Materials Science and Engineering A 424 (2006) 71–76

Low temperature flow behavior of B2 intermetallicphase in Ti–Al–Nb system

Hari Shankar b, N. Eswara Prasad a, A.K. Singh a, T.K. Nandy a,∗a Defence Metallurgical Research Lab., Hyderabad 500058, India

b Indian Institute of Technology, Kanpur 208016, India

Received 5 September 2005; accepted 21 February 2006

Abstract

Low temperature flow behavior of polycrystalline B2 intermetallic in Ti–Al–Nb system has been studied in a range of temperature (77–298 K)and strain rate (10−4 to 10−3 s−1). To assess the rate limiting mechanisms, activation volume and energy have been determined. Relatively low valuesof activation volume (20–50b3) and activation energy (0.2µb3) suggest that a short-range obstacle that is Peierls barrier restricts the dislocationmovement. Work hardening behavior of the intermetallic is also studied and contrasted with that of other B2 intermetallics. The intermetallicexhibits extremely low rates of work hardening that is attributed to localized and planar flow.©

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2006 Elsevier B.V. All rights reserved.

eywords: B2 intermetallic; Ti–Al–Nb system; Peierl

. Introduction

The role of B2 phase in the high temperature titanium alu-inide alloys is well known. This phase, which has an ordered

cc structure (Fig. 1), is stabilized by the addition of Nb. Itnhances the ductility and fracture toughness of Ti3Al–Nb alloys1] because of its ability to deform by multiple slip systems. Thetrengthening effect of B2 phase in multiphase �2 + B2 (Ti3Alase) or O + B2 (Ti2AlNb base) alloys at room temperature haslso been reported [1]. The processing of the alloys is signifi-antly facilitated in the presence of B2 phase. Alloys containing2–24 at.% Al and 20–25 at.% Nb can be easily cold rolled toheets because of the presence of ductile B2 phase and therefore,hey are especially suited to sheet related applications and metal

atrix composites [2,3].Different phases that are present in Ti–Al–Nb systems, in

dditions to B2, are Ti3Al (D019, P63/mmc) and Ti2AlNb (Ohase, Cmcm space group). While considerable attention haseen paid towards the understanding of the mechanical behav-or of Ti3Al [1,4] and Ti2AlNb intermetallics [1,5], the B2hase has attracted limited attention. Since efforts continue to

improved room temperature and high temperature properties,an understanding of the mechanical behavior of the B2 phase, isparticularly crucial. Mechanical properties of a series of com-plex B2 intermetallics have been evaluated by Naka et al. [9–11].Interesting tensile properties were obtained in some of thesealloys. Studies on dislocation substructure of the B2 phase fol-lowing deformation have been carried out by Banerjee et al.[12] and Puissochet [13]. Surprisingly, no detailed study existsas regards the flow mechanisms of the B2 phase at low temper-atures given the influence it has on the properties of both Ti3Aland Ti2AlNb alloys. Therefore, such a study is useful for a betterunderstanding of the mechanical behavior of current generationmulti-phase, �2 + O + B2 alloys. Present study is a preliminaryattempt in this direction.

2. Experimental

As cast alloys were received in form of 12 mm × 12 mmsquare cross-section rods extracted from 120 mm dia. doublemelted ingot. Compositions of the alloys are given in Table 1.

evelop superior compositions of Ti2AlNb alloys [6–8] with

∗ Corresponding author. Tel.: +91 40 2434 2123; fax: +91 40 2434 2123.E-mail address: nandy@dmrl.ernet.in (T.K. Nandy).

The blanks were subjected to a solution heat treatment at 1473 Kfor 45 min followed by water quench. The microstructure in allcases comprised of single-phase equiaxed B2 as shown in Fig. 2.Cylindrical specimens of 6 mm diameter and 9 mm length weremachined from as heat-treated rods and tested in compression

921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2006.02.048

72 H. Shankar et al. / Materials Science and Engineering A 424 (2006) 71–76

Fig. 1. Crystal structure of B2 phase.

Table 1Nominal composition (at.%) of the alloys investigated

Alloy Al Nb Oa Na Ti

Alloy 1 22 25 0.30 0.04 BalanceAlloy 2 24 25 0.33 0.05 BalanceAlloy 3 22 20 0.29 0.04 Balance

a Determined using Leco gas analyzer.

Fig. 2. Optical microstructure of alloy 1 (Ti–22Al–25Nb, 1473 K/45 min/waterquench) showing single-phase microstructure.

using a screw driven Instron 1185 test system. Compression testswere performed in a temperature range of 77–298 K at strainrate of 10−3 s−1. Additionally, strain rate jump tests (10−4 to10−3 s−1) were also carried out. Appropriate lubricants such asgraphite paste were used in order to minimize friction betweenthe specimen and compression platens. For low temperaturetests, a minimum of 20-min soak time was given for the equili-bration of temperature before commencing the tests.

3. Results

A typical set of true stress versus true plastic strain curvesof Ti–22Al–25Nb at three temperatures – 77, 228 and 298 K –are shown in Fig. 3. The curves are characterized by an initialincrease in the flow stress followed by an apparent saturationstress at larger strains. Flow stress curves are shifted up at lowertemperatures that is, for a constant plastic strain, the flow stress

Fig. 3. Stress strain curves of Ti–22Al–25Nb (Alloy 1) in single-phase B2microstructural condition.

increases with decreasing temperature. Flow stress (at 2% plasticstrain) versus temperature plots of all alloys (Fig. 4) show adecreasing trend with increasing temperature and the data pointsoverlap suggesting relatively minor variations in the flow stressvalues from alloy to alloy.

Strain rate sensitivity, S, has been determined from strain ratejump tests using equation:

S = σ1 − σ2

ln(ε1/ε2), (1)

where σ1 and σ2 are the flow stress at two different strain ratesat ε1 and ε2.

Strain rate sensitivity values versus plastic strain plots of allthree alloys are shown in Fig. 5. The sensitivity values overlapwithin a broad band and exhibit a marginally decreasing trend

F

ig. 4. Flow stress of single-phase B2 alloys as a function of temperature.

H. Shankar et al. / Materials Science and Engineering A 424 (2006) 71–76 73

Fig. 5. Strain rate sensitivity (determined from strain rate jump tests) plottedagainst plastic strain.

with increasing strain. At a constant strain, the S values are higherat lower temperature for all three alloys.

4. Discussion

4.1. Rate kinetics

Fig. 6 shows flow stress versus temperature plot of one of theB2 alloys (investigated in the present study) along with thoseof representative bcc (Ta) [14], fcc (Cu) [15] and cph (Ti) [16]metals. FCC copper (that deforms by a forest intersection mech-anism at low temperature) exhibits relatively lower values flowstress (in the given scale) that is almost independent of temper-ature. Flow stress of tantalum and titanium are comparativelyhigher and relatively steeper temperature dependence can beseen in both the cases. Both tantalum and titanium deform by aPeierls mechanism at low temperatures [17–19]. The B2 alloysstudied in the present investigation exhibit significantly higher

Fig. 7. Activation volume of all alloys as a function of strain.

flow stress values and a much greater dependence on tempera-ture, almost comparable to that of another B2 intermetallic NiAlthat has been studied extensively [20].

Activation volumes (V*) have been determined from strainrate sensitivity values using following equation:

V ∗

b3 = MkT

b3

�ln(ε)

�σ∗ , (2)

where b is the Burger’s vector of 1/2 〈1 1 1〉 dislocation [12], Mthe Taylor factor (taken as 3), T the temperature in K, ε the strainrate and σ* is the thermal component of flow stress. Activationvolumes are plotted as a function of plastic strain in Fig. 7. Thevalues range from 20 to 50b3 and are comparable to those oftantalum and molybdenum, which have been shown to deformby Peierls mechanism [17,18]. No alloying effects are apparentsince the activation volumes lie within a broad band at a giventemperature.

The total energy of the barrier to dislocation motion is givenas:

G0 = 1

M

∫ σp

0V ∗ dσ∗ (3)

where G0 is the total activation enthalpy, V* the activa-tion volume, σ* the thermal component of the flow stress(σ* = σtotal − σathermal) and M is the Taylor factor. σp, the upperlsobpToft

ae

Fig. 6. Flow stress vs. temperature plots of metals and B2 intermetallics.

imit of the integral, is the mechanical threshold stress corre-ponding to the peak of the barrier that is the thermal componentf the flow stress at V* = 0. The value of athermal stress haseen assumed to be the flow stress of Ti–26Al–25Nb (in single-hase B2) at 573 K, taken from the work of Puissochet [13].hermal component of the flow stress at 77, 228 and 298 K arebtained by subtracting modulus compensated athermal stressrom total flow stress. The modulus is assumed to be same ashat of Ti2AlNb orthorhombic phase [21].

Fig. 8 shows a plot of V* versus σ* for Ti–22Al–25Nblloy. In view of the limited data points, especially near thextremes, a bilinear behavior has been assumed and the area

74 H. Shankar et al. / Materials Science and Engineering A 424 (2006) 71–76

Fig. 8. Activation volume as a function of thermal stress (σ*) for the low temper-ature deformation of B2 phase (alloy 1). The activation volume at liquid nitrogentemperature (corresponding σ* of 1000 MPa) has been determined from dedi-cated tests at two different strain rates.

of the shaded region has been approximated as total activationenthalpy. Table 2 shows the total activation enthalpy value thusdetermined of the B2 phase along with those of different types ofobstacles [22] and it is obvious that the deformation of Ti–Al–NbB2 intermetallics is characterized by low activation enthalpies.This again suggests the presence of short-range obstacles.

Detailed studies have been undertaken to understand lowtemperature deformation of another B2 intermetallic NiAl thatexhibits similar flow behavior in terms of temperature depen-dence and activation volumes. Compression tests have beencarried out on single crystals of NiAl [23] and activation volumeshave been determined for both 〈1 1 1〉 and 〈1 0 0〉 slip. 〈1 1 1〉 slipthat occurs in the hard orientation is associated with a strongtemperature dependent flow stress and very low activation vol-umes (20b3) thereby suggesting a highly restricted mobility. Theactivation volumes are consistent with Peierls frictional forceson screw dislocations that have been experimentally observedin the low temperature deformation by transmission electronmicroscopy [24,25]. On the other hand, the deformation of softorientation that exhibits 〈1 0 0〉 slip is characterized by lower ratesensitivity that is lower flow stress dependence on temperatureand relatively higher activation volumes (50–70b3) suggestinga relatively broader obstacle profile. While this has been tenta-tively rationalized with Peierls barrier, subsequent work on in

TD

O

S

M

W

P

situ deformation of NiAl has shown that the motion of 〈1 0 0〉dislocations controlled by unpinning from small extrinsic obsta-cles [26].

Steep temperature dependence and low activation volumesobtained in the present B2 alloys are clear pointers towardslimited mobility of dislocations. One of the three possible mech-anisms can be envisaged in this scenario. (1) Peierls forceson dislocations, (2) solute hardening and (3) hardening due todefects such as antisites and vacancies [27]. If Peierls mecha-nism were rate controlling, one would rather expect rectilineardislocation along simple crystallographic directions arising fromhighly anisotropic growth of dislocation loops [17]. While nosuch study was undertaken in the present investigation, detailedsubstructure characterization already exists in the work of Baner-jee et al. [12] and Puissochet [13] on similar B2 alloys. Indeed,immobile linear segments of screw 〈1 1 1〉dislocations have beenreported in the deformation substructure lending credence tothe present hypothesis. It is interesting to note that screw dis-locations have also been observed by Naka et al. [10] in theroom temperature deformation of more complex disordered andordered B2 alloys with widely different compositions.

4.2. Stress–strain curves and work hardening behavior

Stress–strain behavior at all temperatures is characterized bya discernible increase in the flow stress at low plastic strains fol-ltsomslwvsw

Fs

able 2ifferent types of obstacles and their activation enthalpies [22]

bstacle strength G0 Example

trong 2µb3 Dispersions, large or strongprecipitates

edium 0.2–1.0µb3 Forest dislocations, small orweak precipitates

eak <0.2µb3 Lattice resistance, solutionhardening

resent work 0.16µb3 Short-range obstacles

owed by a relatively flat region at larger strains (Fig. 3). At loweremperatures (77 and 228 K), there is a tendency for the flowtress to drop at larger strains, which is clearly seen in the ratef work hardening plots (Fig. 9). For the room temperature defor-ation, the work hardening rate decreases rapidly in the initial

tages and attains vanishingly small values at larger strains. Atower temperatures (77 and 228 K), the plots are shifted to higherork hardening rates at lower strains and the rates attain negativealues at larger strains, which is consistent with mildly droopingtress–strain curves. Flat stress–strain curves indicative of lowork hardening rates, have also been observed by Puissochet

ig. 9. Rate of work hardening of Ti–22Al–20Nb (alloy 3) as a function oftrain.

H. Shankar et al. / Materials Science and Engineering A 424 (2006) 71–76 75

[13] in her study on single-phase B2 (Ti–27Al–21Nb) alloy andby Gogia [28] in �2 + B2 alloys especially in the microstructurescontaining B2 as a major phase.

Low work hardening rate appears to be a typical featurein a class of � titanium alloys that deform by slip [29,30].Detailed studies have been carried out to understand the roomtemperature work hardening behavior of Ti–21S (Ti–15Mo–2.7Nb–3Al–0.2Si) and Timetal LCB (Ti–6.8Mo–4.5Fe–1.5Al)alloys. These alloys exhibit a pronounced yield point (notobserved in the present alloys) followed by a rapid drop in theflow stress values resulting in negative work hardening rates.While the presence of yield point is attributed to the initialscarcity of mobile dislocations [31], low or negative work hard-ening rates are ascribed to planar inhomogeneous flow thatappears to be accentuated by a localized and near adiabaticincrease in temperature. This is further corroborated by trans-mission electron microscopy studies that show the presence ofintense planar slip bands. Deformation substructure of single-phase B2 alloys studied by Banerjee et al. [12] and Puissochet[13] is dominated by the presence of planar slip bands com-prising of 〈1 1 1〉 dislocations. Based on these observations,low work hardening exhibited by Ti–Al–Nb B2 alloys may beattributed to the presence of inhomogeneous slip.

Planar slip bands have also been observed in single-phase� Ti–Mo system [32]. The origin for such a behavior is tracedto the defect structure of � comprising of fine � precipitatestofsPip

cTic

is a common attribute in intermetallics is generally ascribed toreduced propensity for cross slip of dislocations and significantlylower diffusion leading to reduced recovery kinetics [34]. Lowtemperature deformation of NiAl [35] and FeAl [36] is associ-ated with uniform or cellular distribution of dislocations. Planarslip bands extensively observed in disordered � titanium [32]and B2 Ti–Al–Nb alloys [12] are not present in these systems.Thus, in terms of the stress–strain and work hardening behav-ior, Ti–Al–Nb B2 alloys bear a much closer resemblance to �titanium alloys than its other B2 counterparts possibly becauseof the presence of � phase.

5. Conclusions

(1) Low temperature deformation of the B2 phase in Ti–Al–Nballoys is characterized by steep temperature drop and lowactivation volumes suggesting limited mobility of disloca-tions. The proposed operative mechanism is one of Peierl’sforces on screw dislocations.

(2) Stress–strain behavior of the B2 phase is associated withlow or negative work hardening rates, which is attributed toinhomogeneous planar slip.

(3) Variations in Al and Nb content in the B2 phase do not resultin any change in either the deformation mechanism or workhardening behavior.

A

Datufi

R

[

hat are easily shearable. Shearing and subsequent dissolutionf � precipitates result in slip plane softening that leads tourther localization of deformation. Presence of athermal � iningle-phase B2 has also been shown by Banerjee et al. [12] anduissochet [13]. Thus, stress–strain and work hardening behav-

or of the B2 alloys that is attributed to inhomogeneous slip,ossibly originates from the presence of shearable � particles.

Fig. 10 shows room temperature stress–strain curves of poly-rystalline NiAl and FeAl along with that of single-phase B2,i–22Al–25Nb alloy. Both NiAl [33] and FeAl [34] show pos-

tive slopes indicating much higher work hardening rates inomparison to B2 Ti–Al–Nb alloy. Higher work hardening that

Fig. 10. Stress–strain curves of different B2 alloys.

cknowledgements

The authors acknowledge Dr. A.M. Sriramamurty, Director,efence Metallurgical Research Lab., for his encouragement

nd permission to publish this work. The authors are gratefulo Dr. D. Banerjee, CCR&D (AMS), DRDO, New Delhi, forseful discussions. Authors wish to acknowledge DRDO for thenancial support.

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