-7

alloy [2 4] and elongations to rupture close to 800% have been reached at 980 °C. However, the mechanisms of superplastic deformation of this alloy have not yet been clearly established. In particular, large values of activations energies were reported in the superplastic regime which were not related to a well defined physical cause, except to solute effects in the soft fl phase.

The aim of this study was to investigate the super- plastic properties of a super-~2 alloy with particular attention given to determination of the creep parame- ters (such as the apparent activation energy). Since the superplastic properties of the alloy are expected to be closely related to the sizes, morphologies and volume fractions of phases in this alloy, these parameters and their variation during superplastic deformation were estimated and discussed.

2. Experimental procedure

The titanium aluminide used in this study was sup- plied by T I M E T corporation in the form of a 1.7 mm thick sheet. The alloy has the nominal composition T i - 2 4 . 4 A l - 1 0 . 6 N b - l . 0 M o - 3 . 0 V (at.%) and the oxy- gen content is 800 ppm, as given by the supplier. The processing route is not known in detail.

D. Jobart, J,J, Blandin I Materials Science and Engineering A207 (1996) 170-180 171

P1

~ r

Fig. 1. As-received super-ee 2 microstructure.

Mechanical testing at high temperature was carried out in tension under an argon atmosphere in the tem- perature range 850-1190 °C for various strain rates ranging from 5 x 10- 5 s 1 to 5 x 10 - 3 s 1 The defor- mation temperature was achieved in approximately 1 h and 30 min were required to homogenize the tempera- ture in the vicinity of the sample. After the tensile tests, the specimens were cooled at a rate of about 15 °C min-~ . Samples with a gauge length of 10 mm and a cross-section of 8.6 mm 2 were prepared with a spark cutting machine.

3. Results

3.1. Microstructure of the as-received material

The microstructures were characterized by scanning electron microscopy (SEM) on etched specimens (8% HNO3, 2% HF, 90% H20). Phase area fractions were measured by image analysis. The phase size was esti- mated from linear intercept measurements. Each inter- cept corresponds to an ~2-fl interface and consequently does not take into account the 0~2-0~ 2 and fl-fl boundaries. Measurements were carried out in three planes of observation: the sheet plane P~ and two perpendicular cross-section planes P2 and P3. In each plane of observation, mean intercepts were measured along several directions, allowing the determination of a minimum (respectively maximum) L ~ ~ (respectively L m~x) value. From the quantification of these intercept measurements, several assumptions concerning the shape of the grains have been made to allow the calculation of a mean phase size. These assumptions are detailed in Section 3.1. Pole figures limited to 75 ° tilt were deter- mined on a texture goniometer using Cu K~ radiation. The pole intensities were normalized by integrating over the whole measured range and dividing the particular intensity of one pole figure by this integral value.

The three-dimensional microstructure of the as-re- ceived material is shown in Fig. 1. The alloy consists of a fine mixture of e2 (dark) and fl (light) phases. The ~2 phase is rich in aluminium, whereas niobium, molybde- num and vanadium are concentrated in the fl phase. Fig. 1 reveals that both phases are elongated in the transversal and longitudinal directions. A lack of ho- mogeneity in the phase distrbution must be noted in the sheet plane, at the surface and in the bulk of the material. This phenomenon is reduced significantly in the cross-section planes.

The relative volume fraction of each phase is about 50%, as measured by image analysis. In the sheet plane P1, the microstructure has no preferential orientation (L]ni" ~ LTa~). In the cross-section planes P2 and P3, L~ ax ~ L_~ ~ ~ LT a× and L ~ in ~ L ~ in. These results in- dicate that the material was cross-rolled, at least par- tially, and that the microstructure can be modelled as a pile-up of parallelepipeds defined by the lengths L max 2 , L ~ ~x and L ~ m. The mean phase size L~2//j is therefore estimated from measurements in cross-section (PJ , ac- cording to:

l \ \ /

/ / J

"/ , b \ / . . . . . // \

( 1 \ \ ...,~;,o. / \~ " . .--"/

~-~---~- {22~-0}cz 2

Fig. 2. {200}fl and {22-40}~ 2 pole figures of the as-received alloy.

172 D. Jobart, J.J. Blandin / Materials Science and Engineering A207 (1996) 170 180

III

o8 980°C

930°C

0 6

840~C

1050°C 02

o

[] . . . . . Q _ ~ _ * - .

o _ " " ) ' ~

- ~ - g : z . . . . . . " X -

o< " --o

/ 0 o i l - I ) 5 i 0 0 E - 0 4 1 00E-03

Strain Rate (s q)

Fig. 3. Variat ions of strain-rate sensitivity m with strain rate.

L~2//~ = 34L~a in (L~nax)2 (1)

In the as-received conditions, L~2~ ~ 2.5 /,m. {200} pole figures in the fl phase and {2240} in the c~2 phase were obtained from the as-received sheet plane and are presented in Fig. 2. For the fl phase, a (I00) [011] component is obtained which corresponds to that usu- ally observed in rolled b.c.c, materials. For the ce 2 phase, (I12=0) [ I I I0] and (I 120) [0001] components are observed. These texture components have already been reported for a hot-rolled super-~2 alloy [5]. This indi- cates that both prismatic and basal systems were acti- vated in the ~2 phase during hot rolling and confirms that non-basal slip systems can be activated signifi- cantly in the ~2 phase during high temperature defor- mation.

3.2. Mechanical behaviour

3.2.1. Determination of the superplastic domain The effect of testing temperature on the dependence

of the strain rate sensitivity parameter m on the strain rate is shown in Fig. 3. These data are deduced from strain rate change tests. Except for low strain rates at a temperature of 1050 °C, a decrease in the strain rate sensitivity parameter m is observed when the strain rate is increased. Large values of m are obtained at interme- diate temperatures (930 °C and 980 °C). A maximum of 0.7 is reached when the material is deformed from 5 x 10 - s s ~ to 2 x 10 -4 S 1 at 980 °C, but large values of m (close to 0.6) are maintained when the temperature is lowered to 930 °C. It must be noted that the same value (m~0.5) , , i s obtained for these two temperatures at 4 x 10 4 s - l . When T < 9 0 0 °C, m decreases strongly with increasing strain rate and values close to 0.3-0.4 are obtained when g > 10 -4 S 1. Simi- larly, when the temperature is increased to 1050 °C, a significant decrease in the value of m ( ~ 0.3) is also measured.

From these results, it can be concluded that super- plastic properties are obtained in the temperature range 9 3 0 4 T~<980 °C. When T~> 1000 °C, a decrease in m is observed, which is probably associated with plastic deformation of the phases.

3.2.2. Effect of the temperature of deformation at k = 5 x l O 4s-1

The effect of testing temperature on the (a,e) curves at a strain rate of 5 x 10 4 s - - 1 and up to a total strain of 1.0 is shown in Fig. 4 in the temperature range 890~<T~< 1070 °C. Strain softening is obtained at T = 890 °C at strain larger than 0.3, which results partly from the development of a neck in the gauge length of the strained specimen. This result is in agree- ment with the low value of m ~ 0.3 under these condi- tions. At 930 °C and 950 °C, softening is still observed but is not related to necking development. At T ~ 980 °C, a constant flow stress is reached when deformation exceeds 0.1 and is roughly maintained during the whole test. At T,~ 1020 °C, a significant strain-hardening is observed. When the temperature is increased further, some strain-hardening can be detected but it remains limited.

3.2.3. Effect of strain-rate at T= 930 °C Fig. 5 shows the influence of the strain rate on the

stress vs. strain curves of the alloy deformed at a temperature of 930 °C up to a total strain of 1.0. Except for the lower strain rate (10 -5 s - l ) , strain softening is observed systematically. When the strain rate is increased, the magnitude of the softening is also generally increased. Data concerning superplastic defor- mation of a super-~2 alloy strained at a strain rate of 2 x 10 4 s 1 and at 940 °C also show a softening behaviour during deformation [3]. Except for the

10!)

140

12i)

I OU

,C~ ,gel

~l 60

40

2 0

890°C

-- . . . . . . . . . . . . . 930°C 950o C

~ " 1020oC ~ 1050oC

1070oc

0 0 2 04 0 6 08 1 1 2 Strain

Fig. 4. Effect of temperature on the a - g characteristics of super-e 2 at a s t r a i n rate of 5 x l 0 - 4 s 1.

D. Jobart, J.J. Blandin / Materials Science and Engineering A207 (1996) 170 180 173

250

200

p.~ 150

100 ra~

50

5.10 -3 S "1

5.10 "4 S -1

?1~ 0-4 -i -~S-1 s

---:; 7 7 ~ 7 ~ ~ I ° - ! s-~ f 0.2 04 0.6 08 1 1,2

Strain

Conversely, when the material is tested at T = 1050 °C (Fig. 6(c)), the microstructure is strongly modified during the homogenization period. An important de- crease in the :¢2 phase volume fraction occurs and the elongated morphology of the phases is no longer ob- served. The size of the ~2 phase increases whereas fl grain growth occurs. A needle-like e2s phase is also observed in the fl phase but it is likely to appear during the cooling (fl--* fl + gZs) step which follows holding at high temperature. Consequently, the 72s phase is as- sumed not to be present at 1050 °C.

Fig, 5. Effect o f s t ra in ra te on the ~r-c charac ter i s t ics o f super-~ z al T = 930 °C.

highest strain rate (5 x 10 3 S - - 1 ) for which strain softening results mostly from necking, no detectable necking was observed at a strain of 1.0 when ~ < 2 x 1 0 3s 1

3.3. Microstructural characterizations after deformation

3.3.1. Effect of temperature on microstructural transformations under static conditions

A way to estimate the microstructural transforma- tions which occur respectively under static (tempera- ture) and dynamic (temperature and deformation) conditions is to compare the microstructural changes in the undeformed head and in the deformed gauge length of the specimen. Moreover, significant microstructural transformations can take place during the time required to reach the testing temperature and particularly to homogenize it. It appears necessary to characterize the microstructure of the alloy when the tensile test starts. To estimate such transformations, heat treatments cor- responding to the pretesting heating sequence to each temperature were carried out. These resulting mi- crostructures are labelled "homogenized" in the text.

During the homogenization period, depending on the temperature of deformation, significant microstructural transformations can take place, as it is shown in Fig. 6. When the material is tested at T = 930 °C (Fig. 6(a)), the microstructure, when deformation starts, is very similar to the as-received microstructure: no phase growth has been induced and the phases are still elon- gated in the longitudinal direction. When the deforma- tion temperature is increased to 980 °C (Fig. 6(b)), some coalescence of the phases occurs before deforma- tion, leading to a coarser but less elongated microstruc- ture than in the as-received condition. However, the extent of the microstructural transformation remains limited, in particular no significant variation in the phase volume fractions is observed.

a

b

C

Fig. 6. Effect o f the t empera tu re of de fo rma t ion on the micros t ruc- ture of the al loy when de fo rma t ion starts: (a) T = 930 °C, (b) T = 980 °C, (c) T = 1050 °C.

174 D. Jobart, J.J. Blandin / Materials Science and Engineering A207 (1996) 170-180

a

ited variations in the phase volume fractions. At T = 1050 °C, no additional significant grain growth of the c~2 phase occurs in the head of the sample in reference to the homogenized condition.

3.3.2. Effect of temperature on microstructural transJormations under dynamic conditions

Fig. 8 shows the effect of the deformation tempera- ture on the microstructures in the gauge lengths of the specimens ( c = l.0) at g = 5 x 10 4s ~. On the micro- graphs, the stress axis is horizontal. At T = 930 °C (Fig.

b

a

Fig. 7. Effect of temperature on annealed microstructures (head of the sample, c = 1, ~ = 5 x 10 4 s 1): (a) T = 9 3 0 °C, (b) T = 9 8 0 °C, (c) T = 1050 °C.

Fig. 7 shows the effect of the deformation tempera- ture on the microstructures in the heads of the de- formed specimens (c = 1.0) at i = 5 x 10 4 s 1. It must be remembered that the deformation time required to reach a strain of 1.0 is approximately 30 min for the chosen experimental conditions. For the three tempera- tures investigated (930 °C, 980 °C, 1050 °C), the mi- crostructures are very similar to those observed after the homogenization periods. The elongated structure is still roughly maintained when the material is tested at T = 930 °C and T = 980 °C, associated with very lim-

13

(3

Fig. 8. Effect of temperature on superplastically deformed microstruc- tures (gauge length, e = l , ~ = 5 x 10-4 s ~): (a) T = 9 3 0 °C, (b) T = 980 °C, (c) T = 1050 °C.

D. Jobart, J.J. Blandin / Materials Science and Engineerin~ A207 (1996) 170-180 175

a b

c d

Fig. 9. Effect of strain rate on the deformed microstructure at e -1.0: (a) ~= 10 ~ s- ~, (b) k= l0 4 s ~, (c) ~= 5 x 10 4 S I ( d ) ~ - 2 x l 0 3s i.

8(a)), microstructural t ransformat ions remain relatively limited but an equiaxed structure has been developed and preserved during the test, not associated with any impor tan t concurrent grain growth. At T = 980 °C (Fig. 8(b)), an equiaxed microst ructure is also devel- oped but with a slightly larger phase size than that obtained after testing at T = 930 °C. This means that some dynamic grain growth takes place when the mate- rial is tested at T = 980 °C. Dynamic grain growth has already been reported during superplastic deformat ion o f the super-~ 2 alloy at this temperature [3,6]. It results generally in strain hardening. This is not the case for the alloy investigated in this study, since a relatively constant flow stress is maintained during the whole test. When the temperature o f deformat ion is increased up to 1050 °C, addit ional t ransformat ions are observed in the gauge length o f the specimen (Fig. 8(c)). Intense dynamic grain growth takes place, leading to nodules o f ~2 phase o f size larger than 10 /~m. Moreover , an

alignment o f the microstructures along the strain axis is observed.

3.3.3. Effect o f strain rate on microstructural

transformations at T = 930 °C Fig. 9 shows the microstructures o f the gauge lengths

o f specimens strained up to e = 1.0 at 930 °C for different strain rates. On the micrographs, the stress axis is horizontal . The deformat ion time ranges be- tween 7 m i n ( 2 x 10-3 s ~) and 28 h (10 5 s ~). For the lowest strain rate (10 5 s-~) , the microstructure after deformat ion is significantly coarser than before deformat ion but an equiaxed structure has been devel- oped and preserved. For the highest strain rate (2 x 10 3 s - ~), some dynamic refinement o f the struc- ture seems to occur, but the resulting microstructure is relatively elongated along the stress axis, with curved ~2 fl phase boundaries. For the strain rates o f 10 4 s ~ and 5 x 10 4 s ~, the phase size does not evolve

176 D. Jobart, J.J. Blandin / Materials Science and Engineering A207 (1996) 170 180

significantly during deformation and an equiaxed struc- ture is produced and preserved up to a strain of ~ = 1.

4. -Discussion

It is well known that optimum superplastic properties are associated with high values of the strain-rate sensi- tivity parameter m (generally m ~> 0.5) and a minor contribution of intragranular plasticity (by diffusion or dislocation creep) to macroscopic deformation. The main mechanism of deformation is thus sliding along interfaces (grain and/or phase boundaries), which is promoted by a fine, stable and equiaxed microstructure. From the mechanical tests performed on the investi- gated alloy, it can be concluded that the optimum superplastic properties correspond roughly to a temper- ature range from 930 °C to 980 °C, in a relatively large strain rate domain from about 5 x 10 5 s - i to 5 x 10-4 s ~. These properties are associated with a strain rate sensitivity parameter close to 0.6 and with the development and preservation of a fine and equiaxed structure. These results are in good agreement with previously reported data concerning superplastic properties of Ti 2 5 A l - 1 0 N b - 3 V - 1 M o alloy [2 4]. When the deformation temperature is too high (i.e. at T = 1050 °C), straining generates an alignment of the microstructure along the stress axis, which indicates that plastic deformation of the phases contributes sig- nificantly to the total deformation. It is also associated with a decrease in the value of the strain-rate sensitivity parameter (m = 0.3).

A way to obtain data concerning the respective con- tributions of phase boundary sliding and intragranular plasticity (under the assumption that this plasticity involves intragranular slip) to the macroscopic defor- mation is to perform texture measurements. Fig. 10 shows the {200} pole figure of the fl phase and {22740} figure of the 0~ 2 phase after deformation carried out at 950 °C (Fig. 10(a)(b)) and at 1020 °C (Fig. 10(c)(d)). In the superplastic regime ( T = 950 °C), a drastic decrease in all the as-received texture components, (100) [011] for the fl phase and both (1120) [ i l l0] and (1120) [0001] components for the 0~ 2 phase, is observed. This reduction in the initial texture during superplastic de- formation confirms the predominant role of grain movements as the main deformation mechanism. It indicates also that both phases are concerned with these movements. When the material is tested at a tempera- ture higher than 1000 °C (i.e. T = 1020 °C), the varia- tion of the texture components is different for the two phases: the (100) [011] component for the fl phase is reinforced whereas the components for the ~2 phase are only weakly affected by deformation. This result sug- gests that most of the plastic deformation of the mate- rial is concentrated in the fl phase.

For the investigated material, at a strain rate of 5 x 10 -4 S ], a peak stress is obtained in the first steps of superplastic deformation (~ < 0.2) when the tempera- ture is lower than 980 °C (see Fig. 4) and a constant flow stress is maintained at T = 980 °C. In the case of previously studied super-z2 alloys, a significant harden- ing has been reported [2] at T = 950 °C for a strain rate of 6 x 10 - 4 s ] up to as t ra in of 1.0. At T = 9 8 0 °C, strain hardening was also observed at low strain rates (5 x 10 s s - l and 1.5 x 10 -4 S ]) and was attributed to strain-enhanced grain growth [4], whereas at high strain rates (4 x 10 -4 S l), softening was observed only after a strain of about 0.7. Softening was then attributed to dynamic recrystallization [4].

As already mentioned, during the superplastic defor- mation of the investigated alloy, the as-received elon- gated structure is converted progressively into an equiaxed structure. To quantify this microstructural transformation, measurements of the variation with strain of the shape parameter defined as the ratio

m a x m i n L2 /L2 were carried out. From Figs. 7(a)(b) and 8(a)(b), this ratio is decreased from approximately 1.7 in the as-received condition to about 1.4 + 0.05 after superplastic deformation in the temperature range 930 ~ T ~< 980 °C. This process requires, among others, fragmentation of the 0~ 2 phase in the first steps of deformation. Such a process has already been reported for conventional titanium alloys [7,8]. It is attributed to penetration of the fl phase along subboundaries which are formed in the hexagonal phase during the forging

a b

\ -U*.

, '\ , ,\ // \\ /'

\ "..

. i - \ \ /

. '\ // . ',,

(' ~...~ ...~. i .... \ .;,

.~ . ° . . \

" \

\ / /

I I

,\

/ - - ~ - - / / ~--- . \

/ S ' - )

• o ~t ;7- c"~:' ~

"\ ,

d

Fig. 10. {200}fl and {22-40}~ 2 pole figures after deformation: (a)(b) after superplastic deformation (T = 950 °C), (c)(d) after plastic defor- mation ( T = 1020°C).

D. Jobart, J.J. Blandin / Materials" Science and Engineering A207 (1996) 170 180 177

2I_

A e'm

g V

1 0

8

4

2

0

, , r [ ~ ,

i

(Ivnamic - /

J

/

o / / ' "

[] • / ~ , / / s t a t i c

Q - / /

_~.*. 5.1. . . . . . . . . . . . . . . . . . ( Ig ' - I 'CCCIFCJ

g 5 0 9 0 0 9 5 0 1 0 0 0 1 0 5 0

Temperature (°C)

Fig. 11. Mean phase size vs. temperature in the head of the sample (static conditions) and in the gauge length (dynamic conditions) at g = 5 x l O 4 s 1.

step [8]. Moreover, it can result in a peak stress in the (a, ~) curves [7].

The variation of the mean phase size L~2/, with deformation temperature ( ~ = 5 x 10 4 s l, t ~ 3 0 min) under both static and dynamic conditions is shown in Fig. 11. Two domains can be distinguished. When the temperature is lower than 980 °C, L~2//~ is not affected significantly by temperature, even under dynamic con- ditions. This means that the fragmentation process which occurs during the first steps of deformation in this temperature domain is likely to be associated with some phase coalescence since it does not induce a refinement effect, as far as can be detected through the measurement of L~2//j. This also indicates that the soft- ening process cannot be interpreted directly in terms of a reduction in the mean phase size. Moreover, at T = 980 °C, dynamic grain growth takes place since an increase of about 20% is obtained for L~2//~. Neverthe- less, this increase in the mean phase size does not result in hardening, as could be expected from the constitutive law which is generally used to describe superplastic deformation:

In this equation, ~ is the applied strain rate, A is a constant, /t is the shear modulus, b is the Burgers vector, Q is the activation energy, d is the grain size (which is assumed to be equal to L~2//~), p is the grain size exponent, and R and T have their usual meanings. In the case of T i - 6 A 1 - 4 V (wt.%) alloy, values of p = 1.7 [9] and p = 2.3 [10] have been reported, which is in relative agreement with models proposed for super- plastic deformation of two-phase alloys [11,12]. At T = 9 8 0 °C and for a strain rate of 5 x 10 4 s - l , for

which m ~ 0.5, the phase growth during deformation up to a strain of 1.0 (AL~2/, ~ 20%) would result in an equivalent increase in the flow stress (Aa ~ 20%). Such an increase is not observed experimentally since the flow stress remains nearly constant. It confirms the key influence played by the morphology of the phases on the mechanical response of the alloy and the difficulty in accounting satisfactorily for the superplastic be- haviour by using relation (2).

The determination of the apparent activation energy Q~gp in the superplastic regime of the investigated alloy is difficult since, depending on the temperature of defor- mation, a variation with strain of the flow stress can be observed. Nevertheless, in the temperature range 890 ~< T ~< 980 °C, the strain rate sensitivity parameter is relatively constant (m ~ 0.5 + 0.05) for a strain rate equal to 5 x 10-4 s - ] and the grain size is also roughly constant when deformation starts, as illustrated by Fig. 11. The dependence on temperature of the shear mod- ulus of the two-phase alloy has been assumed to be equal to that of the/~ phase [13]. When the peak stress is considered, a value of Q~pp ~ 340 + 30 kJ m o l - 1 is found (see Fig. 12). This value can be compared with the results of Yang et al. [6], who obtained different values of the activation energy depending on the ap- plied stress for superplastic deformation of a super-~2 alloy. In the low stress regime, a value of about 270 kJ mol l was reported whereas a value close to 350 kJ mol ~ was obtained for high stresses.

The apparent activation energy of Ti3A1-Nb based alloys has been widely studied under creep conditions [14 16]. Mendirat ta and Lipsitt [14] have studied near- stoechiometric Ti3A1 and Ti3A1-10wt.%Nb in the tem- perature range 550 ~< T~< 825 °C for stresses ranging from 69 MPa to 312 MPa. When the temperature of deformation is higher than 650 °C, they show that the addition of niobium increases the apparent activation

_=

T p

75

peak stress

mQ = 170 kJ/lnol /

steady-state regime

mQ = 85 k J/tool

ot~ + [3 field

- 8 , , , , i , t , , i , , , ,

7 8 8.5 9

10000/T (K "])

Fig. 12. Estimation of the apparent activation energies in the super- plastic domain at ~ = 5 × 1 0 4 s '.

178 D. Jobart, J.J. Blandin / Materials' Science and Engineering A207 (1996) 170-180

-6

-7 mQ = 60 kJ/mol

[3 field

v J

-8 6.5 7

IO000/T (K ~)

T~

75

Fig. 13. Estimation of the apparent activation energy in the fl phase domain (1090 < T < 1190 oC) at ~ = 5 x 10 4 S i.

energy from 206 kJ mo l -1 to 285 kJ mol 1. From the values of the stress exponents and despite the lack of data concerning the diffusion coefficients, Mendiratta and Lipsitt considered that dislocation climb was the most probably rate-controlling mechanism of deforma- tion. In the case of Ti 2 5 A l - 1 0 N b - 3 M o 1V alloy deformed in the temperature range 650 ~< T~< 870 °C, an apparent activation energy of 140 kJ tool -1 was determined at low stresses [15] whereas a value of 305 kJ mol ' was deduced at high stresses [16]. The first regime corresponds to diffusional creep, since a strain rate sensitivity parameter m close to unity was ob- tained, whereas the large value of Q~p is associated with a climb-controlled creep regime (m ~ 0.25). The difference between the values obtained for Ti3A1- 10wt.%Nb [14] and Ti-25A1 1 0 N b - 3 M o - I V [16] un- der dislocation creep was attributed to the large amount of fl-phase stabilizers.

Da ta concerning this effect can be obtained from estimation of the apparent activation energy in the single-phase fl domain. Additional mechanical tests were carried out in the temperature range 1090~< T~< 1190 °C. In this domain (T~> 1100 °C), a value of rn equal to 0.3 + 0.05 is measured systemati- cally at ~ = 5 x 10 4 s - ' and is attributed to disloca- tion creep for which p = 0. The apparent activation energy Q~pp w a s estimated, taking into account the dependence on temperature of the shear modulus ~t/j in the fl phase [13], and it is found that Q ~ p p ~ 2 0 0 - t - 3 0

kJ m o l - l (see Fig. 13). The activation energy Q/sTi of lattice diffusion in fl-Ti is equal to 150 kJ mo l -~ [13]. The difference between Q~pp and Q/~Ti tends to confirm the effect of the fl stabilizers on the creep behaviour of the Ti-25A1 1 0 N b - 3 V - 1 M o alloy, particularly of the slow diffusing solutes, such as molybdenum.

Nevertheless, the value Q~pp obtained in the fl field remains much smaller than the apparent activation energy Q~pp measured in the superplastic regime (i.e. in the 0~ 2 -~-fl domain). As a consequence, the high value of Q~pp cannot be attributed completely to a solute effect in the fl phase, even if the amount of solute in this phase is increased by decreasing the temperature [17]. The contribution of the c~ 2 phase to the superplas- tic deformation of the investigated alloy should proba- bly be taken into account.

In the case of superplastic deformation of the super- ~2 alloy, the mechanism of accommodat ion of inter- phase sliding is not yet clearly established. In particular, the respective contributions of intragranular plasticity of the hard (~2) and soft (fl) phases to the macroscopic deformation have to be specified. This aspect of super- plastic deformation of a two-phase alloy was studied and modelled by Su6ry and Baudelet [12] in the case of c~-fl brasses. In the so-called hydrodynamical model, they described the superplastic behaviour of the two- phase material as the flow of a viscous non-Newtonian fluid (i.e. the soft phase) containing non-deformable particles (i.e. the hard phase). In such conditions, all the accommodat ion mechanisms of phase boundary sliding are concentrated in the soft fl phase. In the case of conventional titanium alloys, it seems that the hard phase cannot be considered as non-deformable during superplastic deformation [18]. For the super-~ 2 alloy, since the gap between the self-diffusion coefficients of the fl and 0~ 2 phases is larger than the gap between the fl and ~ phases in conventional titanium alloys, the plastic deformation of ~2 in the superplastic regime is expected to be significantly smaller than that of the phase in conventional titanium alloys. Nevertheless, the peak stress is likely to be related, at least partially, to a fragmentation of the ~2 phase. Some localized deforma- tion is therefore expected to take place in the hard phase and thus to affect the mechanical response of the two-phase material and its apparent activation energy. Additional work is, however, needed before any definite conclusion.

Moreover, as previously mentioned, the value of Q~pp has been calculated by considering the peak stress. I f the same calculation is carried out with the steady- state flow stress which is reached after a strain of about

= 0.6 in the temperature range 930 ~< T ~< 980 °C, it is found that rnQ~ pp ~ 85 kJ m o l - ' (see Fig. 12). Since the strain-rate sensitivity parameter m can vary with strain, the estimation of Q~pp remains quite delicate. From Fig. 5, rn can be calculated at various strains for strain rates in the range 10 4 < k ~ < 1 0 3 s ~. In the first steps of deformation, m ~ 0.45 + 0.05, which confi- rms the results obtained during the velocity jump tests. After a strain of 0.7, m appears to be in the interval 0.35 ~< m ~< 0.45. These values lead to Q~pp ~ 180 + 30 kJ mol 1. Since the activation energy Q~pp measured in

D. Jobart, J.J. Blandin / Materials Science and Engineering A207 (1996) 170 180 179

the single-phase domain lies in this interval, this may indicate that most of the plastic deformation is concen- trated in the soft fl phase, as generally admitted in the case of ~-f l brasses [12]. Nevertheless, these results show that complementary data must be obtained before reach- ing a precise conclusion concerning the interpretation of the value of the activation energy during superplastic deformation of the Ti 25Al-10Nb 3V-1Mo alloy.

An additional difficulty in relation to the interpreta- tion of the value of Q~s p£ results from the variation with temperature of the phase volume fraction. Fig. 14 shows this variation for the ~2 phase volume fraction in the temperature range [890 °C;l l00 °C]. This curve was determined by two independent methods (image analysis and variation of the compositions of the phases), the details of which are reported elsewhere [17], and corre- sponds roughly to thermodynamical equilibrium in ref- erence to the as-rolled microstructure. This curve shows three domains. When T > 1070 °C, only the fl phase is detected. In the interval [950 °C;1050], the phase volume fractions are strongly dependent on the temperature. This dependency is significantly lower when the temper- ature is still decreased (T < 950 °C), the two phases showing then quite similar phase volume fractions.

In the case of Ti3Al-based alloys (respectively conven- tional titanium alloys), two kinds of ~2 (respectively ~) phase are generally distinguished. The primary ~2 p is produced during the first treatment performed in the two-phase domain after a prior treatment in the fl field. When further heat treatments are carried out at lower temperatures, additional 0~ 2 phase is nucleated and is described as secondary ~) phase. This phase can coalesce with pre-existing primary ~P grains or grow inside fl grains. It must be underlined that the data of Fig. 14 do not take into account the ~ inside the fl grains. In the case of the investigated alloy, the nucleation of sec- ondary ~ is expected to occur only when T < 9 5 0 °C.

70

~, 60

50

40

-~ 30

20

~< 10

850 950 1050 1150

Temperature ('C) Fig. 14. Variation with temperature of the ~2 phase volume fraction.

In this temperature range, the data plotted in Fig. 14 thus underestimate, probably to a quite limited extent, the total ~2 (i.e. ~P + ~ ) phase volume fraction.

Fig. 14 demonstrates that optimum superplastic prop- erties are achieved for an ~2 phase volume fraction of about 50%. This confirms the similarities between the superplastic deformation of the super-~2 alloy and of the conventional titanium alloys, for which this optimum has been widely demonstrated to coincide generally with approximately equal phase proportions [19,20].

Fig. 14 shows also that superplastic properties are maintained in a temperature domain which corresponds to various phase volume fractions, namely from about 40% to 60% ~2 phase. This aspect must be taken into account in the interpretation of the apparent activation energy which corresponds to the superplastic regime. From a qualitative point of view, it is expected to increase Q~pp since a decrease in the temperature of deformation is associated with an increase in the ~2 phase volume fraction. Complementary work is in pro- gress to estimate more quantitatively this effect o n Qasp, p.

5. Conclusions

Superplastic properties of a Ti-25Al 10Nb 3V- 1Mo (at.%) alloy were observed in the temperature range 930 ~< T~<980 °C and for strain rates ranging from 5 × 1 0 Ss ~ t o 5 × 1 0 4s l. At T = 9 3 0 ° C , a p e a k stress is obtained during the first steps of superplastic deformation, which is related to fragmentation of the ~2 phase to produce an equiaxed structure. This equiaxed structure is preserved after a strain of e = 1.0 and no significant dynamic grain growth occurs. For such con- ditions, the initial rolling texture is reduced in both phases during superplastic deformation. The apparent activation energy Q app w a s estimated under different conditions: in the fl domain (Q~PP) and in the ~2 + fl field (Q~.ppp), for the peak stress and for the steady-state flow stress. It is shown that Q~.pp varies with strain in the superplastic regime. This regime corresponds to an ~2 phase-volume fraction of about 40 50 vol.%.

Acknowledgements

The authors acknowledge Dassault Aviation and Di- rection des Recherches, Etudes et Techniques (D.R.E.T.) for financial support.

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