Anisotropic creep of Ni3(AlTiTa)

Sctipta Materialia, Vol. 37, No. 10, pp. 1491-1498, 1997 Elsevier Science Ltd Copyright 0 1997 Acta Metallurgica Inc. Printed in the USA. All rights reserved 1359-6462/97 $17.00 + 00

Pergrrmon

PI1 S1359-6462(97)00299-6

ANISOTROPIC

CREEP OF Nis(AlTiTa)

C. Knobloch, V.N. Toloraia*and U. Glatzel** Technische Universimt Berlin, Mikrostrukturanalyse, BH 18, 10623 Berlin, Germany *Russian l.nst. of Aviation Materials (VIAM), 17 Radio Street, 107005 Moscow, Russian Fed. **Friedrich-Schiller-Universimt Jena, Metallische Werkstoffe, Lobdergraben 32,07743 Jena, Germany (Received April 2 1, 1997) (Accepted July 3, 1997) Introduction The intermetallic phase N&Al, besides being a major constituent of nickel based superalloys, is itself of importance for high temperature structural applications. Therefore research has extended to N&Al-base composites in recent time (1,2). For the better understanding of its mechanical properties, numerous investigations were carried out. In respect of the temperature anomaly of the yield stress, see (3). The explanation of creep mechanisms still shows some lacks. Caron, Khan and Veyssiere (4) have investigated active slip systems during creep and the dependence of creep rate on the chemical composition, the orientation
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melt of a low melting point metal (aluminium). The parameters for structure forming, such as the axial thermal gradient was chosen to 7-9”/mm, the growth rate to 5 mm/min. The given o~entations were achieved by using seed crystals. Initially for this purpose y’-phase seeds out of arbitrary oriented single crystals had been used, but their application failed. The reason is that during heating of the mold for metal pouring (up to T - 1520- 1550°C) a dense film of thermodynamically stable oxides (A120J, NiAhO) is formed on the contact surface of the seed. It prevents the contact between seed and melt and consequently the seed crystal can not act as starting point for single crystalline solidification. Therefore seeds of an alloy based on a nickel solid solution that did not contain elements forming stable oxides (Al, Ti, Cr) were used for y’-phase single crystal production. As was earlier shown (9) precipitation of the eutectic y’-phase in single crystal nickel alloys practically under all casting conditions takes place in orientation correspondence with the y-matrix. This results in eutectic y’-y globules originating at second and third order dendrite arms. During pouring of the liquid metal into the mold the y’ melt dissolves partially the y-phase seed resulting in the formation of a layer with alternating al~inium content along its height. The depth of this layer amo~ts to about 300400 pm. In the course of subsequent directional solidification, solid solution dendrites with y/y’ eutectic precipitates are formed in this layer. The amount of precipitates increases throughout its height. Dendrites of the base solid solution as well as y’ precipitates grow together. Inasmuch as these precipitates are in the orientation confo~i~ with the y solid solution, y’ phase single crystals having the orientation given by the seed are formed. The single crystal y’ phase structure is a results of dendritic segregation heterogeneity. The value of subgrain misorientation amounts to about 2”. From the rods, tensile creep specimens were prepared. Specimen orientation was verified using a Laue back-resection X-ray technique. The creep tests were performed in air under a constant load of 350 MPa at 1123 K. The tests were halted at different stages of creep. Samples were cooled down under load in order to preserve the microstructure. Transmission electron microscope (TEM) analysis was carried out using a Jeol2OOC and a Philips CM 30 microscope. Results and Discussion Fig. 1 shows the creep behavior of different oriented samples at a temperature of 1123 K and an applied stress of 350 MPa. Creep rates are initially low and increase gradually within a region of less than 1% strain. This is also known as inverse primary creep (8,lO). Inverse primary creep can be _____I_______-

A2

l$(AlTiTa) I 123 K, 350MPa --r.,,...,..I..~1.0

2.0 swain [%]

3.0

WI

IO0 2o” 30” 4o”wl]

4.0

Figure 1. Creep curves of Ni,AI single crystals plotted as strain rate versus strain. The orientation of each specimen is marked in the unit triangle.

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weakly deformed

jeformed

Figure 2. a) Specimen A 1, obtained strain: 0.05%; b) A 2, obtained strain: 6.7%; c) Specimen B 1, obtained strain: 0.03%; d) B 2, obtained strain: 6.0%; e) Specimen C 1, obtained strain: 0.07%; fJ C 2, obtained strain: 4.5%.

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understood by viscous dislocation glide, where the dislocation velocity v is independent of dislocation density p. Using B = p . b . v, the strain rate S thus increases with dislocation density until a constant density and creep rate is reached. All orientations achieve a steady-state region. The [OOl]-direction is the orientation with the highest creep rate and the steady-state creep rate of [ 11 I] is the lowest, one third that of [OOI]. The [01 I] orientation is only slightly faster than [ 11 I]. This result is in qualitative agreement with Shah and Cetel(7).

50 40 zap30 20 10 0

HO-11 [01-l][lOl] [01-l][-1011[Oil] [loll [Oli] act. [Oil] pass.

iiiij

jiiij

(iii) &rirj il-11)(l-11)

(11-l) (11-l)

?

(01-l) gs

a) [OOl]weakly deformed sample A 1, 35 dislocations considered

octahedral glide planes

(111) (111) (-ill)(-lll)(l-11) (1-ll)(ll-1)(11-l) (011) (101)

gs .

b) [001]stronglydeformedsample A 2,44 dlslocatlons constdered

octahedral glideplanes cube glideplanes

[lo-l][-llO] [loll [llo] [llq [l-lo] [loll [IO-I] act. pass. (111) (111) (-111)(-111) (001) (001) (010) (010) 7 gs

c) [Oll]weaktydeformed

sample B 1,32 dislocations considered

Figure 3. Active octahedral and cubic slip systems for the different orientations and obtained strain. All passive systems (Schmid factor close to zero) are grouped together in the last column. (Figure continued.)

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1495

(10-l][-llq [loll [ilO] [llO] [loll[10-l](llq (-1011 act. pass. (111) (111)(-111)(-111)(001) (010) (010)(l-10)(101) 7 gs

d) (011]stronglydeformedsample B 2,46 dislocationsconsidered 50

40 -8%

I

perfect screws

cube glide octahedral glide planes

.

.

planes

[loll [llq [llq [Oil] [loll [Oil] [llq [loll IO111 act. pass. 7 gs

(-111)(-111)(l-11)(1-11)(11-l)(11-l)(001) (010) (100)

e) [11l] weaklydeformedsample C 1,50 dislocationsconsidered

octahedral glide planes .

[loll [llq (llq [Oil] [loll [Oil] [llq (1011 [Oil] pass. (-111)(-111)(l-11)(l-11) (11-l) (11-l) (001) (010) (loo) gs f)[11l]

stronglydeformedsample C 2,42 dislocationsconsldered Figure 3. (Figure continued.)

Fig. 2 compares dislocation structures of slightly deformed samples with the higher deformed state for each orientation. Starting with a low dislocation density, a subsequent dislocation multiplication takes place during primary creep, see Fig. 4. In Fig. 2 a), c) and e) (weakly deformed samples), long straight screw dislocations which glide on only few planes can be seen. As a result of higher deformation, reactions between dislocations take place, see Fig. 2 b), d) and f). The dislocati80n distribution has been found to be homogeneous for all orientations and creep stages. Subgrain boundaries or dislocation networks are observed only occasionally. This is in agreement with

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0%

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undeformed sample

0

I

2

3

Strain

4

5

6

7

[ %]

Figure 4. Change of dislocation density during creep for different oriented samples. The lines show how one would expect the dislocation densities to develop between the available data points.

Hemker and Nix (8) and Nathal (10). Scanning electron microscope observations show that subgrain boundaries correspond to dendriteiinterdendritic transitions for all orientations and creep stages. Change of subgrain size with creep deformation was not observed. Second phase-precipitates and stacking faulfs have not been detected. Through tilting experiments, Burgers vector b, line vector Z and, with sxb or by minimal curvature, the normal vector to the glide plane was determined. Fig. 3 shows the ascertained slip systems for the different specimens. Burgers vector of the type have not been verified yet. The weakly deformed [OOI] sample A 1, see Fig. 2 a) and 3 a) is mainly deformed by glide on two systems: [ lOl]( i 11) and [Ol l]( 11 i), whereas for the strongly deformed sample A 2, Fig. 2 b) and 3 b), the observed dislocations are distributed homogeneously on 8 active octahedral slip systems. This observation is similar for the other orientations. During primary creep a gradually increasing number of slip systems is activated. When the secondary creep range is reached all possible active slip systems contribute to the deformation. For [0 1l] and [ 11 I] orientations cube glide systems are activated even in weakly deformed samples, see Fig. 3 c) to 3 f). Schmid factor considerations are listed in Table 1 together with the comparision of package density of different glide systems. Bold printed fields indicate that these systems have been observed by TEZM

TABLE 1 Comparison of Glide Systems with Burgers Vector Gliding onEither { Ill}, { 100) or { 110) Planes. The Package Density Compared to the { 111) Most Densely Packed Planes, Schmid Factors and Theoretical Number of Active Systems for Different Load Axis arelisted. Bold Printed Combinations Have Been Observed by TEM in the Stationary Creep Stage package of planes (1111 glide planes, 100% b= (001) glide planes, 87% b= (01 I) glide planes, 61%

I

i;=<110>

1

I

[OOI] load axis m, = 0.41

active systems: 8

m,=O active systems: 0 m. = 0.5 active systems: 4 1

I [Ol I] load axis

[lll]loadaxis

m, = 0.41 active systems: 4 m, = 0.35 dive systems: 4

m, = 0.27 active systems: 6 m3 = 0.47 active systems: 3

m, = 0.25 active3systems: 4

m=O active &Wna: 0

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80

7o T

h -

60 50

840

'a

p 30 c=: 20 I 10 0 I 0

[ 11I] stress avis orientation

z

43e .

SCtXW

1

2

3

60” .30”

4

5

6

7

Strain [“h]

Figure 5. Change in dislocation character during creep for [ 11l] oriented samples

in strongly deformed samples. Obviously all glide systems with Schmid factors more than 0.3 have been activated in the secondary creep stage, regardless of the package density of these planes. Fig. 5 shows the change of dislocation character during deformation as an example for the [ 1111 orientation. The amount of screw dislocation decreases drastically, whereas the fractional appearance of mixed dislocations increases. The character of the dislocations has changed from predominant screw dislocations during primary creep to mixed and edge dislocations during secondary creep. This indicates interactions between slip systems, also resulting in a higher fraction of passive systems in the secondary stage of creep, see Fig. 3 b), 3 d) and 3 f). Conclusions

Creep tests with Ni3Al single crystals with different orientations close to the edges of the standard orientation triangle were performed. A number of active slip systems was found to contribute to the deformation. Activation of cube slip could be verified. With growing dislocation density more and more active slip systems are activated, interactions occur which lead to a change in the character of the dislocations and an increasing amount of passive glide systems. All slip systems with high Schmid factors (m, > 0.3) have been activated. These observations give insight in the important role of dislocation interactions in-between different slip systems. The components of an interaction matrix between different slip systems can be extracted by comparing absolute dislocation densities of slip systems as function of plastic strain. These interaction matrix will be needed for a model based on evolution equations for d.islocation densities of different slip systems. Acknowledgments

This work was financially supported by the Deutsche Forschungsgemeinschaft (DFG) within the Gerhard-Hess Programm and by the BMBF in the frame of “New Materials”. The authors wish to thank Dr. J. Ziebs, Dr. H. Frenz and E. Kramer from the Bundesanstalt ftir Materialforschung und -prtifung (BAM) for performing the creep tests. Many helpful discussions with Prof. Dr. M. Feller-Kniepmeier are gratefully acknowledged.

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References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10.

N.S. Stoloff, Intern. Materials Reviews, 34 (1989) 153. G.L. Povirk, J.A. Horton, C.G. McKamey, T.N. Tiegs and S.R. Nutt, J. Mater. Sci. 23 (1988) 3945 D.P. Pope and S.S. Ezz, Intern. Metals Reviews, 29 (1984), 136. P. Caron, T. Khan and P. Veyssiere, Phil. Mag. A, 60A (1989), 267. B.H. Kear and H.G.F. Wilsdorf, Trans. AIME 224 (1962), 382. D.M. Shah, Scripta Metall., 17 (1983), 997. D.M. Shah and A. Cetel, Superalloys 1996, Minerals, Metals & Materials Sot. (1996), 273. K.J. Hemker and W.D. Nix, Metal]. Trans. A, 24A (1993), 335. G. Lutzau, E.P. Kostyukova and V.N. Toloraia, Izv. AN SSSR, Metally, N3 (1978), 167. M.V. Nathal, J.V. Diaz and R.V. Miner, Mat. Res. Sot. Symp. Proc., 133 (1989), 269.

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