The orientation and temperature dependence of the 0.2% proof stress of single crystal Ni3 (Al, Ti)

Scripta

METALLURGICA

V o l . 17, pp. Printed in

209-214, 1983 the U.S.A.

Pergamon P r e s s Ltd. All rights reserved

THE ORIENTATION AND TEMPERATURE DEPENDENCE OF THE 0.2% PROOF STP~SS OF SINGLE CRYSTAL Ni3(AI,Ti)

A.E. Staton-Bevan Department of Metallurgy & Materials Science, Imperial College, London SW7 2BP, U.K.

( R e c e i v e d November 18,

1982)

Introduction The intermetallic compound Ni3(AI,Ti) , the y' phase, has an ordered L12 structure. It is of interest because it occurs as a precipitate in many Ni-based superalloys and is responsible for their high temperature/high strength properties. The mechanical properties of y' have several unusual features; in particular the flow stress increases with increasing temperature to a peak value (1-3) with primary slip occurring on cube planes at temperatures above the peak (2-3). There is now much evidence to suggest that many of the unusual mechanical properties of y' result either directly or indirectly from the low cube plane antiphase boundary (APB) energy of the LI 2 crystal structure (4). The resulting tendency for screw dislocation segments on octohedral {Iii} planes to cross-slip on to cube planes forms the basis of the most widely accepted theories of the deformation behaviour of 7' With these points in mind, crystal compression axes for the present study were chosen with special reference to the Schmid factors for primary slip and cross-slip on octohedral and cube planes. In this paper, measurements of the 0.2% [,roof stress perature, orientation and composition. The proof stress dependent, with double peaks occurring in the C.R.S.S. vs orientated specimens. The results are compared with the cation mobility theory (5). Experimental

are presented as a function of temis shown to be strongly orientationtemperature curves for the <]II> predictions of the Kear-Obl~k dislo-

Procedure

Single crystal compression specimens of Ni~(AI,Ti), 6.0 x 2.5 x 2.5 mm 3, were prepared and tested as described elsewhere (I). A strafn rate of approximately 3 x 10 -4 s -I was employed and test temperatures were in the range -llO°C to iOOO°C. The specimen compression axes and compositions are listed in Table i. Results i.

The temperature dependence of the 0.2% proof stress. The 0.2% proof stress is plotted as a function of temperature for the three crystal orientations <1123> and in Figs. i, 2 and 3 respectively. In order to investigate the effect of deviations from stoichiometry, three compositions were tested for each orientation as listed in Table I. By comparing Figs. i, 2 and 3 it may be seen that there are marked differences in the shapes of the proof stress vs temperature curves for the three orientations. The <123> orientation (Fig. i) shows single, broad peaks and the orientation (Fig. 3) also shows single peaks but they are higher and occur at much higher temperatures. The most interesting curves however are those for the orientation (Fig. 2) which show 2 peaks for all three compositions tested. 2.

Deviations from stoichiometry. For the <123> and orientations (Figs. i and 3) there is no significant difference between the ]proof stress values for the Ni-rich and stoichiometric compositions. Making the com-

209 0036-9748/83/020209-06503.00/0 Copyright (c) 1 9 8 3 P e r g a m o n Press

Ltd.

210

0.29

PROOF

STRESS

OF

Vol.

Ni3(AI,Ti)

17,

TABLE I SPECIMEN COMPRESSION Compression Axis

Composition

AXES ANO COMPOSITIONS at

Co de

Ni

A1

< 123>

75.7

19.2

5.1

< 123>

75.1

16.8

6.I

$1

< 123>

74.4

20.3

5.2

A1

N1

< Iii>

75.5

19.3

5.2

N2

< iii>

75.1

18.8

6.1

$2

< 111>

74.3

20.4

5.3

A2

N3

< 001>

76.0

18.7

5.3

< 001>

75.1

18.8

6.1

$3

< 001>

74.4

20.4

5.2

A3

*

N, S and A denote Ni-rich,

t

*

Ti

stoichiometric

and Al-rich

respectively.

(123) AXIS

02%

xl02

PROOF STRESS MNIma

2 0 "200

;

I

6;0

t

ooo

TEMPERATURE "C

FIG. 1 The temperature-dependence of the 0.2% proof stress of single-crystal Ni3(AI,Ti). <123> compression axis (Compositions • = AI, • = SI, • = NI, see TabIe I, curve N1 previously published (i).)

No.

2

Vol.

17,

No.

10

T 0.2%

2

0.2%

xl0 z

PROOF

STRESS

OF

Ni3(AI,Ti)

211

<111)AXIS

8

PROOF STRESS 6 MNlm2

FIG. 2 The temperature-dependence of the 0.2% proof stress of single-crystal Ni3(AI,Ti) compression axis. (Compositions A= A 2 , • = S2, o = N2, see Table I)

2 0

-

-200

0

'

'

'

'

200

t.O0

600

800

TEMPERATURE

)(10~ ',

T

I

1000

C

<001)A X I S ~

0 2% PRO0F STRESS MNImz

FIG. 3

-200

I

I

ZOO

~00

n

6OO

TEMPERATURE

BOO ~

0

tOO0

The temperature-dependence of the 0.2% proof stress of single-crystal Ni3(AI,Ti ) <001> compression axis. (Compositions A = A3, • = $3, • = N3, see Table i. Curve N3 previously publlshed(~. Dashed lines show the orientation curves of Fig. 2 for comparison.

212

0.2% PROOF STRESS OF Ni3(AI,Ti)

Vol.

17, No. 2

position Al-rich however has a strengthening effect and increases the Critical Resolved Shear Stress (C.R.S.S.) by approximately 75MN/m 2 for the <123> and 50 MN/m 2 for the <001> orientations. For the orientation (Fig. 2) an increase in strength for the Al-rich composition is also observed but, in addition, the magnitudes and temperatures of the 'double peaks' referred to in I. above are dependent on composition. Referring to Fig. 2, it may be seen that on going from composition N2 + $2 ~ A2, the double peaks and, in particular, the lower temperature peaks, become more prominent and occur at lower temperatures. Discussion Deformation mechanisms responsible for the temperature-dependence of the 0.2% proof stress. (a) The <123> orientation. The special features of the slip geometry of specimens having a <123> orientation are given in Table II. This orientation was chosen because it has single primary octohedral or cube slip systems having similar Schmid factors. Under these conditions the proof stress vs temperature curves (Fig. I) show single, broad peaks similar to those obtained for polycrystalline specimens (2). As has been shown in a previously published TEM study of specimens NI (6), primary slip occurs on octohedral planes below the peak temperature and on cube planes above the peak. The anomalous increase in the octohedral proof stress with increasing temperature can be explained most satisfactorily by the dislocation-mobility theory of Kear and Oblak (5). Superlattice dislocations in y' mainly consist of pairs of ~ superlattice partials, separated by antiphase boundary (APB). The low cube plane APB 2 energy of the LI 2 structure results in screw dislocation segments on {III} cross-slipping on to cube planes and in some cases becoming sessile to form so-called "Kear-Wilsdorf locks" (7). Measurements of the separation of crossslipped ~ superpartials in Ni3AI by Oblak and Rand (8), indicate that there is a lattice friction2force opposing slip on cube planes which decreases with increasing temperature. This implies that cube cross-slip becomes easier and double cross-slip back on to {IIi} planes becomes mere difficult with increasing temperature. According to Kear and Oblak (5), the resulting decrease in the average slip distance on {iii} accounts for the anomalous increase in the proof stress. The Kear-Oblak theory is supported by the HVEM observations of Nemoto et al (9),who reported that dislocations in y' nucleate suddenly in localised, dense bands and then move either in a slow, jerky manner at low temperatures, or quickly at higher temperatures. If the theory is correct, factors which make {III} + {iOO} slip easier, or {III} sllp more difficult, may be expected to increase the proof stress. The orientation should therefore have a higher C.R.S.S. for octohedral slip than the <123> orientation since it has a much lower Schmid factor (S.F.) for octohecral slip (0.272 as opposed to 0.467) and a slightly higher S.F. for cube cross-slip (0.472 as opposed to 0.404, see Table II). TABLE II The Characteristics of the Slip Geometry for <123>, <211> and Orientated Compression Specimens. PRIMARY

SPECIAL FEATURES

Single

CUBE

SLIP

SLIP

NO. o f

systems

1

primary octohedral or cube slip s y s -

tems.

CROSS-SLIP

~IM~Y

OCTOHEORAL

-

factor

factor

I i i ~ I00 Favoured Schmid factor

III

Fsvoured

~ III

Schmid factor

primmry I00 + II] Favoured

5chmld factor

tBms

0.487

i

High

0.455

Yes

0.404

No

0.000

0.350

Yes

0.472

NO

0.272

0.272

(or O)

(or O]

0.406

0.406

High

hay-

stmiler Schndd factors



Octohedrel s l l p not favoured

6

Cube s l i p

B

0.272

3

0.406

suppressed High

0.472

High

Low

0

0.000

NO

0.000

Yes

Vol.

17, No. 2

0.2% PROOF STRESS OF Ni3(AI,Ti)

213

(b)

The orientation. The temperature-dependence of the 0.2% proof stress for specimens having a <111> compression axis is shown in Fig. 2. For all three compositions two peaks occur in the proof stress vs temperature curves, a feature not previously observed for L12 structures. The high proof stress for octohedral slip, predicted by the Kear-Oblak theory for this orientation is indeed observed and appears to be responsible for the lower temperature peaks. Earlier Yield Stress measurements on binary Ni3AI (88.2 wt % Ni) (IO) did not detect an orientation-dependence of the C.R.S.S. at room temperature, 250°C or 450°C. This was probably due to the gap of 2OO°C between measurements and the reduced effect for Ni-rich compositions (see Fig. 2). It is interesting to note that a S.F. dependence of the C.R.S.S. has been reported for the highly-ordered L12 compounds Ni3Ga (ii) and Ni3Ge (12), but that no strong S.F. dependence is observed for ordered Cu3Au which disorders at 664K (13). All three compounds, like Ni3(AI,Ti) , show an anomalous increase in the flow stress with increasing temperature, but the S.F. evidence suggests that two different mechanisms are responsible. In the case of Cu3Au , Pope has proposed that the main cause of the anomalous increase in the flow stress is a modulus interaction between dislocations and local regions of disorder within the ordered matrix, which increases with increasing temperature (13). For the more highly ordered compounds Ni3Ga , Ni3Ge and Ni3(AI,Ti) on the other hand, the S.F. evidence indicates that the Kear-Oblak cross-slip mechanism (5) is responsible for the increase in flow stress. Returning to Fig. 2, the lower temperature peak in the proof stress for the <111> orientation may be explained in terms of the Kear-Oblak theory by an increase in the number of crossslipped screw dislocations with increasing temperature. At low temperatures these segments form sessile "Kear-Wilsdorf locks" (7). At higher temperatures however, TEM has shown that the segments become mobile on the cube cross-slip plane (6), thus explaining the fall Ln the proof stress above the lower temperature peak. It is interesting to note that the lower peak becomes more pronounced and moves to a lower temperature on moving from the Ni-rich to the Al-rich composition, suggesting a possible composition-dependence of the {iOO} ! APB energy. 2 Critical resolved shear stress calculations (e.g. Fig. 4) suggest that the change in the slip mode from low temperature primary octohedral slip to high temperature primary cube slip, In Fig. 4, calculations of the C.R.S.S. are occurs at temperatures between the two peaks. shown for orientated stoichiometric crystals, assuming octohedral slip (solid squares) or cube slip (open squares). The solid line shows an estimate of the octohedral C.R.S.S. xlO z obtained from the cube-orientated specimens in which cube-slip was suppressed although values may be lower than those for specimens because of suppression of the Kear3 ! (111) A X I S Wilsdorf mechanism. For the orientation, C.R.S.S. the calculated values of the octohedral C.R.S.S. (solid squares) fall well below the solid line MNImZ at temperatures above 3OO°C and it may therefore be assumed that the primary cube slip o D mode is operative, thus the true C.R.S.S. is 2 that shown by the open squares. The change in slip mode for the orientation there-

t

i

c.R.s.s.(.O--//~

•

FIG 4. The temperature-dependence of the critical re-solved shear stress (C.R.S.S.) for orientated, stoichiometric Ni3(AI,Ti) crystals (composition $2), calculated assuming octohedral or cube slip. ( • C.R.S.S. assuming octohedral slip; D C.R.S.S. assuming cube slip.) The solid curve shows the octohedral C.R.S.S. calculated for cube-orientated specimens, composition $3.

•

[]

•

-200

|

in

D

i

i

i

i

i

0

200

4.00

600

800

T E M P E R A T U RE

"C

1000

214

0.2%

PROOF

STRESS

OF N i 3 ( A I , T i )

Vol.

17,

No.

2

fore occurs at lower temperatures than for the <123> orientation as predicted by the Kear-Oblak theory. Fig. 2 also suggests that the change-over occurs at lower temperatures for the Al-rieh than for the Ni-rlch compositions. The higher temperature peak for the orientation appears to occur when the cube slip mode is operative and can be explained in terms of dislocation widths. It is known that the separation of superlattice dislocation pairs on cube planes increases with increasing temperature (8). This would ~ k e the cross-slip or climb required for dislocation multiplication increasingly difficult. The effect on cross-slip would be accentuated in the orientation by the low S.F. for {IOO} + {III} cross-slip (see Table II). (c) The orientation The temperature-dependence of the 0.2% proof stress for specimens having a compression axis is shown in Fig. 3. The proof stress curves for the orientation (dashed lines) have been included for comparison. The important feature of the orientation is the absence of a resolved shear stress on any cube system (Table II). This must suppress dislocation motion on both the primary and the cross-slip cube planes. The effect of suppressing the high temperature primary cube slip mode may be seen in Fig. 3. The proof stress continues to increase with increasing temperatures up to very high temperatures, as high as 950°C for the Ni-rich composition. As with the <123> and orientations, the anomalous increase in the proof stress for tile orientation may be explained in terms of the Kear-Oblak mechanism. Although the movement of dislocations on cube planes is suppressed by the specimen geometry, screw segments on {iii} may still be expected to cross-slip on to cube planes because of the low cube plane APB energy. However, for this orientation, plastic strain requires the movement of dislocations on {IIi} at all temperatures and therefore double cross-slip back from {iOO} to {iii} is neces~ sary. This becomes increasingly difficult with increasing temperature because of the greater separation of the cross-slipped ~ superlattice pairs, and accounts for the anomalous increase in the proof stress. 2 The effect of suppressing cross-slip on to cube planes may also be seen in Fig. 3. The lower temperature peak observed for <119 orientated crystals and attributed to Kear-Wilsdorf locking, is absent from the
A.E. Staton-Bevan and R.D. Rawlings, Phys. Star. Sol. (a) 29, 613 (1975). R.D. Rawlings and A.E. Staton-Bevan, J. Mat. Sci. I0, 505 (1975). P.H. Thornton, R.G. Davies and T.L. Johnson, Met. Trans. i, 212 (1970). P.A. Flinn, Trans. M.S. AIME, 218, 145 (1960). B.H. }~ear and J.M. Oblak, J. de Physique 35, C7-35 (1974) A.E. Staton-Bevan and R.D. Rawlings, Phil. Mag. 32, 787 (1975). B.H. Kear and H.G.F. Wilsdorf, Trans. AIME 224, 382 (1962). J.M. Oblak and W.H. Rand, Proc. 32nd Annual EMSA Meeting, 502 (1974). M. Nemoto, J. Echigoya and H. Suto, Proc. 5th Int. Conf. on HVEM Microscopy, Kyoto, 467 (1977). S.M. Copley, B.H. Kear, Trans. AIME 239, 977 (1967). S. Takeuchi and E. Kuramoto, Acta Met. 21, 415 (1973). H.R. Pak, T. Saburi and S. Nenno, Proc. 2nd Kyoto Symp. on the Mechanical Behaviour of Materials, Vol. II, (Jap. Soc. Mat. Sci.), 23 (1974). E. Kuramoto and D.P. Pope, Phil. Mag. 33, 675 (1976).