Creep behavior of intermetallic FeAl and FeAlCr alloys

Creep behavior of intermetallic FeAl and FeAlCr alloys

MATERIALS SCIENCE & ENGINEERING ELSEVIER Materials Science and Engineering A220 (1996) 93-99 A Creep behavior of intermetallic Fe-A1 and F e - A 1 ...

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MATERIALS SCIENCE & ENGINEERING ELSEVIER

Materials Science and Engineering A220 (1996) 93-99

A

Creep behavior of intermetallic Fe-A1 and F e - A 1 - C r alloys J.A. Jimenez*, G. Frommeyer Max Planctc Institut fiir Eisenforschung GmbH, Max Planclc Strasse t, D-40237 Diisseldorf, Germany

Received 1 April t996

Abstract A study of creep behavior of several alloys based in the Fe-A1 and F e - A I - C r systems ranging in aluminium from 21.7 to 48 at.% was undertaken. The alloys were produced by induction melting and possess a coarse and equiaxial microstructure, with a grain size of about 500 gin. Compression tests at rates from 10 - 5 to 10-2 s - 1 were conducted at temperatures ranging from 700 to 1000°C under a protective atmosphere of argon to minimize oxidation. Analysis of the stress-strain data revealed in the binary alloys a stress component of about 3 that suggest that creep is controlled by viscous glide of dislocations. For an aluminium content above 30 at.%, the activation energy for creep does not vary very much with the aluminimn concentration and values ranging from 360 to 395 kJ tool- 1 were obtained. On the other hand, the ternary alloys present an improvement in strength in the high temperature compressive creep. A stress exponent of 4-5 is observed in this case that suggest that creep is controlled by dislocation climb. An activation energy for creep of about 505 kJ m o l - 1 was deduced for the alloy containing 30 at.% A1 and 10 at.% Cr. Keywords: Activation energy; AIuminium; Creep behavior

1. Introduction Ordered intermetatlic i r o n - a l u m i n i u m alloys are currently being investigated for use as structural material for high-temperature. The aluminides of iron possess an attractive combination of high melting point and a good oxidation resistance at temperatures above I000°C since alumininm can form a protective oxide layer. Besides, the presence of aluminium confers lower densities than the conventional alloy used at elevated temperatures. The main disadvantage of these alloys, as with m a n y intermetallic alloys, is the low ductility at room temperature. Plastic deformation is more difficult in intermetallic compounds because the ordered structure has associated a strong b o n d and a lower symmetry of deformation. The iron rich end of the FeA1 system shows the presence of disordered ~.-Fe and the cubic crystal ordered structures B2 and D O 3 with stoichiometric compositions FeA1 and Fe3A1 , respectively, [1,2]. F o r an aluminium content ranging from 35 to 50 at.%, the alloy has the ordered B2 crystal structure from r o o m

* Corresponding author. 0921-5093/96/$15.00 © 1996 - - Elsevier Science S.A. All rights reserved pIT S0921-5093(96) 10480-0

temperature up to the melting. On the other hand, for an amount o f aluminium in the neighbourhood of 25 at.%, the ordered D O 3 crystal structure is present in the alloy. This DO3 structure transforms to the B2 structure above the critical ordering temperature of 814 K [31. Creep behavior of FeA1 intermetallic alloys have not been extensively studied because of low temperature brittleness. Mechanical properties at r o o m temperatures have been studied in detail on both single crystal and polycrystal alloys possessing both the D O 3 and B2 structures [3,4]. In general, it was found that the yield strength and work-hardening rate increase and the ductility decrease with increasing the alm-ninium content. There is sufficient evidence that iron rich FeA1 alloys are not inherently brittle at r o o m temperature and for an A1 content ranging from 25 to 45 at.% elongations up to 3% have been reported when testing in tension. Recent works have shown that the r o o m temperature ductility of polycrystalline DO3 and B2 structured FeA1 alloys is influenced by extrinsic factors such as addition of ternary elements, test environment and heat treatment [5-10]. Boron doped Fe-40at.%A1 alloy exhibits an improved r o o m temperature ductility compared with the b o r o n free alloy [6]. Boron segregates strongly at

94

AA. Yimenez, G. Frommeyer 1 Materials Science and Engineering A220 (t996) 93-99

the grain boundaries and thus grain-boundary strengthening is obtained. It has also been reported that the processing conditions strongly affect the mechanical properties of FeA1 alloys and a wide range of strengths and ductilities has been reported for some compositions. The higher the cooling rate after a thermal treatment at elevated temperature the higher the yield stress and strength and the lower the ductility at room-temperature [7]. The cooling rate effect is associated to vacancy retention. Furthermore, the beneficial effect of slow cooling after annealing is additive to the environmental effect. Recent studies associate the decrease in room temperature ductility in FeA1 alloys with the interaction of the aluminide with water vapour [8-10]. This environmental embrittlement involves the reaction of water vapour with aluminium atoms at the crack tips and release of hydrogen atoms. The hydrogen atoms diffuse into the metal and thereby enhance the crack propagation. The high temperature plastic flow properties of iron aluminides have been studied in alloys with an aluminium content ranging from 40 to 50 at.% [11,12]. Compression tests in the temperature range from 1100 to 1400 K showed two different concentration-independent creep mechanisms with the same activation energy ( g 4 5 0 kJ mol-t). At low deformation temperatures the FeA1 alloys creep with a stress of about 6 and at high deformation temperatures with a stress exponent of about 3. Both deformation mechanisms are dependent on grain size, but in an opposite manner. For the high stress exponent regimen strength increase with a decreasing grain size, while for the lower stress exponent regimen a decreasing strength with decreasing grain size was observed. The strength of the binary alloys at temperatures of interest can be improved by addition ternary elements [13,14]. Two different strengthening mechanisms in the FeA1 intermetallic system containing ternary additions have been reported. Alloys containing 1-5 at.% Cr, Ti, Mn, Co exhibit a single phase microstructure and a solid-solution strengthening mechanism operate. On the other hand, a secondphase strengthening mechanism operates in the alloys containing 0.8-5 at.% B, Zr, Ta, Nb, Re and Hf, in which precipitation of ternary intermetallic compounds pins the dislocations. The aim of the current work is to study the strain rate-flow stress behavior at high temperature of a series of three FeAl alloys ranging in aluminium from 25 to 48 at.% and two ternary FeA1Cr alloys. Mechanical properties were evaluated by strain-ratechange tests performed in compression in the temperature range from 750 to 1000°C.

2. Experimental procedure Ingots of three FeA1 binary alloys containing 25, 30 and 48 at.% A1 and two ternary FeA1Cr containing 21.7at.%A1-23.4at.%Cr and 30at.%Al-10at.%Cr were produced in a laboratory induction furnace under 40 kPa of argon. Parallelepiped compression samples were prepared by spark erosion from casted material. The compression samples had a square cross section, 6 x 6 ram, and 11 mm in height. All compression tests at high temperatures were performed under a protective atmosphere of argon to minimize oxidation. Elevated temperature mechanical properties were determined by strain-rate-change tests at strain rates ranging from 10 .5 to 10 .2 s -1 in a universal testing machine. The mechanical tests were performed in compression in the temperature range from 750 to 1000°C. In order to achieve the steady state, the samples were pre-strained to about 15% at a strain rate of about 1 x 10 .3 s - t True compressive strains, stresses and strain rates were calculated from the load-deformation curves with assumptions of uniform deformation and conservation of volume.

3. Results The high temperature behavior of FeA1 rich alloys has been studied due to their potential applications at elevated temperature. The as cast alloys present a coarse and equiaxial microstructure, with a grain size of about 500 gm. These alloys were tested in compression at 750, 800, 850, 900, 950 and 1000°C. The high temperature strain rate-flow stress data of crystalline materials are usually analyzed using a relationship of the form:

where ~ is the strain rate in the steady state, cr is the flow stress, E the Young's modulus, Qo the activation energy for creep, R the universal gas constant and T is the absolute test temperature. Young's moduli for the studied alloys at the various temperatures were taken from K6ster [15]. Therefore, n can be obtained from the slope of a log ~ versus log ~;/E data plot. From Eq. (1) it deduced that the activation energy for creep, Qo, can be calculated from the compression data using the well-known relation:

/0 In k\ Qo:

-

The results obtained are outlined below.

J.A. Jimenez, G. Frommeyer / Materials Science and Engineering A220 (i996) 93-99 i 0 "I

I

1 0 -2

--

1

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'1

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1 0 -3

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. . . . .

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1 0 -3

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I000°C 950°C 900°C 850~C 800°C 750°C 700°C

10-2

,, ¢

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95

.j

i

.d 10 -4

10 -4

i0 -s

i0 -s

J

(a) 1 0 -6

f

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1;.4

l l J , r l

'

(b)

......

1 0 -6

10 -3

1 0 -2

l

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Ir+Irl

r

1

10-4

1 0 .5

i

t

rtql[

i

10.3

+

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+

~

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• '+'+ //

o I:I

10-2

°



1 0 "3

7 .d

IO00°C 950°C 900°C

9

8~°c soo°c

/ / i

~a

///, p p

/

10 -4

10 4

(c) 1 0 -6

1 1 0 "s

, ......

I 1 0 -4

.....

~,,I 1 0 .3

....... 1 0 -2

o/E Fig. 1. Logarithm strain rate vs. logarithm Young's modulus-compensated flow stress for the FeA1 binary alloys (a) Fe-25at.%A1, (b) Fe-30at.%A1 and (c) Fe-48at.%AI.

3.1. Binary alloys The results o f strain-rate-change tests for the three binary alloys studied are plotted in Fig. l(a), (b) and (c) as logarithm strain-rate versus logarithm Young's modulus compensated flow stress. At temperatures above 850°C, the three alloys present a stress exponent of about 3.5. In this temperature range, the flow stressstrain rate-temperature characteristics are identical for the alloys with an aluminium content of 30 and 48 at.%. On the other hand, the alloy with an aluminium content of 25 at+% presents a lower creep resistance compared with these alloys.

At 800°C the alloy with a 25 at.% present a stress exponent of about 3.5, while for the alloy with 48 at.% a stress exponent of about 5 was deduced. The alloy with an aluminium content of 30 at.% present at this temperature a change from n ~ 5 at high strain rates to 3.5 at low strain rates. Finally, at testing temperatures o f 750°C an increase in the stress exponent in the three alloys studied, was observed. The activation energy was determined from the data of Fig. 1 (a), (b) and (c), and the results are shown in Fig. 2(a), (b) and (c), respectively. Values o f 280, 360 and 395 kJ m o l - i were obtained for the alloys F e 25at.%A1, Fe-30at.%A1 and Fe-48at.%, respectively.

96

J.A. Jimenez, G. Frommeyer f Materials Science and Engineering A220 (1996) 93-99

10-~

+

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'

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3x104

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Q = 395 -* 15 kJ/mol

8.0

8.5 9.0 lfr, k "l

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384 kJ/mol 368 kJ/mol '",n 358 kJlme Q = 360 _* 20 k J / m o l / ,

I

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Fig. 2. Activation energy for plastic flow for the FeA1 binary alloys (a) Fe-25at.%At, (b) Fe-30at.%A1 and (c) Fe-48at.%A1, from comression strain-rate-change tests.

3.2. Ternary alloys The results of strain-rate-change tests for two ternary alloys studied are plotted in Fig. 3(a) and (b) as logarithm strain-rate versus logarithm Young's modulus compensated flow stress. The alloy containing 21.7%at.%A123.4at.%Cr presents a behavior similar to that described for the binary alloys: at 750°C the alloy creeps with a stress exponent of about 5 and for higher temperatures with a stress exponent n of 3-4. A similar creep resistance of this alloy with the binary alloys containing 30 and 48 at.% of ahiminium is observed comparing the values of the flow stress-strain rate-temperature present in Fig. 3(a) with that of Fig. l(b) and (c). A value for the activation energy of 290 kJ tool - i was obtained from the data of Fig. 3(a) as shown in Fig. 4(a). The alloy containing 30 at.% A1 and 10 at.% Cr presents a creep behavior quite different to the other studied alloys. Fig. 3(b) shows two stress exponent regimes of creep at temperatures above 800°C. At low strain rates, a stress exponent of n ~ 3.5 and at high strain rates, an exponent of n ~ 5 is deduced. Each mode of deformation can be described by a different activation energy. Values for the activation energy of

350 and 505 kJ m o l - t were obtained from the data of Fig. 3(b) for the low and high stress regime, respectively, as shown in Fig. 4(b). Comparing the flow stress-strain rate-temperature data for the F e - 3 0 a t . % A l - 10at.%Cr alloy with that for the binary alloy with 30 at.% A1, an improvement in the creep resistance is observed for the alloy with chromium at strain rates lower than 2 x 10 -4 s -1. At higher strain rates, the alloy with chromium presents a higher n value, and the differences in creep resistance become smaller with the Fe-30at.%Al. Similar values of flow stresses were obtained at strain rates ranging from 10 . 3 to 10 . 2 s -1 in both alloys.

4. Discussion

The general correlation of both theoretical and experimental constitutive relations for high temperature deformation of crystalline material is usually given by the Dorn equation for dislocation creep:

ADo b Y V

~= ~

/

Qo

t~ ) expL---~-~)

(3)

J.A. dimenez, G. Frommeyer / Materials Science and Engineering A220 (1996) 93-99

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(a) ........ 1 0 -61 0 -5

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(b)

.......

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........ 10-5

I 10 4

/E

4. I. Binary alloys The present work shows a low stress exponent of n z 3.5 at temperature deformation above 800°C for the three binary alloys. This value indicates that creep is controlled by viscous glide of dislocation, a feature typical of class I of pure-metal creep. Creep of FeA1 as a solid-solution class I was reported by several authors and confirmed by Rudy and Sauthoff using transmission electron microscopy [17]. These authors found no

10 "2

o/E

Fig. 3. L o g a r i t h m strain rate vs. l o g a r i t h m Y o u n g ' s m o d u l u s - c o m p e n s a t e d flow stress for the F e - C r - A 1 23.4at.%Cr and (b) F e - 3 0 a t . % A l - 1 0 a t . % C r .

where D Ois the diffusion coefficient, G the shear modulus, k is the Boltzman's constant and A the microstructural and mechanism dependent constant. Several deformation mechanisms can occur in metallic and other materials depending on temperature and deformation rate. Each mechanism has characteristic values of the parameters A, Qo and n by which they can be defined uniquely. An extensive analysis of creep data of conventional disordered alloys shows that the stress exponent tends to cluster around the values 3 or 5 [16]. A stress exponent close to 5 is generally attributed to intragranular dislocation creep controlled by the climb of dislocations at dislocation pile-up, which gives rise to a subgrain structure. On the other hand, for metallic solid solution alloys, intragranular dislocation creep is often attributed to the sequential operation of glide and climb process. When solute drag reduces the rate of glide, dislocation glide becomes rate controlling and then a stress exponent of 3 is obtained. In this late case, dislocation tangles without formation of subgrains, are observed.

.......

'ternary alloys (a) F e - 2 1 . 7 a t . % A 1 -

subboundaries or subgrains structure in a FesoA150 sample after creep at 900°C. Flow stress-strain rate-temperature data show that the creep behavior of binary alloys is determined by the ordered structure present in the alloys. The FeA1 equilibrium diagram shows that a Fe-A1 alloy with an aluminium content up to 25 at.% has disordered e-Fe in the temperature range used, while for an aluminium content above 30 at.% has the ordered cubic B2 structure. The rupture of bonds between like atoms in the disordered Fe-25at.%A1 alloy makes the glide of dislocations easier. The velocity of glide of dislocation in the glide plane will be increased and then a lower creep resistance is obtained. On the other hand, the presence of the B2 long-order in the Fe-30at.%A1 and F e 48at.%A1 alloys at the tested temperatures determine the similar creep characteristic observed in both alloys. Two values of the activation energy for deformation have been estimated in function also of the degree of structural disorder in the alloy at the deformation temperatures used. The activation energy obtained for the disordered Fe-25at.%A1 alloy was 280 kJ m o l - i . This value agrees with those for lattice interdiffusion or tracer diffusion reported in the literature for FeA1 alloys with an aluminium content up to 25 at.%, which range from 280 to 310 kJ mo1-1. An activation energy for creep ranging from 360 to 395 kJ mo1-1 was estimated when the order B2 is present in the Fe-30at.%A1 and Fe-48at.% alloys. These values are quite similar to the values reported by Lawley et al. The authors, studying the creep behavior of FeA1 alloys containing between 10 and 27 at.% A1,

98

J.A. Jimenez, G. Frommeyer / Materials Science and Engineering A220 (I996) 93-99

found a value of 370 kJ m o l - ~ when the B2 order is present [18]. However, these activation energies are much smaller than the value of about 460 kJ m o l measured by Whittenberger in alloys with an aluminium content ranging f r o m 4t to 49 at.% [12]. 4.2. Ternary alloys

The alloy Fe-21.7at.%A1-23.4at.%Cr presents a similar creep behavior to the binary Fe-25at.%A1, since it presents also a disordered lattice in the temperature range used. A value of n ~ 3.5 and an activation energy of 290 kJ t o o l - a were measured also in this case. However, this alloy presents a higher creep resistance than the binary alloy. The results of the compression strain-rate-change tests at high temperatures clearly indicate a substitutional solution strengthening comparing with the binary B2 FeA1 alloy. The alloy containing 30at.%Al-10at.%Cr presents two different deformation regimes with a different associated activation energy. The low strain rate deformation mechanism has a value of n g 3.5 and Q ~ 350 kJ mol - 1 and the high strain rate deformation mechanism has a value of n ~ 5 and Q ~ 505 kJ m o l - ~. This stress

10-~ 10 -2

!290ld/mol 280 kJ/mol

...

10-3

.d

104 " ~ 5

10 4"~

7 x 10

10 4

Q = 290 ± 10 kJ/mol

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10-~ 10-s Q = 350 -* 10 k J/tool 10.6 r , , , I , , , i I i i r i 1 i .7.5 8.0, 8.5 9.0 l/T, k" 1

. (b) ,

i

i

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9.5

t

,

5. C o n c l u s i o n s

Creep of the three FeA1 binary alloys containing 25, 30 and 48 at.% A1 and the ternary Fe-21.7at.%A123.4at.%Cr at temperature deformation above 800°C occurs with a stress exponent of g 3.5, a feature typical of creep mechanisms controlled by viscous glide of dislocation. Two values of the activation energy, Qo, related with the degree of structural disorder in the alloy have been measured when the material creep with n ~ 3.5. The alloys with an aluminium content up to 25 at.% present the disordered c~-Fe structure and Qo ~ 280 kJ tool-1. However, the alloys with an aluminium content above 30 at.% have the ordered cubic B2 structure and Qo 360 kJ m o l - 1. The alloy containing 30at.%Al-10at.%Cr present two different deformation regimes. At low strain rate, similar values to that for the binary Fe-30at.%A1 alloy v~as obtained (n ~ 3.5 and Q ~ 350 kJ m o l - t). At high strain rate the alloy shows an increased creep resistance and a higher value of the activation energy for creep (n ~ 5 and Q ~ 505 kJ t o o l - 1).

~

b/E=7x 10~, 10:3Q = 505 ± 20 k J/tool

103

exponent indicates that creep is controlled by the climb of dislocations at dislocation pile-up, a feature typical of class II of pure-metal creep. Two deformation regimes have also been reported in the literature for binary F e - A 1 alloys with finer grain size (about 40 Ixm) under identical test conditions: a high stress exponent (n ~ 6) at low temperature or high strain rate a high, and a low stress exponent ( n g 3) at high temperature or low strain rate [12]. However these authors found the same activation energy for both mechanisms (about 460 kJ t o o l - ~). This activation energy is between that for the low stress regimen (350 kJ tool -1) and high stress exponent (505 kJ m o l - t ) found in the present work.

p

tO x 104

Fig. 4. Activation energy for plastic flow for the Fe-Cr=A1 ternary alloys (a) Fe-21.Tat.%Ai-23.4at.%Cr and (b) Fe-30at.%A110at.%Cr, from compression strain-rate-change tests.

References

[I] M. Hansen, Constitution of Binary Alloys, McGraw-Hill, New York, 1958, p. 90. [2] H.J. Goldschmidt, Interstitial Alloys, Butterworths, London, 1967. [3] K. Vedula, in J.H. Westbrook and R.L. Fleisher (eds.), Intermetallic Compounds; Vol 2, Practice, Wiley, New York, 1994, p. 199. [4] I. Baker, in J.H. Schneibel and M.A. Crimp (eds.), Processing, Properties and Applications of h'on Ahtm#Tides, The Minerals, Metals & Materials Society, 1994, p. 101. [5] Z.Q. Sun, Y.D. t-Iuang,W.Y. Yang and L. Chen, in I. Baker, R. Darolia, J.D. Whittenberger and M.H. Yoo (eds.), High-Temperature Ordered Intermetaltie Alloys V, Mat. Res. Soc. Syrup. Proc., Vol. 288, Materials Research Society, Pittsburgh, PA, 1993, p. 885. [6] C.T. Liu and E.P. Georg, Scripta MetalI. Mater., 24 (1990) 1285.

J.A. Jimenez, G. Frommeyer / Materials Science and Engineering A220 (t996) 93-99 [7] M.A. Crimp, K.M. Vedula and D.J. Gaydosh, in N.S. Stoioff, C.C. Koch, C.T. Liu and O. Izumi (eds.), High-Temperature Ordered Intermetallie Alloys II, Mat. Res. Soc. Symp. Proc., Vol. 81, Materials Research Society, Pittsburgh, PA, 1987, p. 499. [8] C.T. Liu, C.G. McKamey and E.H. Lee, Scripta Metal. Mater., 24 (I990) 385. [9] D.J. Gaydosh and M.V. Nathal, Seripta Metal. Mater., 24 (1990) 1281. [I0] C.G. McKamey and E.H. Lee, in I. Baker, R. Darolia, J.D. Whittenberger and M.H. Yoo (eds.), High-Temperature Ordered Intermetaltic Alloys V, Mat. Res. Soc. Syrup. Proc., Vol. 288, Materials Research Society, Pittsburgh, PA, 1993, p. 983. [11] J.D. Whittenberger, Mater. Sci. Eng., 57 (1983) 77. [12] J.D. Whittenberger, Mater. Sci. Eng., 77 (I986) 103.

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[13] K. Vedula and J.R. Stephens, in N.S. Stoloff, C.C. Koch, C.T. Liu and O. Izumi (eds.), High-Temperat~o'e Ordered Intermetaltic Alloys II, Mat. Res. Soc. Symp. Proc., Vol. 81, Materials Research Society, Pittsburgh, PA, 1987, p. 381. [14] M.G. Mendiratta, S.K. Ehlers, D.M. Dimiduk, W.R. Kerr, S. Mazdiyasni and H.A. Lipsitt, in N.S. Stoloff, C.C. Koch, C.T. Liu and O. Izumi (eds.), High-Temperature Ordered Intermetallic Alloys II, Mat. Res. Soc. Symp. Proc., Vol. 81, Materials Research Society, Pittsburgh, PA, 1987, p. 393. [15] W. K6ster and G. G6decke, Z. Metallkde., 73 (1982) 111. [16] O.D. Sherby and P.M. Burke, Prog. Mater. Sei., 13 (1967) 323. [17] M. Rudy and G. Sauthoff, Mater. Sci. Eng., 8t (1986) 525. [18] A. Lawley, J.A. Coll and R.W. Cahn, Trans. ]ffetall. Soc. AIME, 218 (1960) I66.