Superalloys 2008
Roger C. Reed, Kenneth A. Green, Pierre Caron, Timothy P. Gabb, Michael G. Fahrmann, Eric S. Huron and Shiela A. Woodard
TMS (The Minerals, Metals & Materials Society), 2008
TENSION/COMPRESSION ASYMMETRY IN YIELD AND CREEP STRENGTHS OF
NI-BASED SUPERALLOYS
N. Tsuno1, S. Shimabayashi1, K. Kakehi1C.M.F. Rae2 and R.C. Reed3
Dept of Mechanical Engineering, Tokyo Metropolitan University; 1-1 Minami-Osawa, Hachioji, Tokyo 192-0397, JAPAN
2
University of Cambridge, Department of Materials Science and Metallurgy;
Pembroke Street, Cambridge, CB2 3QZ, UK
3
Dept of Metallurgy and Materials, The University of Birmingham; Edgbaston, Birmingham, B15 2TT, UK
1
Keywords: Single crystal, Creep, Twinning, Stacking fault, Tension, Compression
width of the core of the a/2[101] superpartial dislocation must be
changed by the effect of the components of the stress tensor when
the superpartial dislocation in the γ′ phase undergoes cross-slips
from {111} to {100}. Since these a/2[101] superpartial
dislocations dissociate into Shockley partials on the {111} plane,
any component of the stress tensor that constricts superpartial
dislocations promotes cross-slip, and any component that extends
them impedes cross-slip. According to the “core width effect” [7],
the tensile flow stress is greater than the compressive flow stress
for orientations near [001].
Abstract
The tension/compression asymmetries of yield and creep strengths
of three kinds of single-crystal superalloys—PWA1480, CMSX-4,
and TMS-75—and a DS superalloy, Mar-M247LC, were investigated at intermediate and high temperatures. In PWA1480, tensile
yield strength was higher than the compressive strength from
20°C to 750 °C. From the TEM observation, it was found that the
asymmetry of yield strengths is primarily due to the microtwin
formation associated with a superlattice extrinsic stacking fault
(SESF). In CMSX-4 and TMS-75, tensile/compressive yield
strengths were comparable at every temperature, and shearing of
γ′ precipitates by a/2<110> dislocations pairs was the dominant
deformation mechanism in both tensile and compressive tests at
750 °C. The creep response of these materials were quite different
than their yielding response. CMSX-4 and TMS-75 showed distinctive creep tension/compression asymmetry. These two alloys
showed large creep strain caused by {111}<112> slip at 750 °C
under tensile stress, and mechanical twins at 750 °C and 900 °C
under compressive stresses. Tension/compression asymmetry of
CMSX-4 and TMS-75 was larger at 900 °C than at 750 °C because <112> viscous slip was not observed under tensile stress at
900 °C. The asymmetric nature of PWA1480 was small at both
750 °C and 900 °C because the dominant deformation mode in
both tension and compression is a combination of both a/2<110>
dislocation's bowing on the {111} plane in the matrix and climbing along the γ/γ′ interfaces.
Under thermo-mechanical fatigue (TMF), the number of cycles to
failure of a single-crystal superalloy was found to decrease drastically with an increase of the compressive strain hold time [8]. The
compressive creep strength of single-crystal superalloys is lower
than the tensile creep strength on [001] loading [2]; therefore,
high-tensile residual stress is produced by the larger stress relaxation during the compressive hold. This would decrease the TMF
strength of the single-crystal superalloy. These phenomena are
due to mechanical twinning associated with the operation of a
{111}<112> viscous slip. Hence, in order to improve TMF life
under condition of a compressive strain hold, it is important to
clarify the mechanism of tension/compression asymmetry in the
creep behavior in order to inhibit the operation of <112> viscous
slip. Little attention has been given to the tension/compression
asymmetry at intermediate temperatures. In this paper, emphasis
is placed on creep and yield strengths in the temperature range
750-900°C to explain the phenomena of asymmetry behavior in
single-crystal superalloys.
Introduction
Experimental Procedure
Ni-based single-crystal superalloys exhibit tension/compression
asymmetry in yield strength and creep behavior at low
temperatures [ 1 , 2 ]. The flow stress behavior of Ni-based
superalloys has been explained by the thermally activated crossslip model suggested for the single γ′ phase [3,4,5]. Takeuchi and
Kuramoto predicted the crystallographic-orientation dependence
of the flow stress of Ni3Ga single crystals assuming that
dislocations on the {111} slip planes become locally pinned by
thermally activated cross-slip onto {100} planes [ 6 ].
Tension/compression asymmetry in the yield strength of Ni-based
superalloys was explained in terms of the “core width effect” by
Lall, Chin and Pope [7 ]. They modified the "cross-slip effect"
model proposed by Takeuchi and Kuramoto, and assumed that the
Co
The investigation was carried out for a first generation SX superalloy, a second generation SX superalloy, as well as a third generation SX superalloy; i.e., PWA1480, CMSX-4 and TMS-75. A
directionally solidified (DS) superalloy, Mar-M247LC, was also
tested to study the influence of grain boundaries on the intermediate temperature properties. The nominal chemical compositions of
the alloys tested in this study are listed in Table I. The alloys were
solution heat-treated and subsequently aged using the respective
conditions shown in Table II. After analyzing of the crystallographic orientation by the back-reflection Laue method, specimens were cut from a single as-aged bar using a spark cutter. The
Table I. Chemical composition of alloys (mass %)
Cr Mo W Al Ti Ta Re
C
B
Zr
Hf
Ni
PWA1480
5
10
-
4
5
1.5
12
-
-
-
-
-
bal.
CMSX-4
TMS-75
MAR-M247LC
9
12
9.4
6.5
3
8
0.6
2
0.5
6
6
9.5
5.6
6
5.6
1
0.7
6.5
6
3.2
3
5
-
0.07
0.015
0.01
0.1
0.1
1.5
bal.
bal.
bal.
433
Table II. Heat treatment condition of alloys
Solution heat treatment
PWA1480
CMSX-4
TMS-75
MAR-M247LC
1282°C/1 hr + 1287°C/2 hr + 1294°C/1 hr/GFC
1277°C/2 hr + 1288°C/2 hr + 1296°C/3 hr + 1304°C/3 hr +
1313°C/2 hr + 1316°C/2 hr + 1318°C/2 hr + 1321°C/2h/GFC
1240°C/1 hr + 1280°C/2 hr + 1300°C/2 hr + 1320°C/8 hr/GFC
1230°C/2hr/GFC
Aging treatment
1080°C/4 hr/AC+ 870°C/32
hr/AC
1140°C/4 hr/GFC +
870°C/20 hr/GFC
1150°C/4 hr/GFC +
870°C/20 hr/GFC
980°C/5 hr/GFC + 870°C/20
hr/GFC
1300
Tension
Compression
Yield Stress MPa
1200
(a)
1100
1000
900
(a)
800
0
1300
400
600
Temperature °C
Tension
Compression
1200
Yield Stress MPa
200
800
―
g=111
200 nm
(b)
(b)
_
1100
g=111
1000
900
800
0
200
400
600
Temperature °C
800
1300
Tension
Compression
Yield Stress MPa
1200
(c)
200 nm
Fig. 2. TEM images of the deformed structure in CMSX4
specimen after yielding: (a) tension, ε=1.5% and (b) compression, ε=4.3%.
1100
1000
tensile specimens and 10.0×3.0×3.0 (mm) in compressive specimens. Tensile tests were performed at 20 °C, 400 °C and 750 °C
under strain rates of 4.25×10-4 s-1 (tension) and 8.33×10-4 s-1
(compression). Tensile and compressive creep tests were performed at 750 °C under 750 MPa and 900 °C under 392 MPa in
the <001> orientation. The specimens for TEM observation were
cut from interrupted specimens along the desired crystallographic
plane and electro-polished using a twin jet with a solution of 15%
perchloric acid in methanol at 25 V and 5 °C. Specimens were
observed with a JEOL 2000FX microscope at 200 kV and a JEM3010 microscope at 300 kV. The Burgers vectors and the natures
of stacking faults were determined using the contrast visibility of
dislocations and stacking fault fringes [9].
900
800
0
200
400
600
800
Temperature °C
Fig. 1. Yield strengths of tensile and compressive tests under
initial strain rates of 4.25×10-4 s-1 (tension) and 8.33×10-4 s-1
(compression): (a) PWA1480, (b) CMSX-4 and (c) TMS-75.
The points are the results of single tests.
orientation of all the specimens were within 4 degrees of <001>.
The dimensions of the gage length were 19.6×2.8×3.0 (mm) in
434
(a)
(b)
―
g=111
＿
g=111
200 nm
200 nm
Fig. 3. TEM microstructure of the yielded PWA1480 specimens: (a) bright-field image at 1.9 % tensile strain, the foil normal [111],
Beam Direction (BD)～112; and (b) g-3g weak beam image at 2.7 % compressive strain, the foil normal [111], BD～112.
(a)
(b)
―
g=111
γ′
S.F.
_ _
a/3[211]
_ _
a/6[211]
200 nm
(d)
(c)
γ′
γ′twin
γ′
(111)
―
g=111
200 nm
1 nm
Fig. 4. TEM microstructure of PWA1480 of compressive specimens after yielding at 2.7%. (a) Dislocation structure, Beam Direction
－－
(BD)～112. (b) Schematic diagram illustrating that the a/3[211] dislocation cuts the γ′ phase and leaves a superlattice stacking fault
in the γ′ phase. (c) Microstructure showing micro twin, BD~110 and (d) lattice image of circled area shown in (c).
CMSX-4 and TMS-75 at every temperature, whereas PWA1480
showed tension/compression asymmetry of the yield strengths.
The tensile yield strength of PWA1480 was higher than the compressive strength from 20 °C to 750 °C. The tensile yield strength
increased from 400 °C to 750 °C. In CMSX-4, the shearing γ′ by
Results
Yield strength
The yield stress as defined by the 0.2% proof stress is shown in
Fig. 1. Tensile/compressive yield strengths were comparable in
435
sion Electron Microscopy (HRTEM) image (Fig. 4(d)) of the linear structure circled in Fig. 4(c), it was found to be a mechanical
twin having 9-12 atomic planes in its width. This twin was called
a “microtwin” [11]. These results indicate that a stacking fault is
indeed associated with twin formation in the γ′ phase.
Table III. The proportion of SISF and SESF observed in tensile and compressive PWA1480 specimens at 750°C in (%).
Tension
Compression
SISF
SESF
75.0
37.5
25.0
62.5
Creep strength
a/2<110> dislocation pairs was the same (Fig. 2(a), (b)) for both
tension and compression. Tensile and compressive yield microstructures in PWA1480 are shown in Fig. 3. The tensile microstructure showed matrix dislocations and many stacking faults in
γ′ (Fig. 3(a)). The percentages of the numbers of superlattice intrinsic stacking faults (SISF) and superlattice extrinsic stacking
faults (SESF) in the tensile and compressive specimens of
PWA1480 at 750 °C are shown in Table III. The relative percent
occurs in reverse for tension and compression: The applied
stresses favored SISF formation in tension and SESF formation in
compression. The compressive microstructure revealed matrix
dislocations, many stacking faults and linear dislocations in the γ′
phase, which are shown by arrows (Fig. 3(b)). Using the invisibility criterion [9], the analysis of the Burgers vectors and line directions [10,11] revealed that linear dislocations in the γ′ phase were
of the a/2<110> type on {111} planes. Cutting of the γ′ phase by
the superlattice dislocation and the stacking fault, colored black,
was also observed (Fig. 4(a)). Through the analysis of the Burgers
vectors and line directions [9,10], it was determined that the Burgers vector of the dislocation in the γ′ phase was a/3<112>. The γ′
phase was cut by a a/3<112> dislocation, as illustrated in Fig. 4(b).
The compressive TEM microstructure of the (110) foil is shown in
Fig. 4(c). This image is an edge-on view of the fault using the
<110> zone axis. From the enlarged High Resolution Transmis-
Tensile and compressive creep curves at 900 °C are presented in
Fig. 8. Anisotropy between tension and compression was observed in CMSX-4 and TMS-75. As shown in Fig. 9, in all three
alloys, it was observed that dislocations were concentrated into
3
(a)
(b)
Creep strain ε (%)
Creep strain ε (%)
15
Fig. 5 shows the tensile and compressive creep curves at 750 °C.
In CMSX-4 and TMS-75, the compressive creep curve showed
inferior strength compared to the tensile creep curve. In contrast,
the creep curves of PWA1480 were comparable. The microstructures of tensile and compressive interrupted creep tests in TMS-75
and PWA1480 are presented in Fig. 6. In TMS-75 the <112> ribbons which were composed of extended stacking faults and
a/3<112> dislocations were observed both in tension and compression (Fig. 6(a), (b)). The dislocations in tension and compression, which are indicated by arrowheads, were type a/3<112> on
{111}. The tangled matrix dislocations were observed both in
tension and compression in PWA1480 (Fig. 6 (c), (d)). The tangled dislocations in tensile specimens were of the type a/2<110>
on {111}. The dominant mechanism in creep deformation was the
a/2<110> dislocation motion, which occurred mainly in the γ
matrix channels. In TMS-75 and CMSX4, mechanical twin lamellae in the compressive creep microstructure were confirmed by
{011} plane observation, as shown in Fig. 7. They were not observed in the bright-field image of PWA1480 (Fig. 7(c)).
Compression
10
Tension
5
2
Compression
1
0
0
0
20
20
40
60
Time, t /h
80
0
100
30
(c)
Compression
15
10
Tension
5
Creep strain ε (%)
Creep strain ε (%)
Tension
0
20
50
100
Time, t /h
150
200
(d)
Compression
Tension
10
0
100
150
200
0
50
100
150
200
Time, t /h
Time, t /h
Fig. 5. Tensile and compressive creep curves of (a) Mar-M247LC, (b) PWA1480, (c) CMSX-4 and (d) TMS-75 at 750°C under
750 MPa.
0
50
436
(a)
(b)
― ―
g=111
― ―
g=111
200 nm
200 nm
(d)
(c)
―
g=111
―
g=202
200 nm
200 nm
Fig. 6. Dislocation microstructure of interrupted creep specimens: (a) TMS75 at 1.4 % tensile strain, the foil normal [111], Beam
Direction (BD)～112; (b) TMS75 at 1.8 % compressive strain, the foil normal [111], BD～112; (c) PWA1480 at 0.4 % tensile strain,
the foil normal [111], BD～111; and (d) PWA1480 at 0.8 % compressive strain, the foil normal [111], BD～112.
(a)
g=200
200 nm
(b)
200 nm
(c)
g=200
1 μm
g=200
Fig. 7. Compressive microstructure of interrupted creep specimens at 750 °C, the foil normal [011], Beam Direction (BD)～011: (a)
TMS75 at 25 % strain; (b) CMSX4 at 27 % strain; and (c) PWA1480 at 2 % strain.
the γ phase, and very few dislocations occurred in the γ′ phase. It
is likely that the tensile creep deformation occurs by a combination of <110> slip in the matrix and dislocation's climbing along
the γ/γ′ interface. The creep curves under tensile stress showed the
transition from secondary creep to tertiary creep (Fig. 8). Fig. 10
shows the TEM microstructures of compressive creep specimens.
The PWA1480 microstructure showed a concentration of dislocations in the γ phase and a few stacking faults in γ′. On the other
hand, in CMSX-4 and TMS-75 specimens, distinctive stacking
fault pairs were observed. The Burgers vectors of partial dislocations that formed stacking fault pairs were determined. The Bur-
gers vectors of dislocation A were identified as a/3<112>, and
those of dislocation B were identified as a/6<112>. Primary creep
strains in CMSX-4 and TMS-75 (Fig. 8) would be brought about
by the motion of the stacking fault pairs. These creep curves of
these alloys showed the transition from primary creep to secondary creep (Fig. 8(c), (d)). Moreover, deformation twins were
observed in the compressive specimens in these two alloys (Fig.
10(d)).
437
4
5
(b)
Creep strain ε (%)
Creep strain ε (%)
(a)
3
Compression
2
Tension
1
Compression
3
Tension
2
1
0
0
0
5
10
Time, t /h
15
20
0
20
3
3
40
60
Time, t /h
80
100
(d)
(c)
Tension
Creep strain ε (%)
Creep strain ε (%)
4
2
Compression
1
Compression
2
Tension
1
0
0
0
100
200
Time, t /h
300
400
0
100
200
Time, t /h
300
400
Fig. 8. Tensile and compressive creep curves of (a) Mar-M247LC, (b) PWA1480, (c) CMSX-4 and (d) TMS-75 at 900°C under
392 MPa.
However, if long-range ordering produces a superlattice structure,
the ordinary twinning mode will become a pseudo mode which
results in incorrect ordering in the sheared lattice. In the superalloys, an a/6 <112> Shockley partial is unable to pass through the
γ′ precipitate easily because it creates a high-energy fault. If
a/6<112> Shockley dislocations shear the γ′ precipitates, the L12
ordered structure will be destroyed. Kolbe [12] suggested a mi-
Discussions
Deformation mechanisms of creep
The tension/compression asymmetry was distinctive in the creep
strengths of CMSX-4 and TMS-75. In CMSX-4 and TMS-75, the
compressive creep curve showed inferior strength compared to the
tensile creep curve. Under compressive stresses, as shown in Fig.
7, the TEM micrograph of the [001] specimen showed microtwin
lamellae extending over both the γ and γ′ phases. The results suggest that the asymmetry is also caused by mechanical twinning
under compressive stress. Tension/compression asymmetry of
creep strength in the [001] orientation is attributed to the directional-formation characteristic of the mechanical twins; i.e., twin
formation occurs only under [001] compressive stress conditions,
and its formation stress is lower than the critical resolved shear
_
_
_
dislocation glides on the (111) plane, a CSF (Complex Stacking
_
_
dislocation cannot follow the first on the same (111) plane because it produces an energetically unfavorable structure. There_
fore, it is assumed that it glides on a neighboring (111) glide plane.
_
_
The third dislocation follows on the next (111) plane. However, in
the L12 ordered structure, the movement of dislocation produces
an energetically unfavorable pseudotwin. A pair of double steps of
diffusion reestablishing order in the pseudotwin, or "interchange
shuffles", will follow the dislocation movement.
Twinning is not a dominant deformation mechanism in metals
which possess many possible slip systems; it generally occurs
when the slip systems are restricted or when something increases
the critical resolved shear stress so that the twinning stress is less
than the stress for slip. The yield stress of γ′ precipitates increases
with increasing temperature up to a peak temperature (650 °C –
800 °C). Since the critical stress for twinning is lower than that for
the slip near the peak temperature, novel mechanical twins can be
expected to occur in the Ni-based superalloys. The twinning shear
is always directional in the sense that the shear in one direction is
not equivalent to the shear in the opposite direction. Twinning in
_
_
Fault) is produced behind the dislocation. The second a/6[112 ]
stresses for <110>{111} and [112](111) slip systems.
_
_
crotwin-formation model for the superalloys. After an a/6 [112]
The final result of the process is the creation of thin twin lamellae
that extend through both the γ and γ′ phases as shown in Fig. 7.
The twin dislocation rotates around its pole on every successive
{111} plane [13]. By such a mechanism, a twin can propagate
very quickly with the interchange shuffles, and is not subjected to
the need for a separate partial dislocation source on every (111)
plane [14]. Therefore, the nucleation of the twin must be the ratedetermining step rather than the propagation process. Massive
microtwins (Fig. 7) were observed at the same time as SISF/
SESF stacking fault pairs (Figs. 6, 10). The nucleation of the microtwin was then associated with the nature of superlattice stacking faults.
_
FCC metals occurs on (1 11) planes in the [1 12 ] direction. In
[001]-compressive deformations, the resolved shear stress on the (
_
111) plane acts in the twinning direction, and the formation of the
mechanical twin results in weakness.
In contrast with CMSX-4 and TMS-75, in PWA1480 and DS
Mar-M247LC, the extent of creep asymmetry was small. The
438
promotion of the Suzuki locking effect of the {111}<112> slip as
a result of the segregation of tantalum into the SSF [19]. It is reasonable to expect that the dislocation motion in the γ′ phase would
be restrained by the effects of tantalum. As a result, the dominant
deformation mode of PWA1480 both in tension and compression
should be a combination of both a/2<110> dislocation bowing on
the {111} plane in the matrix and the climbing process along the
γ/γ′ interface; that is to say, the creep rates would be controlled by
the dislocation climb process.
(a)
g=200
Single-crystal Mar-M247LC shows a distinct {111}<112> slip
[ 20 ]; however, in case of the DS alloy, tension/compression
asymmetry associated with the {111}<112> slip was not observed.
It seems reasonable to suppose that the grain boundaries inhibit
{111}<112> slip system operation.
400 nm
(b)
Effect of temperature on creep deformation mechanisms
g=200
The ratio of compressive to tensile primary-creep strain as a function of the test temperature is illustrated in Fig. 11. Tension/compression asymmetry was more pronounced at 900 °C
than at 750 °C in CMSX-4 and TMS-75. In general, the dominant
deformation mechanism of creep in the temperature range above
850 °C is viscous slipping of a/2<110> dislocations and the
climbing process along the γ/γ′ interface, and the <110> slip does
not show a clear primary creep strain [21, 22 ]. In the present
study, tensile creep specimens showed no primary creep strain
(Fig. 8) and the Burgers vector was determined as <110> by the
contrast visibility of the dislocation (Fig. 9) at 900 °C. The resolved critical shear stress for the {111}<110> slip falls above the
peak temperature of yield stress. As a result, at 900°C, tensile
creep deformation occurred from a combination of an a/2<110>
slip and climbing along the γ/γ′ interface. We can speculate on the
cause of tension/compression asymmetry of creep at 900 °C based
on our observation of tension/compression asymmetry at both
750°C and 900 °C in CMSX-4 and TMS-75. In general, as mentioned above, the dominant deformation mechanism in the temperature range above 850 °C is the viscous slip of a/2<110> dislocations and the climbing process along the γ/γ′ interface; the
stacking fault mode of deformation plays a minor role at high
temperatures. However, in this study, the γ and γ′ phases were
sheared by <112> dislocation pairs, and microtwins were observed under compressive stress at 900 °C. The critical shear
stress of twin formation associated with <112> slipping appears to
be lower than that for <110> slipping even at 900°C. One explanation for this may be that twinning nucleation is associated with
the critical radius of the extrinsic and intrinsic stacking fault loop,
as described below.
400 nm
(c)
g=200
400 nm
Fig. 9. Dislocation structures of interrupted tensile crept
specimens at 900 °C.(a) PWA1480 (ε=0.5%), (b) CMSX-4
(ε=0.5%) and (c) TMS-75 (ε= 0.4%). BD~001.
crept PWA1480 specimen in this study did not show the mechanical twin structure that was observed in CMSX-4 and TMS-75 (Fig.
7). The dominant deformation mode of PWA1480 both in tension
and compression is a combination of a/2<110> dislocation bowing on the {111} plane in the matrix and climbing along the γ/γ′
interface (Fig. 6). In an Ni-based superalloy containing a high
level of tantalum, PWA1480, the {111}<112> slip was prohibited
and the critical stress of mechanical twin formation would be
increased because of the strengthened γ′ phase. A high level of
tantalum is expected to induce the following effects: the solidsolution strengthening of the γ′ phase resulting from lattice parameter changes [15, 16], the increase of APB energy on (111)
[15] (The APBE of Ni3Al is increased from 180 mJm−2 to 240
mJm−2 when substituting 1 at% Ta for 1at% Al) [17, 18], and the
Knowles [23] described the correlation between the formation of
deformation twins during creep and the nature of superlattice
stacking faults in CMSX-4. Tension/compression asymmetry is
brought about by the following mechanism associated with <112>
slip. SISF (Superlattice Intrinsic Stacking Fault) tends to form
under tensile stress and SESF (Superlattice Extrinsic Stacking
Fault) under compressive stress. Superlattice stacking faults are
formed by two Shockley partial dislocations. Under compressive
stress, the leading partial is subjected to a high Schmid factor;
therefore, the partial enters the γ′ phase. After it cuts the γ′, the
trailing partial enters the γ′ phase to form SESF. On the other
hand, under tensile stress, the Schmid factor for the trailing partial
is higher than that of the leading partial. As a result, it is possible
for the trailing dislocation to cross the leading dislocation and
439
(a)
(b)
__
B
g=111
A
__
g=111
400 nm
200 nm
(c)
(d)
g=200
B
A
___
g=111
400 nm
200 nm
Fig. 10. Dislocation structures of interrupted compressive crept specimens: (a) PWA1480 (ε = 2.1 %), (b) CMSX-4 (ε = 1.1 %), and
(c) TMS-75 (ε = 1.2 %), Beam Direction ~112, (d) mechanical twins in the TMS-75 specimen (ε = 5.4 %), Beam Direction ~011.
creases at higher temperatures, which indicates that twinning becomes less likely to occur at higher temperatures. However, at
950 °C and 350 MPa, the critical size of a twin nucleus in compression reaches a minimum at n (number of atomic layers in a
nucleus) =4, and is smaller than the critical size of an SESF (n=2).
Therefore, a twin nucleus of four atomic layers is more likely to
form than an SESF. It should be clear from the above consideration that, in this study, the compressive stress favored twin nucleation and SESF formation even at 900 °C.
Ratio of Compressive to
Tensile Strain
15
CMSX-4
10
TMS-75
5
PWA1480
Mar-M247LC
0
700
To sum up, at 900 °C, the deformation mechanisms in tension and
compression are completely different; i.e., the dominant deformation mode in tension occurs by an a/2<110> dislocation motion;
however, the creep mechanism in compression was associated
with <112> slip and twin formation. At 750 °C, however, tension
and compression mechanisms were both associated with <112>
slip. The tension/compression asymmetry of CMSX-4 and TMS75 was larger at 900 °C than at 750 °C because the <112> viscous
slip does not occur under tensile stress at 900 °C.
750
800
850
900
950
Temperature, T (℃)
Fig. 11. Ratio of compressive to tensile creep strain at 20 h. MarM247LC is DS alloy.
enter the γ′ precipitate where the nucleation of a third Shockley
partial allows a super Shockley partial formation to trail an SISF;
i.e., crossing between the leading and trailing partial forms an
SISF [24]. This superlattice-stacking-fault formation mechanism
clearly indicates that the applied stress along [001] will favor
SISF formation in tension and SESF formation in compression.
As shown in Table III, the applied stress favored SISF formation
in tension and SESF formation in compression when testing along
[001]. At a given temperature and applied stress, the critical size
associated with an SESF twinning mechanism is smaller than that
for an SISF mechanism at intermediate and high temperatures,
suggesting that twinning is likely to form through an SESF
mechanism at both temperatures. The critical size generally in-
In PWA1480, at temperatures of 750 °C and 900 °C, the dominant
deformation mode in both tension and compression are a combination of a/2<110> dislocation bowing on the {111} plane in the
matrix and climbing along the γ/γ′ interface; that is to say, the
creep rates are controlled by the dislocation climbing process.
This is thought to be the reason why there is no difference in the
creep deformation mechanism for tensile and compressive stresses,
and why the asymmetry was smaller in PWA1480 than in the
other single-crystal superalloys.
440
Yield deformation mechanisms
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－
－－
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Acknowledgements
This research was partially supported by Grant-in-Aid for Scientific Research (C), 19560706, 2007. The authors would like to
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442