- Email: [email protected]
Single and duplex creep tests at intermediate temperatures on Ni3Al
MATERIALS SCIENCE & ENGWEERING
ELSEVIER
A
Materials Science and Engineering A233 (1997) 26-32
Single and duplex creep tests at intermediate N&Al
temperatures on
I.P. Jones *, T.S. Rong, R.E. Smallman School
of Metallurgy
and Mareriais,
rind IRC
in Mutevials for High Performance Applications, Birmingham B15 2TT, UK
The Unirersity
of Birmingham,
Edgbaston,
Received 29 August 1996; received in revised form 18 November 1996
Abstract Creep tests and parallel transmission electron microscopy (TEM) studieson N&Al in the intermediatetemperatureregime (wherestrengthis increasingwith temperature)are described.Under certain circumstancesinversecreepis observed.Lower stress and higher
temperature
both inhibit
the transition
to cube cross-slip.
Pretesting
at constant
strain rate can both increase or
decreasethe subsequent creep rate, depending on the temperature of the pretest. Pre-creep testing also influences subsequent constant strain rate deformation. All of these observations are rationalised in terms of the dislocation geometry. 0 1997 Elsevier Science S.A. Keywords:
Creep testing; Intermetallic compounds; Ni,Al
1. Introduction In constant strain rate testing N&AI, like many other LI, intermetallic compounds, displays three different temperature regimes (Fig. 1). In region II there is an ‘anomalous’ increase in strength with temperature,
\
III Temperature Fig. 1. Schematic of the three temperature regimes of strength for Ll, compounds. * Corresponding
author.
0921-5093/97]$17.00 Q 1997 Elsevier Science S.A. All rights reserved. P1~s0921-5093(97)00045-2
which is linked to thermally activated cross slip. Such phenomena, in fact, are not restricted to ordered compounds: the same thing happens to some h.c.p. elements [l]. Region IT is what we have chosen here to call ‘intermediate temperatures’. Creep of N&AI at intermediate temperatures is particularly fascinating because under certain conditions inverse creep occurs [2]. This is when, during a creep test, the rate of creep suddenly starts to increase and is quite different from the tertiary stage where voids start to form and which precedes fracture. Inverse creep of N&Al was shown by Hemker and colleagues [2] to coincide with a movement of the dislocations from the primary octahedral slip plane to the cube cross-slip plane. Depending on other factors this can lead to an increase in creep rate. Dislocation glide during constant strain rate testing is now thought [3,4] to be accomplished by the movement of edge superkinks separated by Kear-Wilsdorf locks, where screw dislocations have cross-slipped from the octahedral plane to the cube cross-slip plane. It is an extraordinary and significant observation by Hemker et al. [2] that the dislocations inhabit the cube cross-slip plane only in the later stages of creep, but never the primary cube slip plane. One is led to the conclusion that the Kear-Wilsdorf locks referred to
I.P. Jones et al. /Materials
Scien?emGld
above are acting as very efficient sources for slip on the cube cross-slip plane. The fact that the transition from octahedral to cube slip takes place not at the beginning of the test but, if at all, at some point later, implies that some condition or other has to be met. The obvious requirement is that the length of the source dislocation (KW lock) has to be sufficiently great that the propagation stress /lb/L is less than the applied stress g. There is thus a very intimate connection between the microstructural development and the mechanical properties. We have performed both single and duplex tests, involving both constant strain rate testing and creep. We will first give a brief summary of the results of our single creep tests on Ni,Al at various stresses and temperatures and the associated microstructural development and then we will describe the duplex tests.
2. Experimental
details
In the present investigation, single and duplex creep tests were carried out on polycrystal and single crystal N&Al, respectively. A polycrystalline Ni,Al button with a nominal composition Ni75at.SA125at.% was prepared by plasma melting under an argon atmosphere. Before cutting into specimens for single creep tests, it was homogenised at 1050°C for 3 days in vacuum and slowly cooled in the furnace. This procedure resulted in a material with a weak texture. For the duplex creep tests, a single crystal of Ni,,Ni,,,,Ta,,, (at.%) was employed. The creep tests were performed at constant load in compression in air or vacuum, on specimens with dimensions about 4.5 x 4.5 x 9.0 mm for polycrystals and 3.5 x 4.0 x 7.5 mm for single crystals, In an attempt to retain the high-temperature deformation structure, the specimens were quenched into an ice bath within a few seconds of removal after creep tests. After mechanical tests, thin foils for transmission electron microscopy (TEM) observation were cut at 45” with respect to the compression axis for polycrystal specimens and for single crystal specimens sectioned to make either the primary octahedral or cube cross-slip plane parallel to the specimen. Normally, six or seven TEM foils were examined for each different mechanical test. Transmission electron microscopy involved a Philips CM20 operated at 200 kV and a JEOL 4000FX operated at 200 or 400 kV.
3. Single creep tests Fig. 2 shows results of two creep tests at 580°C (which lies below the peak stress r,) at different stresses. The higher stress curve illustrates the phenomenon of inverse creep. The creep rate ceases subse-
Engineering
A233
(1997)
27
26-32
0.7
0.2:
o0
,,,,.,,,,,,,,,, 10
,), 20
30
40
),, ,,, 50 60 Time (hour)
,,, 70
,,, 80
,,,. 90
1
Fig. 2. Creep curves for polycrystalline N&Al at 580°C at 250 and 400 MPa.
quently to increase because the test is in compression and the stress effectively drops. Note that at the lower stress there is no inverse creep. Fig. 3 shows the microstructures after 196 h at the higher stress. All the dislocations are on a cube plane which we infer, following Hemker et al. [2], to be the cube cross-slip plane. (This is a polycrystalline specimen.) The inset shows the dislocation structure at an early stage, i.e., the primary regime where the creep rate is decreasing: the dislocations are on an octahedral plane. TEM of the lower stress test reveals a low dislocation density which remains on the octahedral plane. A similar test at 250 MPa, but at 380°C also shows no inverse creep regime, but here the dislocations are on the cube cross-slip plane (as we infer) throughout most of the test, having made a very early transfer from the primary octahedral plane. The transfer from octahedral to cube cross-slip is partly a question of the operation of the Kear-Wilsdorf locks as sources and partly a question of the blocking of the superkink movement on the primary octahedral slip plane. Fig. 4 shows the edge end of a loop after 10 h at 580°C and 400 MPa. The two partials have climbed one above the other (Fig. 5). Either the kink has been stopped by the increasing forest density, followed subsequently by climb, or the climb itself has halted the kink. This type of observation has never been reported after constant strain rate deformation below Tp [5,6] (except by Takasugi and Yoshida [7], who did pot quench their specimen after the test). We therefore believe that it is the climb dissociation which terminates primary creep. Once the dislocations have started to glide on the cube cross-slip plane. they multiply via a double crossslip mechanism (see, e.g., Fig. 6). As a jogged screw moves on the cube cross-slip plane, an edge dipole is dragged out. If the height of the dipole is large enough, two edge segments glide apart to form a single-turn helix. Then, if cross-slip occurs, an expanding loop may be created [8].
I.P. Jones et al. ,’ Mnterinls
Science
and Engineeriq
A.?33 (1997)
26-32
Fig. 3. A low magnification micrograph showing the dislocation structure in polycrystalline Ni,AI crept at 400 MPa for 196 h at 58O”C, where two kinds of (110) dislocations operate and share a common cube plane. Dislocations consist of nearly pure edge {e) and screw (s) segments. The insert shows the dislocation morphology after a creep test was interrupted at 10 h, in which screw segments (s) predominate.
To summarise the observations above
250 MPa
400 MPa
380°C
580°C
Octahedral and cube cross-slip; no inverse creep
Octahedral slip; no inverse creep Octahedral and cube cross-slip + inverse creep
We believe that these observations can be understood in terms of the geometry of the dislocations on the primary octahedral slip plane. A series of theoretical papers, e.g., Paidar et al. [9], led up to two simulation papers (Mills and Chrzan [IO], Devincre et al.’ [I 11) which attempted to predict the shape of the primary octahedral slip dislocations, i.e., the kink and by implication the Kear-Wilsdorf lock distribution. Although the models were different-Mills and Chrzan [lOj used a localised pinning model and Devincre et al. ll l] a more realistic Kear-Wilsdorf lock formation model (although two-dimensional)-the conclusions were much the same:
’ We are grateful to the authors for a preprint of this paper.
As time increases at constant stress, the larger kinks run out and thus the whole kink size distribution (including average length) moves downwards. One would therefore expect the Kear-Wilsdorf lock length distribution to move upwards. l As temperature increases and there is a greater frequency of cross-slip. the kink distribution moves down (smaller average kink size). We would therefore expect smaller Kear-Wilsdorf locks on average. l As the (constant) stress increases, fewer small kinks are left. The experimental observations of Couret et al. [S] confirm at least some of these results, viz. the overall shapes of the distributions, and that as the temperature rises, the kinks decrease in size, averagely. Turning back to the little matrix of results above, if we take 58O”C/400 MPa as the starting condition, then reducing the temperature accelerates the change to cube cross-slip while decreasing the stress has the opposite effect. This is consistent with the trends described immediately above. Decreasing the temperature increases the average Kear-Wilsdorf lock lengfh (because of less frequent cross-slip) and therefore these rapidly become viable sources. Decreasing the stress allows more small kinks to remain, which in their turn decrease the average Kcar-Wilsdorf lock length and inhibit the transformation to cube cross-slip. Because we observe a quicker transition to cube slip at the lower temperature, we l
I.P. Jones er al. /Materials
Science
and Engirmring
AZ33
(1991)
26-32
29
Tilting Path
<11Os=(111)
alip
Fig. 4. A weak-beam tilting sequenceindicates and reveals the planes of the loop and the dissociation for (1 IO) { 11I} slip in a specimencrept for 10 h at 400 MPai580”C.
infer that the lengths of the KW locks are more important than the relative increase of friction on the
4. Duplex tests
cube plane as compared with that on the octahedral plane. We now turn to the duplex tests.
We have investigated the effect of prior creep on constant strain rate tests and the effect of prior constant strain rate deformation on creep.
30
I.P. Jones et al. ,/itkzterinls
(111) -L-
- -A
Science
cwi Engineering
A.233 (1997)
26-32
-L
Fig. 5. A schematic showing climb dissociation of the edge se,ments which terminates the Primary octahedral phase of creep,
4.1. Creepfollotved by constant strain rate testing Fig. 7. Bright-field image revealing the dislocation structure on the cube cross-slip Plane in a single crystal N&AI after conventional constant strain rate deformation (IO-’ s-‘) to ez 0.025 at 380°C. The dislocations are mainly of a long and straight screw type.
Here we have tested single crystals at 380°C in compression along [932]. (This crystal and that referred to in the next section were simply what were available to us.) Fig. 7 shows the microstructure after a conventional constant strain rate (lo-” s-l) test (E ~2.5%). As expected, the microstructure is dominated by octahedral slip and long Kear-Wilsdorf locks. Fig. 8 shows the microstructure after a duplex test with a total plastic strain about 2.5%: 45 h of creep at 225 MPa, followed by a constant strain rate of 10m4 s-l (E z 2%). The slip is now on the cube cross-slip plane. (This time we know this because we are using a single crystal.) Thus the prior creep test has completely changed the nature of the constant strain rate deformation. The creep test has established cube cross-slip and this maintains itself throughout the subsequent constant strain rate deformation despite the friction stress on the cube cross-slip plane being higher then that on the octahedral plane, as it is throughout the intermediate temper-
ature regime. The cube cross slip dislocations are not as rectangular as those in Fig. 3. This is because the deformation is occurring at a lower temperature and implies that the formation of double and super LomerCottrell locks requires thermal activation. Cube crossslip never occurs in ordinary constant strain rate tests because the kinks are constantly being replenished by the rising stress. Once the dislocations are on the cube cross-slip plane they will be reluctant to cross-slip back to the octahedral plane, because this is a state of higher energy. The fact that new octahedral sources do not operate (either in creep or? as here, under constant strain rate conditions) must be a comment on the ratio of the friction stresses on cube and octahedral slip planes. Presumably, at a low enough temperature, this may cease to be true.
d) Fig. 6. Tilting of the foil to reveal the geometry of a single helical dislocation in a specimen crept at 580°C and 400 MPa for 196 h.
I.P. Jones et al. /Materials
4.2. Constant strGimmfe
defonnntior~ followed
Science
and Engi%ZeFini
by creep
Here we tested single crystals of Ni,(AlO.Sat.%Ta) in compressive creep along [155i] at 52O”C, 330 MPa and for 5 h. We used three conditions: 1. a virgin specimen; and, after constant strain rate deformation of 3% at 10n4 s-l, 2. at room temperature (RT), 3. at 520°C. Fig. 9 shows the results. The RT pre-deformation makes the specimen creep more quickly than the virgin specimen; the 520°C predeformation has the opposite effect.
3 c H x g m
A233
(1997)
31
26-32
500 450 400 350 300 250 200 Constantstrain-rate deformation at RT 150 100 50 o,.,,,,,,,.,,,.,,,,,.,,...,,,, 0 1 2 3 4 strain (%)
Creep at
520°C for 5h
5
//,,/ 6
Strain (70)
0
I
2
3
4
5
6
Strain (%) Fig. 9. Stress-strain curves: (a) for a duplex test involving constant strain rate deformation and creep. The specimen was pre-strained at constant strain rate (lOWJ s- ‘) at room temperature for 3% and was then heated up to 520°C and crept at 330 MPa for 5 h. (b) Virgin specimen crept at 520°C and 330 MPa for 5 h. (c) Stress-strain curves corresponding to a duplex test involving constant strain rate (lo-” s- ‘) deformation to 3% strain at 520°C and then creep at 520°C.
Fig. 8. Bright-field low-magnification tilting showing cube cross-slip after a combination test involving creep (45 h and 225 MPa) and constant strain rate deformation (lo-“ s- ‘) at 380°C. The loops are at their widest when the cube cross-slip plane is perpendicular to the foil.
The slower creep after high temperature deformation is hardly surprising and we attribute it to work hardening. The more surprising result, perhaps, is that predeformation at room temperature increases the subsequent creep rate. (Note that the creep deformation is in the primary (octahedral) region, as confirmed for all three situations by slip trace analysis both after the constant strain rate tests and after the creep tests.) We believe that the low temperature prestrain has introduced a distribution of very mobile dislocations, which subsequently give faster creep deformation. They are mobile because the average kink size at lower tempera-
32
I.P. Jones et al. ;‘Materials
Science
ture is larger, as predicted by Mills and Chrzan [lo], Devincre et al. (unpublished) and as observed by Couret et al. [5].
5. Summary
We have attempted to establish that the results of a multiplicity of creep and constant strain rate deformation tests on N&Al can be understood by a consideration of the geometry of the dislocations slipping on the primary octahedral plane and that our observations are consistent with the predictions of Mills and Chrzan [lo], Devincre et al. (unpublished) and with the observations of Couret et al. 151.
Acknowledgements
We wish to thank Professor J.F. Knott for the provi-
and Engineering
A233
(1997)
26-32
sion of laboratory facilities, Dr Y. Bi for the single crystals of Ni,AI, and Professor P. Bowen for access to mechanical testing equipment.
References 111I.P. Jones, W.B. Hutchison, Acta Metallurgica 29 (1981) 951. PI K.J. Hemker, M.J. Mills. W.D. Nix, Acta Metallurgica 39 (1991) 1901.
P.B. Hirsch, Philosophical Magazine A 65 (1992) 569. M.J. Mills, N. Baluc, H.P. Karnthaler, High-Temperature Ordered Intermetallic Alloys, Materials Research Society Symposium Proceedings 133 (1989) 203. PI A. Couret, Y.Q. Sun, P.M. Hazzledine, Materials Research Society Symposium Proceedings 213 (1991) 317. @I Y.Q. Sun, Ph.D. Thesis, Oxford University, UK, 1990. [71 T. Takasugi, M. Yoshida, Philosophical Magazine A 67 (1993) [31 [41
447.
J.S. Koehler, Physics Review 86 (1952) 52. V. Paidar, D.P. Pope, V. Vitek, Acta Metallurgica 32 (1984) 435. [lOI M.J. Mills, D.C. Chrzan, Acta Metallurgica 40 (1992) 3051. 183 I91
ELSEVIER
A
Materials Science and Engineering A233 (1997) 26-32
Single and duplex creep tests at intermediate N&Al
temperatures on
I.P. Jones *, T.S. Rong, R.E. Smallman School
of Metallurgy
and Mareriais,
rind IRC
in Mutevials for High Performance Applications, Birmingham B15 2TT, UK
The Unirersity
of Birmingham,
Edgbaston,
Received 29 August 1996; received in revised form 18 November 1996
Abstract Creep tests and parallel transmission electron microscopy (TEM) studieson N&Al in the intermediatetemperatureregime (wherestrengthis increasingwith temperature)are described.Under certain circumstancesinversecreepis observed.Lower stress and higher
temperature
both inhibit
the transition
to cube cross-slip.
Pretesting
at constant
strain rate can both increase or
decreasethe subsequent creep rate, depending on the temperature of the pretest. Pre-creep testing also influences subsequent constant strain rate deformation. All of these observations are rationalised in terms of the dislocation geometry. 0 1997 Elsevier Science S.A. Keywords:
Creep testing; Intermetallic compounds; Ni,Al
1. Introduction In constant strain rate testing N&AI, like many other LI, intermetallic compounds, displays three different temperature regimes (Fig. 1). In region II there is an ‘anomalous’ increase in strength with temperature,
\
III Temperature Fig. 1. Schematic of the three temperature regimes of strength for Ll, compounds. * Corresponding
author.
0921-5093/97]$17.00 Q 1997 Elsevier Science S.A. All rights reserved. P1~s0921-5093(97)00045-2
which is linked to thermally activated cross slip. Such phenomena, in fact, are not restricted to ordered compounds: the same thing happens to some h.c.p. elements [l]. Region IT is what we have chosen here to call ‘intermediate temperatures’. Creep of N&AI at intermediate temperatures is particularly fascinating because under certain conditions inverse creep occurs [2]. This is when, during a creep test, the rate of creep suddenly starts to increase and is quite different from the tertiary stage where voids start to form and which precedes fracture. Inverse creep of N&Al was shown by Hemker and colleagues [2] to coincide with a movement of the dislocations from the primary octahedral slip plane to the cube cross-slip plane. Depending on other factors this can lead to an increase in creep rate. Dislocation glide during constant strain rate testing is now thought [3,4] to be accomplished by the movement of edge superkinks separated by Kear-Wilsdorf locks, where screw dislocations have cross-slipped from the octahedral plane to the cube cross-slip plane. It is an extraordinary and significant observation by Hemker et al. [2] that the dislocations inhabit the cube cross-slip plane only in the later stages of creep, but never the primary cube slip plane. One is led to the conclusion that the Kear-Wilsdorf locks referred to
I.P. Jones et al. /Materials
Scien?emGld
above are acting as very efficient sources for slip on the cube cross-slip plane. The fact that the transition from octahedral to cube slip takes place not at the beginning of the test but, if at all, at some point later, implies that some condition or other has to be met. The obvious requirement is that the length of the source dislocation (KW lock) has to be sufficiently great that the propagation stress /lb/L is less than the applied stress g. There is thus a very intimate connection between the microstructural development and the mechanical properties. We have performed both single and duplex tests, involving both constant strain rate testing and creep. We will first give a brief summary of the results of our single creep tests on Ni,Al at various stresses and temperatures and the associated microstructural development and then we will describe the duplex tests.
2. Experimental
details
In the present investigation, single and duplex creep tests were carried out on polycrystal and single crystal N&Al, respectively. A polycrystalline Ni,Al button with a nominal composition Ni75at.SA125at.% was prepared by plasma melting under an argon atmosphere. Before cutting into specimens for single creep tests, it was homogenised at 1050°C for 3 days in vacuum and slowly cooled in the furnace. This procedure resulted in a material with a weak texture. For the duplex creep tests, a single crystal of Ni,,Ni,,,,Ta,,, (at.%) was employed. The creep tests were performed at constant load in compression in air or vacuum, on specimens with dimensions about 4.5 x 4.5 x 9.0 mm for polycrystals and 3.5 x 4.0 x 7.5 mm for single crystals, In an attempt to retain the high-temperature deformation structure, the specimens were quenched into an ice bath within a few seconds of removal after creep tests. After mechanical tests, thin foils for transmission electron microscopy (TEM) observation were cut at 45” with respect to the compression axis for polycrystal specimens and for single crystal specimens sectioned to make either the primary octahedral or cube cross-slip plane parallel to the specimen. Normally, six or seven TEM foils were examined for each different mechanical test. Transmission electron microscopy involved a Philips CM20 operated at 200 kV and a JEOL 4000FX operated at 200 or 400 kV.
3. Single creep tests Fig. 2 shows results of two creep tests at 580°C (which lies below the peak stress r,) at different stresses. The higher stress curve illustrates the phenomenon of inverse creep. The creep rate ceases subse-
Engineering
A233
(1997)
27
26-32
0.7
0.2:
o0
,,,,.,,,,,,,,,, 10
,), 20
30
40
),, ,,, 50 60 Time (hour)
,,, 70
,,, 80
,,,. 90
1
Fig. 2. Creep curves for polycrystalline N&Al at 580°C at 250 and 400 MPa.
quently to increase because the test is in compression and the stress effectively drops. Note that at the lower stress there is no inverse creep. Fig. 3 shows the microstructures after 196 h at the higher stress. All the dislocations are on a cube plane which we infer, following Hemker et al. [2], to be the cube cross-slip plane. (This is a polycrystalline specimen.) The inset shows the dislocation structure at an early stage, i.e., the primary regime where the creep rate is decreasing: the dislocations are on an octahedral plane. TEM of the lower stress test reveals a low dislocation density which remains on the octahedral plane. A similar test at 250 MPa, but at 380°C also shows no inverse creep regime, but here the dislocations are on the cube cross-slip plane (as we infer) throughout most of the test, having made a very early transfer from the primary octahedral plane. The transfer from octahedral to cube cross-slip is partly a question of the operation of the Kear-Wilsdorf locks as sources and partly a question of the blocking of the superkink movement on the primary octahedral slip plane. Fig. 4 shows the edge end of a loop after 10 h at 580°C and 400 MPa. The two partials have climbed one above the other (Fig. 5). Either the kink has been stopped by the increasing forest density, followed subsequently by climb, or the climb itself has halted the kink. This type of observation has never been reported after constant strain rate deformation below Tp [5,6] (except by Takasugi and Yoshida [7], who did pot quench their specimen after the test). We therefore believe that it is the climb dissociation which terminates primary creep. Once the dislocations have started to glide on the cube cross-slip plane. they multiply via a double crossslip mechanism (see, e.g., Fig. 6). As a jogged screw moves on the cube cross-slip plane, an edge dipole is dragged out. If the height of the dipole is large enough, two edge segments glide apart to form a single-turn helix. Then, if cross-slip occurs, an expanding loop may be created [8].
I.P. Jones et al. ,’ Mnterinls
Science
and Engineeriq
A.?33 (1997)
26-32
Fig. 3. A low magnification micrograph showing the dislocation structure in polycrystalline Ni,AI crept at 400 MPa for 196 h at 58O”C, where two kinds of (110) dislocations operate and share a common cube plane. Dislocations consist of nearly pure edge {e) and screw (s) segments. The insert shows the dislocation morphology after a creep test was interrupted at 10 h, in which screw segments (s) predominate.
To summarise the observations above
250 MPa
400 MPa
380°C
580°C
Octahedral and cube cross-slip; no inverse creep
Octahedral slip; no inverse creep Octahedral and cube cross-slip + inverse creep
We believe that these observations can be understood in terms of the geometry of the dislocations on the primary octahedral slip plane. A series of theoretical papers, e.g., Paidar et al. [9], led up to two simulation papers (Mills and Chrzan [IO], Devincre et al.’ [I 11) which attempted to predict the shape of the primary octahedral slip dislocations, i.e., the kink and by implication the Kear-Wilsdorf lock distribution. Although the models were different-Mills and Chrzan [lOj used a localised pinning model and Devincre et al. ll l] a more realistic Kear-Wilsdorf lock formation model (although two-dimensional)-the conclusions were much the same:
’ We are grateful to the authors for a preprint of this paper.
As time increases at constant stress, the larger kinks run out and thus the whole kink size distribution (including average length) moves downwards. One would therefore expect the Kear-Wilsdorf lock length distribution to move upwards. l As temperature increases and there is a greater frequency of cross-slip. the kink distribution moves down (smaller average kink size). We would therefore expect smaller Kear-Wilsdorf locks on average. l As the (constant) stress increases, fewer small kinks are left. The experimental observations of Couret et al. [S] confirm at least some of these results, viz. the overall shapes of the distributions, and that as the temperature rises, the kinks decrease in size, averagely. Turning back to the little matrix of results above, if we take 58O”C/400 MPa as the starting condition, then reducing the temperature accelerates the change to cube cross-slip while decreasing the stress has the opposite effect. This is consistent with the trends described immediately above. Decreasing the temperature increases the average Kear-Wilsdorf lock lengfh (because of less frequent cross-slip) and therefore these rapidly become viable sources. Decreasing the stress allows more small kinks to remain, which in their turn decrease the average Kcar-Wilsdorf lock length and inhibit the transformation to cube cross-slip. Because we observe a quicker transition to cube slip at the lower temperature, we l
I.P. Jones er al. /Materials
Science
and Engirmring
AZ33
(1991)
26-32
29
Tilting Path
<11Os=(111)
alip
Fig. 4. A weak-beam tilting sequenceindicates and reveals the planes of the loop and the dissociation for (1 IO) { 11I} slip in a specimencrept for 10 h at 400 MPai580”C.
infer that the lengths of the KW locks are more important than the relative increase of friction on the
4. Duplex tests
cube plane as compared with that on the octahedral plane. We now turn to the duplex tests.
We have investigated the effect of prior creep on constant strain rate tests and the effect of prior constant strain rate deformation on creep.
30
I.P. Jones et al. ,/itkzterinls
(111) -L-
- -A
Science
cwi Engineering
A.233 (1997)
26-32
-L
Fig. 5. A schematic showing climb dissociation of the edge se,ments which terminates the Primary octahedral phase of creep,
4.1. Creepfollotved by constant strain rate testing Fig. 7. Bright-field image revealing the dislocation structure on the cube cross-slip Plane in a single crystal N&AI after conventional constant strain rate deformation (IO-’ s-‘) to ez 0.025 at 380°C. The dislocations are mainly of a long and straight screw type.
Here we have tested single crystals at 380°C in compression along [932]. (This crystal and that referred to in the next section were simply what were available to us.) Fig. 7 shows the microstructure after a conventional constant strain rate (lo-” s-l) test (E ~2.5%). As expected, the microstructure is dominated by octahedral slip and long Kear-Wilsdorf locks. Fig. 8 shows the microstructure after a duplex test with a total plastic strain about 2.5%: 45 h of creep at 225 MPa, followed by a constant strain rate of 10m4 s-l (E z 2%). The slip is now on the cube cross-slip plane. (This time we know this because we are using a single crystal.) Thus the prior creep test has completely changed the nature of the constant strain rate deformation. The creep test has established cube cross-slip and this maintains itself throughout the subsequent constant strain rate deformation despite the friction stress on the cube cross-slip plane being higher then that on the octahedral plane, as it is throughout the intermediate temper-
ature regime. The cube cross slip dislocations are not as rectangular as those in Fig. 3. This is because the deformation is occurring at a lower temperature and implies that the formation of double and super LomerCottrell locks requires thermal activation. Cube crossslip never occurs in ordinary constant strain rate tests because the kinks are constantly being replenished by the rising stress. Once the dislocations are on the cube cross-slip plane they will be reluctant to cross-slip back to the octahedral plane, because this is a state of higher energy. The fact that new octahedral sources do not operate (either in creep or? as here, under constant strain rate conditions) must be a comment on the ratio of the friction stresses on cube and octahedral slip planes. Presumably, at a low enough temperature, this may cease to be true.
d) Fig. 6. Tilting of the foil to reveal the geometry of a single helical dislocation in a specimen crept at 580°C and 400 MPa for 196 h.
I.P. Jones et al. /Materials
4.2. Constant strGimmfe
defonnntior~ followed
Science
and Engi%ZeFini
by creep
Here we tested single crystals of Ni,(AlO.Sat.%Ta) in compressive creep along [155i] at 52O”C, 330 MPa and for 5 h. We used three conditions: 1. a virgin specimen; and, after constant strain rate deformation of 3% at 10n4 s-l, 2. at room temperature (RT), 3. at 520°C. Fig. 9 shows the results. The RT pre-deformation makes the specimen creep more quickly than the virgin specimen; the 520°C predeformation has the opposite effect.
3 c H x g m
A233
(1997)
31
26-32
500 450 400 350 300 250 200 Constantstrain-rate deformation at RT 150 100 50 o,.,,,,,,,.,,,.,,,,,.,,...,,,, 0 1 2 3 4 strain (%)
Creep at
520°C for 5h
5
//,,/ 6
Strain (70)
0
I
2
3
4
5
6
Strain (%) Fig. 9. Stress-strain curves: (a) for a duplex test involving constant strain rate deformation and creep. The specimen was pre-strained at constant strain rate (lOWJ s- ‘) at room temperature for 3% and was then heated up to 520°C and crept at 330 MPa for 5 h. (b) Virgin specimen crept at 520°C and 330 MPa for 5 h. (c) Stress-strain curves corresponding to a duplex test involving constant strain rate (lo-” s- ‘) deformation to 3% strain at 520°C and then creep at 520°C.
Fig. 8. Bright-field low-magnification tilting showing cube cross-slip after a combination test involving creep (45 h and 225 MPa) and constant strain rate deformation (lo-“ s- ‘) at 380°C. The loops are at their widest when the cube cross-slip plane is perpendicular to the foil.
The slower creep after high temperature deformation is hardly surprising and we attribute it to work hardening. The more surprising result, perhaps, is that predeformation at room temperature increases the subsequent creep rate. (Note that the creep deformation is in the primary (octahedral) region, as confirmed for all three situations by slip trace analysis both after the constant strain rate tests and after the creep tests.) We believe that the low temperature prestrain has introduced a distribution of very mobile dislocations, which subsequently give faster creep deformation. They are mobile because the average kink size at lower tempera-
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ture is larger, as predicted by Mills and Chrzan [lo], Devincre et al. (unpublished) and as observed by Couret et al. [5].
5. Summary
We have attempted to establish that the results of a multiplicity of creep and constant strain rate deformation tests on N&Al can be understood by a consideration of the geometry of the dislocations slipping on the primary octahedral plane and that our observations are consistent with the predictions of Mills and Chrzan [lo], Devincre et al. (unpublished) and with the observations of Couret et al. 151.
Acknowledgements
We wish to thank Professor J.F. Knott for the provi-
and Engineering
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(1997)
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sion of laboratory facilities, Dr Y. Bi for the single crystals of Ni,AI, and Professor P. Bowen for access to mechanical testing equipment.
References 111I.P. Jones, W.B. Hutchison, Acta Metallurgica 29 (1981) 951. PI K.J. Hemker, M.J. Mills. W.D. Nix, Acta Metallurgica 39 (1991) 1901.
P.B. Hirsch, Philosophical Magazine A 65 (1992) 569. M.J. Mills, N. Baluc, H.P. Karnthaler, High-Temperature Ordered Intermetallic Alloys, Materials Research Society Symposium Proceedings 133 (1989) 203. PI A. Couret, Y.Q. Sun, P.M. Hazzledine, Materials Research Society Symposium Proceedings 213 (1991) 317. @I Y.Q. Sun, Ph.D. Thesis, Oxford University, UK, 1990. [71 T. Takasugi, M. Yoshida, Philosophical Magazine A 67 (1993) [31 [41
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J.S. Koehler, Physics Review 86 (1952) 52. V. Paidar, D.P. Pope, V. Vitek, Acta Metallurgica 32 (1984) 435. [lOI M.J. Mills, D.C. Chrzan, Acta Metallurgica 40 (1992) 3051. 183 I91