S c r i p t a HETALLURGICA Vol . 26, pp . 8 4 1 - 8 4 6 , 1992 Pergamon P r e s s p l c e t MATERIALIA P r i n t e d in t he U .S .A . A l l r i g h t s r e s e r v e d

JERKY FLOW (THE PLC EFFECT) IN L12 AI3Ti-BASED ALLOYS

L. Potez*, G. Lapasset* and L.P. Kubin ** *ONERA (OM), 29 Av. de la division Leclerc, BP 72, 92322 Chatillon Cedex, France

LEM, CNRS-ONERA, 29 Av. de la division Leclerc, BP 72, 92322 Chat'ilion Cedex, France ( R e c e i v e d December 10, 1991)

( R e v i s e d J a n u a r y 8, 1992)

Introduction

Al3Ti-based alloys are of great interest for aerospace applications because of their low density, good oxidation and corrosion resistance. Unfortunately, AI3Ti has low symmetry, tetragonai, 12022 structure and is extremely brittle at room temperature (1). The 13022 structure is related to the cubic L12 via a displacement vector of 1/2 [110] on every (001) plane and several AI-Ti-X systems have been successfully explored in order to obtain L12 compounds with a stcehiometry close to AI3Ti (2). As the LI 2 structure exhibits five independent slip systems, an increase in the ductili- ty of such alloys can be expected. For this reason, the compressive strength of L12 AI3.xM x Ti compounds has been extensively studied for M = Cu (3), Cr, Mn (4, 5, 6, 7), Ni (8), Fe (9,10) and, more recently, Pd (11) and Ag (12). It has been shown that, despite a good ductility in compression at room temperature, these alloys remain extremely brittle in tension, whatever the ternary element.

The aim of this study is to report on observations of jerky flow and anomalies of mechanical properties in A13.xM x Ti (M = Cu, Cr, Mn ) compounds in the temperature range 0.3-0.4 T m (T m, melting temperature). This behaviour is identified as resulting from the Portevin-Le Chatelier (PLC) effect.

The alloys used in this investigation were prepared from the pure elements AI (99,99%), Ti (99,35%), Cu (99,99%), Cr (99,65%), M~ (99,65%) by are-melting in a water cooled copper hearth. Prior to melting, the chamber was twice eva- cuated to 10"J Pa and baekfilled with high purity argon. The specimens were remelted a minimum of six times and the weight of the specimens was measured before and after melting to check the composition accuracy. Subsequently, the alloys were homogenized during 24h at 1000*C under argon atmosphere. As the homogenized material was quite po- rous, it was further submitted to hot isostatic pressure (HIP) at 1155"C during 3h and under a pressure of 100 MPa.Tbe target compositions were in all cases close to AIsMTi 2 and the results of the chemical analyses are listed in Table 1.

TABLE 1

Analysed compositions of the L12 compounds

alloys AI (at%) Ti (at%) M (at%)

AI5Ti2Cu 62.5 24.5 13

AI5Ti2Cr 59.4 28.1 12.5

AI5Ti2Mn 62 25.5 12.5

In order to analyse the effect of oxygen, two alloys were elaborated simultaneously with the same composition (AI62Ti25Cu 13) but with two titaniums of different impurity contents The oxygen content was analysed for both al- loys by oxido-reduction after the hiping process. The alloy elaborated with a commercial titanium contained 485 ppm of oxygen and the other one, elaborated with a high purity titanium, contained 210 ppm of oxygen.

The compression specimens were spark-cut from the hiped ingots and the surface damage was removed by mechanical polishing. The dimensions of the paraUelepipedic specimens were 3"3"6 mm'~'. Compression tests were carried out on a

841 0056-9748/92 $5 .00 + .00

C o p y r i g h t (c) 1992 Pergamon P r e s s p l c

842 JERKY FLOW IN A13Ti Vol . 26, No. 5

servomechanical ADAMEL deformation machine under vacuum (5"10 -4 Pa), at a strain rate of 10 -4 s "1 and for several temperatures between 20 and 800°C. All the tests were interrupted at a strain of 3%.

Strain rate change tests were carried out with a strain rate ratio (82/81) = 3. The upper strain rate was maintained in the strain interval 1%-2%. The stress changes after an upward strain rate jump, AO = (O:2---~1) were measured by extrapo- lating the two linear portions of the deformation curves in steady state, thus avoidin~ th~ transient zone. The quantity extracted from the experiments is the strain rate sensitivity S = Ao/Ln(E2/E1).

The microstructures were investigated by transmission electron microscopy (TEM), x-ray diffraction and scanning electron microscopy in the backscattered electron mode

Results

For the compounds A162 5Ti24 5Cu13, AI62Ti25 5Mn12 5, A159 4Ti28 1Crl2 5' the as-cast microstructures are den- dritic with a small amoui~t ofs~cond phases in th~ interd~ndritic i~gion~. This'is illustrated by fig. 1-a in the case of the AI62Ti25.5Mn 12.5 as-cast compound. These second phases contain a high amount of ternary element and are pro- bably the result of segregation during solidification. As shown in fig. l-b, they are removed by the homogeneizing treatment. Figure l-c, shows that the porosity formed during the homogenization treatment is significantly reduced after hiping. The presence of a single L12 phase in the final state was checked by TEM and x-ray diffraction. More detail about this microstructural study will be published elsewhere.

The 0.2% compressive proof stress of the three intermetallics is plotted in fig. 2 as a function of temperature. All three curves exhibit a stress anomaly in the range 300-500 °C and there is a portion of positive temperature dependence for the alloys containing Cu and Cr.

These strength profiles do not differ strongly from those previously reported for these compounds (4, 5). By contrast, the stress levels are significantly lower than those reported by Mabuchi et al. (6, 7) on A166Ti25Mn 9 and A167Ti25Cr 8 and by Winnicka et Varin (3) on AI5Ti2Cu. These differences can be ascribed to the HIP process used in the present work, which is known to decrease the strength (5) with respect to the powder processing techniques used by Mabuchi et al. (6). They can also stem from differences in titanium contents in the various L12 alloys (5, 13). Fi- nally, the microstructure may significantly differ from one alloy to the other and the importance of a detailed charac- terization by TEM has already been emphasized (14).

a b c

Fig 1 : SEM micrographs in back scattered electron mode of AI62Ti25.5Mn12.5 : (a) as-cast Co) homogenized (c) ho- mogenized and hiped.

The most striking feature of the deformation curves is the occurrence of serrations at 300°C for the three compounds, as illustrated by fig,3, These serrations also occur at 400°C, at the upper strain rate after the strain rate change, Two other features can be norlced on fig. 3 : - the flow s~ss decreases when the strain rate is increased, i.e., the strain rate sensitivity of the flow stress is negative. - There are critical strains for the onset and the disappearance of serrations, whose values depend on temperature.

V o l . 26, No. 5 JERKY FLOW IN A13Ti 843

The occurrence of jerky flow in such types of compounds has already bccn reported but its interpretation is, as ycL still under discussion (12, 15, 16). Since jerky flow is often associated with dynamic strain ageing and the PLC effect, the strain rate sensitivity (SRS) of our specimens, as deduced from the swain rate change tests, has been investigated. Its temperature dependence is shown on fig. 4. The SRS becomes negative in the range 250-500°C with a marked mi- nimum at 300°C. In fig. 5 arc plotted the temperature dependences of the flow stress and of the SRS for the alloy containing Cu. The maximum of the flow stress occurs at a temperature larger than that of the minimum of the SRS. As discussed in the next Section, all these features arc largely sufficicnt to idcmify this jerky bchaviour with the PLC effect.

Discussion

Dynamic swain aging (DSA) and its subdomaln, the PLC effect, arc related to the dynamic interaction of solute atoms and mobile dislocations. A simple way to describe DSA is to consider the evolution of the flow stress with the waiting time of dislocations arrested at "forest obstacles" (17, 18, 19). This flow stress is the sum of two contributions : one, which is thermally activated, corresponds to the glide resistance and it decreases with increasing waiting time. The other is due to the diffusion of solute species on arrested dislocations. Dynamic ageing takes place during the waiting time and its diffusion kinetics arc governed by a characteristic time x related to solute diffusion. For a given tempera- ture and solute content, the hardening due to ageing increases with the waiting time up to a saturation value. Thus, this last contribution yields a component of the flow stress, which increases with temperature and waiting time. Since the waiting time is inversely proportional to the swain rate, it also yields a negative contribution to the SRS. The latter may, thus, turn negative within a range of strain rates and temperatures. Current models have consistently shown that a negative total SRS cannot be associated with stable, uniform plastic flow, which explains the occurrence of serra- tions.

Three other features reported above arc characteristic of the PLC effect, i) - Contrarily to what is stated by Lerf and Morris (16), the occurrence of critical strains for the onset of serrations as well as the fact that, in some cases, jerky flow starts at the onset of plastic flow, arc quite consistent v, dth the PLC effect. The conditions for the onset of serrated yielding associated with DSA has recently been rationalized by Kubin and Estrin (20) in terms of the evolution of the mobile and immobile dislocation densities with swain, ii) - The temperature shift between the maximum flow stress and the minimum of the SRS (cf. fig. 5) is also consistent with the present explanation. The SRS becomes negative when the characteristic diffusion time is approximatively equal to the waiting time ('t -- tw), while the strengthening is maximum and saturates when % << t w, i.e., when the dislocations arc saturated by solute atoms at the end of their wait- hag time. This effect is currently observed in metallic alloys exhibiting DSA (21). iii) - Finally, when z >> t w , the diffusivity of solute atoms is too small to allow for an effective pinning of the mobile dislocations during their waiting time. This explains, in particular, why the PLC effect cannot occur at low temperatures.

Both Mabuchi ctal. (12) and Kumar and Herring (15) have already pointed out that jerky flow, as observed in such in- termetallic compounds, bears similarities with the PLC effect. Recently, Lcrf and Morris (16), disregarding DSA me- chanisms, attempted to propose another explanation. They attributed the occurrence of serrations at intermediate tem- peratures to the reocated dissociations of <110> supcrdislocations into pairs of mobile superpartials which have a higher mobility. This interpretation induces, therefore, strain softening rather than swain-rate softening. It is doubtful that it can explain the related occurrence of a strain anomaly, which is interpreted here in terms of ageing, and of a ne- gative strain rate dependence of the flow stress. Actually, the occurrence of dissociation in this range of temperature could as well bc regarded as a consequence of the high local stresses and strain rates involved during serrated yielding. To date, the occurrence of strain rate softening through dynamic strain ageing is the only mechanism known to bc as- sociated with a negative SRS, a stress anomaly and the occurrence of jerky flow. Thus, the observations of figs. 2 to 6 arc sufficient to identify unambiguously the strength anomalies exhibited by our alloys with the PLC effect.

Since the diffusion properties of solutes atoms arc Ix)orly understood in ordered alloys, a direct identification of the nature of the solute species is impossiblc.Nonethe!ess, some observations may help to clarify the experimental situa- tion. Oxygen is the major interstitial impurity in titanium and it is known to bc responsible for the PLC effect in Ti al- loys in the same range of temperatures (22). The effect of other interstitials like C and N can bc ruled out because of their extremely low solubility in the matrix (23). To check the influence of the oxygen content, the SRS has been mea- sured on two A162Ti25Cu13 alloys containing different levels of oxygen. The curves of fig. 6 show that the behaviour of the SRS is almost the same is both cases, while the two alloys exhibit serrated yielding at 300°C. This does not rule out oxygen, however, for the following reason. The saturated concentration of interstitials in the L12 matrix can be ra- rather low and it could have well reached its saturation value in the two alloys investigated. The possible effect of substitutional elements has also been considered (24). In the absence of diffusion data on such elements, only the com- parison between various alloys of different compositions will help identify the migrating solute, Finally, it must be

844 JERKY FLOW IN A13Ti Vol . 26, No. 5

kept in mind that the mechanisms of jerky flow in ordered alloys arc not necessarily the same as in disordered metallic alloys, although they still involve a diffusional mechanism. For instance, a simple redistribution of solutes within the distorted regions of the dislocation cores has been quoted as a possible cause of DSA (25).

Conclusion

The L12 alloys stabilised with Cr, Mn, Cu present a slight positive dependence of the yield stress around 500°C, ser- rated yielding and a SRS dip at 300°(2. These features lead to the identification of the PLC effect as the cause of jer- ky flow

Acknowl¢~igement

The authors are grateful to Dr S. Naka for fruitful discussions. This work was supported by DRET grant N°89.34.001

References I. M. Yamaguchi, Y. Umakoshi and T. Yamane, Phil. Mag., 55 A, 103 (1987). 2. K.S. Kumar, Int. Mat. Rev., 356, 293 (1990). 3, M.B, Winnicka and R.A Vadn Scripta Metall., 23, 1199 (1989). 4. S.A Brown, K.S. Kumar and J.D. Whittenberger, Scripta Metall,, 24, 2001 (1990). 5. S. Zhang, J.P, Nic, W.W, MiIligan and D.E, Mikkola, Scripta Metall., 24, 1441 (1990). 6. H, Mabuchi, K, Hirukawa and Y, Nakayama, Scdpta Metall., 23, 1761 (1989). 7. H, Mabuchi, K. Hirukawa, H. Tsuda and Y. Nakayama, Scripta MetaIl., 24, 505 (1990). 8. C.D, Turner, W.O. Powers and J.A. Weft, Acta MetaII., 37, 2635 (1989). 9. K,S, Kumat and J.R, Pickens, Scripta Metall., 22, I015 (1988). 10. Z.L. Wu, D.P, Pope and V. Vitek, Scripta Metall., 24, 2191 (1990). II, W.O. Powers and J,A. Wen, Metall. Trans. A, 21,145 0990). 12. H. Mabuchi, K. Hirukawa, K. Katayama, H. Tsuda and Y. Nakayama, Scripta Metall., 24, 1553 (1990). 13. M.B, Winnicka and R.A. Vafin, $cripta Metall., 25, 2297 (1991). 14. L. Potez, A, Loiseau, S. Naka and G, Lapas~t, to be published in J. of Materials Research 15. K,S. Kumar and R.A. Herring, J, Met, 1990 TMS Annual Meeting (Anaheim, CA), 41, 92 (1989). 16. R. Leffand D.G. Morris, Acta Metall., 39, 2419 (1991). 17. P.G, McCormick, Acta Metall., 20, 35 (1972). 18. A. van den Beukel, Phys. Stat. Sol. (a), 30, 197 0975). 19. U.F. Kocks in Progress in Materials science, Pergamon Press (Oxford), 19, 185 (1981). 20. L.P. Kubin and Y.Estrin, Acta Metall., 38, 697 (1990). 21. E. Bouchaud, L. Kubin andH. Octor, Metall. Trans. A, 22, 1021 (1991). 22. Ch. Perrier, Doctorate Thesis, Universitd Pierre ct Made Curie, Paris (1989). 23. G. Hug, Doctorate Thesis, Universitd Pads-Sud, Centre d'Orsay (1988). 24. H. Mabuchi and Y. Nakayama, Bulletin Jap. Inst. of Metals, 30, 24 (1991). 25. R.B. Schwatz and L.L. Funk, Acta Mctali, 31,299 (1983).

Vol. 26, No. 5 JERKY FLOW IN AI3Ti 845

300"

260"

220:

~ 180-

140

100 • . . , . . . . . . . . . . . , . . •

0 200 400 600 800 1000

Temperature (°C)

0.

v

¢O

3,

2 ,

1

0

-1

-2

-3 • ' ' I ' ' ' I ' • ' I ' ' • I • ' ,

200 400 600 800 1000

T e m p e r a t u r e ( ° C )

Fig 2 : Temperature dependence of the proof stress (00.2%) for the LI2-AI3Ti modified alloys.

Fig 4 : Temperature dependence of the stxain rate sensitivity, S, for the L12-A13Ti modified alloys.

300

260

," , 220 ¢N

1 8 0

1 4 0

100

\ .5 1.

" -. I: -,I

" 26d ' i6d 'g6d's6d "1'00() 0 200 400 600 800 1000 Temperature (°C) Temperature (°C)

Fig 5 : Temperature dependences of the proof suess Fig 6 :Dependence of the strain rate sensitivity, S, (o0.2%) and of the strain rate sensitivity, S in the on the oxygen level for the AI62.sTi24.SCUl 3 al- A162.sTi24.sCu 13 alloy, loy.

846 JERKY FLOW IN A13Ti Vol . 26, No. 5

(a) rl

Y- r I ~1

C

/ . / - e = o . 2 % °

I I T = 4 0 0 C

Fig 3 : Stress-strain curves showing serrations for the A159.4Ti28.1Cr12.5 (A), AI62Ti25.5Mn12.5(B), AI62.5Ti24.5Cu13 (C) alloys. The arrows indicate the onset or disappearance of jerky flow. The strain

rates ~land ~2 are respectively 10 -4 s "1 and 2.10 -3 s "1. (a) T = 300°C and (b) T = 400°C