Acta metall, mater. Vol. 40, No. 3, pp. 581-587, 1992 0956-7151/92 $5.00 + 0.00 Printed in Great Britain. All fights reserved Copyright © 1992 Pergamon Press pie

MECHANICAL BEHAVIOUR OF FINE GRAINED TiA1 INTERMETALLIC COMPOUND--I. SUPERPLASTICITY

R. M. IMAYEV, O. A. KAIBYSHEV and G. A. SALISHCHEV Institute for Metals Superplasticity Problems, U.S.S.R. Academy of Sciences, Ufa 450001, U.S.S.R.

(Received 4 March 1990; in revised form 4 March 1991)

Abstract--TiA1 was used to show that superplasticity is possible in intermetallic compounds with a high ordering energy. Mechanical properties and structural changes are found to be strongly affected by the original grain size and grain boundary structure. It was established that the regularities of superplastic flow in TiA1 are those typical of the structural superplasticity in metals.

R~mn6----On utilise du TiAI pour montrer que la superplastieit~ est possible dans les compos6s interm~talliques ~ grande 6nergie de mise en ordre. On trouve que les propri~t~s m&..aniques et les changements de structure sont fortement affect~s par la taille du grain initial et par la structure des joints de grains. I1 a ~t~ d6montr~ que les propri~t6s de I'~¢oulement superplastique dans TiA1 sont typiques de la superplasticit6 structurale dans les m~taux.

Zusammenfassung--An TiA1 wird gezeigt, dab Superplastizit/i.t bei intermetallischen Legierungen hoher Ordnungsenergie m6glicb ist. Mechanische Eigenschaften and strukturelle ,~.nderungen werden stark yon der Korngr6Be und der Korngrenzstruktur zu Anfang beeinflul3t. Es wird gezeigt, dab das superplastische Flie~n yon TiA1 Gesetzen folgt, die typiscb ffir die strukturelle Superplastizit/it yon Metallen sind.

1. INTRODUCTION

The majority of polycrystalline intermetallics are known [1-3] to be brittle at low temperatures and display limited plasticity at higher temperatures. Their brittleness was explained in various ways, the main reasons being: a limited number of slip systems, impeded cross-slip, and difficulty of the strain transfer from one grain to another, formation of dislocation barriers typical of ordered alloys, and segregation of harmful impurities in grain boundaries [1-10]. At the same time, it is well known that the mechanical behaviour of metals and alloys can be controlled by varying their structural state, in par- ticular, the area and structure of grain boundaries [11-13]. For example, the refining of grains abruptly decreases the threshold of low-temperature brittle- ness and induces superplasticity in a higher tempera- ture interval [14, 15]. On the other hand, the presence of special grain boundaries in the microstructure of metals and alloys affects superplasticity impeding its development [15]; the role played by special grain boundaries under ductile-brittle transition being practically unknown. Hence, it would be interesting to study mechanical properties of intermetallic with large areas and different structures of grain boundaries.

Let us first consider the regularities of the mechan- ical behaviour of intermetallic at high temperatures.

So far, the nature of superplastic flow has been studied mainly on metal-base samples [15-17]. Thus,

the problem of transforming materials traditionally considered to be undeformable, like intermetaUics, to superplastic state is of current importance. There are several reasons why it is so important to prove the possibility of making these compounds superplastic and to study the conditions of its being displayed. First, it is important to know if the regularities of the superplasticity deformation of intermetallic com- pounds correspond to those typical of the structural superplasticity of metals. Second, the knowledge of the peculiarities of microstructural changes during the superplasticity deformation of intermetallics may be useful to control their mechanical properties. No systematic research has been undertaken in these areas so far.

A brittle intermetallic compound TiAI (superlattice L10) was selected for the investigation. This com- pound possesses properties typical of the majority of intermetallics: it is brittle and hard-to-deform even when processed at high temperatures [18, 19]. Since it preserves its ordered structure up to the melting point, no domain structure is formed, which could otherwise make the study more complicated [5]. In earlier papers [20, 21] it was reported that during the hot straining of TiAI its microstructure was easily refined and further heat treatment yielded a lot of annealing twins, with a large total area of grain boundaries being preserved [22, 23]. The latter fact is rather interesting since it allows to find out the dependence of the superplasticity of intermetallics on both the grain size and boundary structure. It is well

581

582 IMAYEV et al.: MECHANICAL BEHAVIOUR OF FINE GRAINED TiAI--I

known [15, 24] that in metals both of these micro- structural parameters exert considerable influence on their superplasticity behaviour.

2. MATERIAL AND EXPERIMENTAL

Table 1. Temperature dependence of mechanical properties of TiA1 in state A (i =8.3 x 10-(s - I )

Mechanical tda. (°C) properties 900 950 975 1000 1025 1050 6 (%) 155 175 195 215 250 175 050 (MPa) 290 225 200 ! 80 165 150

The alloy composition is Ti-35.9wt% AI. The sample was shaped as a Q90 x 120mm rod pro- duced by the compaction of granules. It was sub- jected to 80% compression strain at 1000°C and a strain rate of ~ = 10- 3 s - l whereafter it cut into bars that where annealed partly at 930°C for 5 h (state A) and partly at 1050°C for 2 h (state B). The bars were cooled with the furnace. Then, flat samples with working dimensions of 10 x 5 x 2 mm were cut out of the bars.

Mechanical tensile tests were performed on an Instron machine. At first, the state A samples were deformed in a temperature range of 900-1050°C at an initial strain rate of ~ = 8.3 x 10-4s -~. Then, the samples in both the states were tested at 1025°C and initial strain rates of ~ = (1.6-83) x 10 -4 s - l .

The maximum elongation to rupture 6 and true flow stress at a relative elongation of 50% Os0, were determined based on the tensile diagrams. The strain rate sensitivity coefficient m was evaluated by the log ~rso-log ~ curves slopes.

To assess a mean grain size ~7 and the volume fractions of phases and to plot the histogram of the grain size distribution, the linear intercept method was used, with the number of measurements in each case being 300. To evaluate the mean grain size, the annealing twin boundaries were taken into account.

The electron microscopy analysis was conducted on Tesla BS-540 with an accelerating voltage of 120 kV. The density of dislocations p was evaluated by the linear intercept method by considering from 20 to 25 negatives.

The absorption properties of grain boundaries were examined by measuring the temperature of relaxation tr ("blurring" of the electron microscope contrast) of trapped lattice dislocations in grain boundaries. To define t,, the state B samples sized O10 x 10ram were 1% compression strained and then annealed for 15 min at different temperatures.

3. RESULTS

3. I. Mechanical propert ies

The temperature dependence of mechanical prop- erties of the state A compound is rather complicated (see the Table). With the rise of the testing tempera- ture, the relative elongation to rupture 6 first gradu- ally increases, but after reaching its maximum (6 = 250%) at 1025°C it abruptly drops. The flow stress decreases monotonously with the rise of temperature.

The strain rate dependence of the TiA1 mechan- ical properties in the studied states at 1025°C is shown in Fig. 1. In a certain strain rate interval, state A contrary to state B display an enhanced strain

rate sensitivity of the o50 flow stress [Fig. l(a)]. Thus, in a strain rate interval of (0.83-1.6) x l0 -3 s -1 the values of m in state A vary from 0.33 to 0.43 [Fig. l(b)]. The higher m value is matched with the larger relative elongation to rupture (6 = 200-250%). Samples deformed under these conditions display monotonous elongation to failure. In state B, coefficient m is constant throughout the considered rate interval, its value being 0.26, while the relative elongation depends on the strain rate only slightly (6 = 130-160%); the samples being deformed non- uniformly. Hence, in state A, TiAl displays superplas- tic features while in state B superplastic behaviour is less pronounced.

Figure 2 shows curves cq - , for tensile straining of the state A and B samples at an optimal strain rate (~ =8.3 x 10-4s-t) . It can be seen that the steady flow stage in the state A sample is preceded by a prolonged stage of strengthening. Similar curves arc characteristic of the state A samples at low strain

(el

2,50

200

100

80:

"-'-" . ~ ~ : ~ "

10-4 10-5 (s -1 )

i

10-2

25O

2OO

150 ~_~

100

5O

0

(b) 0.5

0.4

0.3

0.2

0.1

t 10-4

. / i J

10-3 10-2

(s -~) Fig. 1. Dependence of flow stress #50 and relative elongation to rupture 6 (a), and strain rate sensitivity coefficient, m (b) o n the rate of TiAI deformation, i: • state A; O state B

(t = 1025°C).

IMAYEV et al.: MECHANICAL BEHAVIOUR OF FINE GRAINED TiAI--I 583

150.

I 00.

¢ 50,

0 10 20 30 40 50 60 70 80

E (*1.~

Fig. 2. Stress-strain curves of TiAI for states A (1) and B (2) (t = 1025°C, ~ =8.3 x 10-4s-I).

rates. In state B, at the initial stage of deformation, a peak of flow stress is observed. With further straining, the flow stress monotonously decreases and approaches the cr~ level in state A. A similar curve is also observed at high strain rates.

It should be noted that similar mechanical behaviour is observed under superplasticity defor- mation in the Ni3AI intermetallic compound with a mean grain size of 1.6#m at 700°C and

< 1.04x 10-4s -1 (t~ = 160%) [25].

3.2. Microstructure

In state A, TiA1 has a homogeneous fine grained microstructure with a mean grain size of 5 pro. On the contrary, the TiAl microstructure in state B has a remarkable diversity of grain sizes, which is due to a nonuniform distribution of the second phase (ct2-phase), i.e. Ti3AI (superlattice D0~9) its fraction being no more than 3% of the volume. The volume fraction of "fine" grained and "coarse" grained phases constitute 15 and 85% respectively. The sizes of "fine" grains vary from 2 to 8/*m while the average size of "coarse" grains is 15 #m.

The metallographic analysis of sample micro- structures in each of the states immediately before straining after their heating up to 1025°C and holding at this temperature for 30 rain indicates that in state A contrary to state B some of the grains coarsen, this producing a relatively coarse grained fraction with a mean grain size of 12/~m constituting 85 vol.% like in state B. In both the states the size of "fine" grains is the same: from 2 to 8/~m. On the whole, the considered states of the sample can be defined as having very similar grain sizes. In both the states grains are equiaxial. No differences in the quantity, distribution, sizes (0.1-31*m), and shape of the ct2-phase particles were observed. There are twin boundaries in the microstructures of the sample states too. Their fractions in state A and B are 10 and 35% of total number of grains respectively. No ~q-phase particles were found in twin boundaries. Subsequent deformation causes no changes in the microstructures

of the sample heads in states A and B irrespective of the strain rates.

The microstructure of the sample head in both the states and that of the sample neck in state A after superplasticity deformation at optimal con- ditions and a strain level of 160% is shown in Fig. 3(a-c). It can be seen that superplastic treatment results in the making of grain sizes more uniform and disappearance of twins. The comparison of the grain size distribution histograms of arbitrarily selected regions of the sample neck and head indicates that the microstructure homogeneity increases due to the simultaneous growth of "fine" grains and refining of "coarse" grains [Fig. 3(d)]. The distribution of the g2-phase becomes more uniform too. During super- plasticity deformation, grains preserve their equiaxial shape until failure.

At the initial stage deformation (E=15%) mechanical twinning and bulging of grain boundaries occur in the "coarse" grains. With further straining (E = 30%), new grains are generated both in the grain boundaries and within the original grains. The mech- anisms of the generation of these grains are similar to those observed during the deformation of TiA1 with a coarse grained microstructure [20, 21]. Meanwhile, new grains are formed not at the initial stage of deformation, but throughout the superplasticity flow until the failure of samples. This process is also in progress at strain rates that are much lower than an optimal strain rate (Fig. 4). The recrystallized grains do not mechanically retwin.

The growth of strain rates over the optimal value results in an intensive refining of the sample microstructure while their decrease on the contrary causes the growth of part of grains and leads to grain diversity. At the final stage of superplasticity flow, voids are formed in the intermetallic compound, and they are especially numerous at the site of the sample failure.

3.3. Dislocation structure

The nonuniformity of grain sizes and existence of twin boundaries in the original states of the alloy were taken into account while studying the dislo- cation structure since different strain controlling mechanisms can operate in regions possessing differ- ent microstructures.

In "fine" grains (d < 7-8 #m), at different strain levels, individual dislocations prevail in both the states, only some of the grains displaying dislocation walls and nets [Fig. 5(a, b)]. The density of dislo- cations p is from l0 s to 5 x l0 s cm -2. Meanwhile, "coarse" grains (d > 8/~m) in both the states display a more intensive accumulation of dislocations at the initial stage of flow than "fine" grains. The values of p at the 15% strain in states A and B are 109 and 101° cm -2 respectively. By that time, dislocation tangles, nets, walls, subgrains, and mechanical twins can be found in the grains [Fig. 5(c, d)]. Mechanical twinning in state B is more intensive than in state A.

584 IMAYEV et al.: MECHANICAL BEHAVIOUR OF FINE GRAINED TiAI--I

30 (d}

20

CL

8 16 24 32

Fig. 3. Microstructure (a-c) and grain size distribution histograms (d) of TiA1 after 160% deformation at ]025°C and strain rate of 8.3 x ]0 -4 s-=: (a) state A (specimen head), (b) state B (head), (c) state A

(specimen neck); (d) • state A (head), O state A (neck).

Contrary to other boundaries twin boundaries can contain lattice dislocations [Fig. 5(e)].

Fig. 4. Microstructure of TiA1 after 160% deformation at 1025°C and strain rate of 1.6 × 10-4s -I, arrows point at

some of the recrystailized grains.

With the growth of strain level the density of dislocations in "coarse" grained areas decreases, but still remains high enough to induce superplasticity flow and at ~ = 120% it reaches from 109 to 5 × 109 cm -~. The character of the dislocation struc- ture remains the same. No mechanical twinning is observed under these conditions. The growth or decrease of the testing rate do not cause any qualitative changes in the dislocation structure, only p varies. After superplasticity deformation, the ~2- phase particles are located not only in the grain boundaries, but also within the grains [Fig. 5(a)]. It should be noted that a high density of dislocations observed after superplasticity flow is also encoun- tered in the Ni3AI intermetallic compound with boron additives [25].

The process of the relaxation of lattice dislocations in grain boundaries during annealing indicates that in twin boundaries the lattice dislocation contrast does not disappear even after annealing at I100°C [Fig. 6(a)]. Besides, starting with 1000°C on, the

IMAYEV e t al.: MECHANICAL BEHAVIOUR OF FINE GRAINED TiAI--I 585

Fig. 5. Dislocation structure of TiAI after superplasticity deformation: (a, b) "fine" grains, E = 120%; (c, d) "coarse" grains, E = 15%; (e) twin boundary, E = 15%; arrows point at trapped lattice dislocations.

Fig. 6. Relaxation of trapped lattice dislocations in grain boundaries ofTiAl: (a) temperature of annealing 1100°C, twin boundary; 0a) 600°C, random boundary.

586 IMAYEV et a/.: MECHANICAL BEHAVIOUR OF FINE GRAINED TiAI--I

system of lattice dislocations in twin boundaries becomes ordered and quasiperiodic rows and nets are formed. Other boundaries which are probably random are frc¢ of lattice dislocations already at 600°C [Fig. 6(b)].

4. DISCUSSION

Proceeding from the data thereabove, it can be stated that in the TiAI int~rmetallic compound like in metals all the features of structural superplasticity can be realized once a fine grained microstructure is available [15-17]. Similar regularities of phenomenol- ogy and structural changes can be regarded as an indication of the similarity of deformation mechan- isms operating in these materials so much different in their nature.

It should be noted that a specific feature of TiAI is dynamic recrystallization under superplasticity con- ditions, which eliminates the original diversity of grain sizes. The latter fact is quite unexpected since it is well known that in metals the localization of supcrplasticity flow in fine grained areas hinders the development of recrystallization in coarse grains [26]. The easiness of dynamic recrystallization in TiA1 is probably due to a lower mobility of dislocations in it as compared to metals. In fact, a high Peierls barrier and considerable crystallographic dissociation of su- perdislocations a[101], ~t[011] and 1/2a[112] hinder their slip and transition from one plane to another by means of cross-slip and climbing [5-10, 27-29]. As a result, a density of dislocations sufficient for the start of dynamic recrystallization is created in coarse grains while dislocations are observed in many fine grains, which is not typical of the structural super- plasticity of metals [15-17].

Though the uniformity of TiAI microstructure increases by a certain stage of deformation, the formation of new recrystallized grains continues in this intermetallic compound until the sample failure. New grains are probably generated in individual grains grown in the course of superplasticity flow. Similar processes were observed in a titanium base /~-alloy [30]. Dynamic recrystallization is also in progress during the superplasticity deformation of Ni3AI with boron additives [25].

The aforementioned peculiarities of the super- plasticity behaviour of the TiA1 intermetallic com- pound can be explained using a model of grain boundary dislocations motion kinetics [31]. Accord- ing to this model, under supcrplasticity conditions, grain boundaries act as a "channel" through which plastic flow is accomplished via the movement of grain boundary dislocations. When grains grow, the mobility of grain boundary dislocations is insufficient for the strain rate specified by the machine, which results in the enhancement of intragranular slip facil- itating the relaxation of the arising stresses. At the same time, contrary to TiAI, the coarsening of grains during superplasticity flow in metals does not always

result in a considerable accumulation of dislocations because recovery is very fast [15-17].

During the supcrplasticity deformation of TiAI, an approximate dynamic equilibrium between the coarsening and refining of grains is achieved. Any deviation of the temperature-rate parameters of superplasticity deformation breaks this equilibrium and leads to a decrease in the compound plasticity. It should be noted that though normal grain growth in metals during supcrplasticity flow reduces the metal plasticity [15-17], the development of dynamic recrystallization "limiting" the grain growth in TiAI may provide for the enhancement of its plasticity. However, acts of recrystallization do not occur regularly in the microstructure and they probably fail to remove stress concentrations from the grain boundary junctions. The relaxation of stresses via grain boundary sliding, when the accommodation of grains grown during the supcrplasticity flow is insufficient, usually results in cavitation [15-17]. This is what observed in the experiment. The coalescence of voids during deformation leads to the sample failure.

Simultaneous growth of fine grains and refining of coarse grains during superplasticity flow provide for a more uniform microstructure of the intermetallic compound as compared with the original state, which is rather important since, as stated in [23], nonunifor- mity of microstructure is one of the reasons for the low plasticity of TiAI.

Another important observation is a strong effect of the grain boundary structure on the mechanical properties of TiAI under superplasticity conditions. A slight increase in the mean grain size during the transition from state A to state B results in an abrupt drop in the compound deformability. Since states A and B are characterized by almost similar grain sizes and similar amount, shape, and distribution of g2" phase particles, the difference between the intermetal- lic compound properties in these states is probably due to the peculiarities of the grain boundary struc- tures or, to be more precise, to the different number of twin boundaries in each of the states.

The fact that different numbers of twin boundaries produce different effect on the superplastic flow in states A and B can be explained in the following way. The accumulation of dislocations during supcrplastic flow is known to be controlled by the processes of their generation and absorption in grain boundaries [32]. The following expression was obtained for the density of mobile dislocations:

p = ~z/4b(7 (1)

where i is the strain rate, b is the Burgers vector, is the mean grain size, and ¢ is the time required

for dislocations to be absorbed in grain boun- daries. According to (1), the observed difference in the density of dislocations in states A and B can be explained by the fact that time required for the absorption of trapped lattice dislocations is longer in

IMAYEV et al.: MECHANICAL BEHAVIOUR OF FINE GRAINED TiAI--I 587

twin boundaries than in random boundaries. This is proved by the study of the absorption properties of grain boundaries in TiAl: tr of lattice dislocations in twin boundaries is much higher than that in random boundaries.

In the model the kinetics of the motion of grain boundary dislocations [31], the rate of plastic defor- mation under superplasticity conditions can be rep- resented as

m ~GBS m kG/~b~ b (2)

where G is an average Burgers vector, Pb and 6 b are the volume density and average rate of the grain boundary dislocations motion respectively, and k is a coefficient. According to (2), in random boundaries, the rate of grain boundary sliding is higher than in twin boundaries due to the enhanced mobility of grain boundary dislocations [31]. At the steady flow stage the density of grain boundary dislocations in random boundaries tends to be a stationary value while in twin boundaries they are to be accumulated. That is why the development of grain boundary sliding in TiAI is impeded in state B where twin boundaries are more numerous than in state A. As a result, state B is characterized by lower values of the strain rate coefficient m and relative elongation to rupture 6 as compared to state A.

The low mobility of grain boundary dislocations in twin boundaries probably results in a low plas- ticity of cast TiA1 at high temperatures [18]. For in spite of the large area of grain boundaries (d < 5/am), the compound contains no less than 90% of twin type boundaries (transformation twin boundaries).

Thus, the data obtained indicate that for the effect of superplasticity to be displayed in a polycrystal not only a considerable area of grain boundaries, but also their structure are important. A similar conclusion was drawn in [24] where the superplasticity behaviour of a magnesium alloy MA8 containing deformation twins in the microstructure was studied.

The influence of the grain boundary structure on the TiA1 mechanical properties is also manifested under the condition of ductile-brittle transition [22]. This problem will be given a detailed consideration in part 2 of this paper.

5. CONCLUSION

1. The TiA1 intermetallic compound with a fine grained microstructure displays all the features of structural superplasticity in metals at 1025°C and a strain rate of (0.83-1.6)x 10-3s -~, the strain rate sensitivity coefficient m being about 0.33 to 0.43 and the relatively elongation to rupture reaching 200-250%.

2. For the superplasticity effect to be displayed, both the area and structure of grain bound- aries are important. The growth of the fraction of boundaries results in a decrease in the superplastic parameters.

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