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

1Dept of Mechanical Engineering, Tokyo Metropolitan University; 1-1 Minami-Osawa, Hachioji, Tokyo 192-0397, JAPAN 2University of Cambridge, Department of Materials Science and Metallurgy;

Pembroke Street, Cambridge, CB2 3QZ, UK 3Dept of Metallurgy and Materials, The University of Birmingham; Edgbaston, Birmingham, B15 2TT, UK

Keywords: Single crystal, Creep, Twinning, Stacking fault, Tension, Compression

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 investi-gated 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 dis-tinctive 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 be-cause <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 climb-ing along the γ/γ′ interfaces.

Introduction 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 cross-slip 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

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]. Under thermo-mechanical fatigue (TMF), the number of cycles to failure of a single-crystal superalloy was found to decrease drasti-cally 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 relaxa-tion 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.

Experimental Procedure The investigation was carried out for a first generation SX super-alloy, a second generation SX superalloy, as well as a third gen-eration 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 intermedi-ate 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 crystallo-graphic orientation by the back-reflection Laue method, speci-mens were cut from a single as-aged bar using a spark cutter. The

Table I. Chemical composition of alloys (mass %) Co Cr Mo W Al Ti Ta Re C B Zr Hf Ni

PWA1480 5 10 - 4 5 1.5 12 - - - - - bal. CMSX-4 9 6.5 0.6 6 5.6 1 6.5 3 - - - 0.1 bal. TMS-75 12 3 2 6 6 - 6 5 - - - 0.1 bal.

MAR-M247LC 9.4 8 0.5 9.5 5.6 0.7 3.2 - 0.07 0.015 0.01 1.5 bal.

433

Table II. Heat treatment condition of alloys Solution heat treatment Aging treatment

PWA1480 1282°C/1 hr + 1287°C/2 hr + 1294°C/1 hr/GFC 1080°C/4 hr/AC+ 870°C/32 hr/AC

CMSX-4 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

1140°C/4 hr/GFC + 870°C/20 hr/GFC

TMS-75 1240°C/1 hr + 1280°C/2 hr + 1300°C/2 hr + 1320°C/8 hr/GFC

1150°C/4 hr/GFC + 870°C/20 hr/GFC

MAR-M247LC 1230°C/2hr/GFC 980°C/5 hr/GFC + 870°C/20 hr/GFC

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

tensile specimens and 10.0×3.0×3.0 (mm) in compressive speci-mens. 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 per-formed 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 JEM-3010 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].

800

900

1000

1100

1200

1300

0 200 400 600 800Temperature °C

Yie

ld S

tress

MP

a

TensionCompression

800

900

1000

1100

1200

1300

0 200 400 600 800Temperature °C

Yie

ld S

tress

MP

a

TensionCompression

(a)

(b)

(c)

800

900

1000

1100

1200

1300

0 200 400 600 800Temperature °C

Yie

ld S

tress

MP

a

TensionCompression

Fig. 1. Yield strengths of tensile and compressive tests under initial strain rates of

200 nm

g=1―

11

Fig. 2. TEM images of the deformed structure in CMSX4 specimen after yielding: (a) tension, ε=1.5% and (b) compres-sion, ε=4.3%.

200 nm

g=1_11

(b)

(a)

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.

434

g=1＿

11

g=1―

11

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.

200 nm

(a) (b)

a/3[21_

1_

]

S.F.

γ′

a/6[21_

1_

]

(b) (a)

g=1―

11

200 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[21

－

1－

] 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).

γ′ γ′ γ′twin

(d)

(111)

(c)

g=1―

11 200 nm 1 nm

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

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 com-pressive 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

435

a/2<110> dislocation pairs was the same (Fig. 2(a), (b)) for both tension and compression. Tensile and compressive yield micro-structures in PWA1480 are shown in Fig. 3. The tensile micro-structure showed matrix dislocations and many stacking faults in γ′ (Fig. 3(a)). The percentages of the numbers of superlattice in-trinsic 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 invisibil-ity criterion [9], the analysis of the Burgers vectors and line direc-tions [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 Bur-gers 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).

microstructure of the (110) foil is shown in image is an edge-on view of the fault using the

The compressive TEMFig. 4(c). This <110> zone axis. From the enlarged High Resolution Transmis-

sion Electron Microscopy (HRTEM) image (Fig. 4(d)) of the lin-ear 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. Creep strength 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 microstruc-tures of tensile and compressive interrupted creep tests in TMS-75

d PWA1480 are presented in Fig. 6. In TMS-75 the <112> rib-anbons which were composed of extended stacking faults and a/3<112> dislocations were observed both in tension and com-pression (Fig. 6(a), (b)). The dislocations in tension and compres-sion, 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 tan-gled 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 lamel-lae in the compressive creep microstructure were confirmed by {011} plane observation, as shown in Fig. 7. They were not ob-served in the bright-field image of PWA1480 (Fig. 7(c)). Tensile and compressive creep curves at 900 °C are presented in Fig. 8. Anisotropy between tension and compression was ob-served in CMSX-4 and TMS-75. As shown in Fig. 9, in all three alloys, it was observed that dislocations were concentrated into

SISF SESF

Tension 75.0 25.0 Compression 37.5 62.5

Table III. The proportion of SISF and SESF observed in ten-sile and compressive PWA1480 specimens at 750°C in (%).

(a) (b)

(c) (d)

0

5

10

15

(%)

Com

0 20 40 60 80 10Time, t /h

Cre

ep s

trai

n ε

Tension

pression

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0

1

2

3

0 50 100 150 200Time, t /h

Cre

ep s

trai

n ε

(%) Tension

Compression

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5

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Cre

ep

stra

in ε

(%)

Tension

Compression

0

10

20

30

0 50 100 150 200Time, t /h

Cre

ep s

trai

n ε

(%) Tension

Compression

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.

436

(a) (b)

the γ phase, and very few dislocations occurred in the γ′ phase. It is likely that the tensile creep deformation occurs by a combina-tion 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 disloca-tions 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 disloca-tions 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 secon-dary creep (Fig. 8(c), (d)). Moreover, deformation twins were observed in the compressive specimens in these two alloys (Fig. 10(d)).

g=202―

200 nm

g=1―

11

g=1―

11―

g=1

―

11―

200 nm

200 nm

200 nm

Fig. 6. Dislocation microstructure of interrupted creep specimens: (a)Direction (BD)～112; (b) 8 % compressive strain, the foilthe foil normal

(d)(c)

TMS75 at 1.4 % tensile strain, the foil normal [111], Beam normal [111], BD～112; (c) PWA1480 at 0.4 % tensile strain, TMS75 at 1.

[111], BD～111; and (d) PWA1480 at 0.8 % compressive strain, the foil normal [111], BD～112.

g=200 1 μm g=200

200 nm

g=200

200 nm Fig. 7. Compressive microstructure of interrupted creepTMS75 at 25 % strain

(a) (b)

specimens at 750 °C, the foil normal [011], Beam Direction (BD)～011: (a) ; (b) CMSX4 at 27 % strain; and (c) PWA1480 at 2 % strain.

(c)

437

Discussions Deformation mechanisms of creep The tension/compression asymmetry was distinctive in the creepstrengths of CMSX-4 and TMS-75. In CMSX-4 and TMS-75, thcompressive creep curve showed inferior strength compared to thtensile creep curve. Under compressive stresses, as shown in Fig. 7, the TEM micrograph of the [001] specimen showed microtwinlamellae extending over both the γ and γ′ phases. The results sug-gest that the asymmetry is also caused by mechanical twinninunder compressive stress. Tension/compression asymmetry ocreep strength in the [001] orientation is attributed to the diretional-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 she

stresses for <110>{111} and [11_

2](1_

11) slip systems. Twinning is not a dominant deformation mechanism in mewhich possess many possible slip systems; it generally occuwhen the slip systems are restricted or when something increas

e e

g f

c-

ar

tals rs es

nce the critical stress for twinning is lower than that for

interchange shuffles, and is not subjected to e need for a separate partial dislocation source on every (111)

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). Sithe 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

FCC metals occurs on (1_

11) planes in the [1_

12_

] direction. In [001]-compressive deformations, the resolved shear stress on the (

1_

11) plane acts in the twinning direction, and the formation of the mechanical twin results in weakness.

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 superal-loys, 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-

crotwin-formation model for the superalloys. After an a/6 [1_

12_

]

dislocation glides on the (1_

11) plane, a CSF (Complex Stacking

Fault) is produced behind the dislocation. The second a/6[1_

12_

]

dislocation cannot follow the first on the same (1_

11) plane be-cause it produces an energetically unfavorable structure. There-

fore, it is assumed that it glides on a neighboring (1_

11) glide plane.

The third dislocation follows on the next (1_

11) 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. 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 thethplane [14]. Therefore, the nucleation of the twin must be the rate-determining 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 mi-crotwin was then associated with the nature of superlattice stack-ing faults. In contrast with CMSX-4 and TMS-75, in PWA1480 and DS Mar-M247LC, the extent of creep asymmetry was small. The

0

1

2

3

0 100 200 300 400Time, t /h

Cre

ep

stra

in ε

(%) Tension

Compression

0

1

2

Tension

3

0 100 200 300 400

Cre

ep

stra

in ε

(%)

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Time, t /h

0

1

2

3

4

5

0 20 40 60 80 100Time, t /h

p st

rain

ε(%

)

Tension

Compression

Cre

eFig. 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.

(a)

(c)

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Cre

ep

stra

in ε

(%)

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(b)

(d)

438

crept PWA1480 specimen in this study did not show the mechani-cal 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 bow-ing 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 solid-solution strengthening of the γ′ phase resulting from lattice pa-rameter 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

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 rea-sonable to expect that the dislocation motion in the γ′ phase wouldbe restrained by the effects of tantalum. As a result, the dominandeformation mode of PWA1480 both in tension and compressionshould be a combination of both a/2<110> dislocation bowing onthe {111} plane in the matrix and the climbing process along the γ/γ′ interface; that is to say, the creep rates would be controlled bthe dislocation climb process. 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. Effect of temperature on creep deformation mechanisms

t

y

The ratio of compressive to tensile primary-creep strain as a func-tion of the test temperature is illustrated in Fig. 11. Ten-sion/compression asymmetry was more pronounced at 900 °Cthan at 750 °C in CMSX-4 and TMS-75. In general, the dominant deformation mechanism of creep in the temperature range abov850 °C is viscous slipping of a/2<110> dislocations and thclimbing process along the γ/γ′ interface, and the <110> slip does not show a clear primary creep strain [21, 22]. In the presenstudy, tensile creep specimens showed no primary creep stra(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

e e

t in

-

>

sic stacking fault loop,

er compressive

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.

g=200

(a)

g=200

400 nm

(b) g=200

400 nm

(c) g=200

400 nm

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 men-tioned above, the dominant deformation mechanism in the tem-perature range above 850 °C is the viscous slip of a/2<110> dislo-cations 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 ob-served 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 expla-nation for this may be that twinning nucleation is associated with the critical radius of the extrinsic and intrinas described below. 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. Undstress, 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

enter the γ′ precipitate where the nucleation of a third Shocklepartial allows a super Shockley partial formation to trail an SISF

., crossing between the leading and trailing partial forms an

y ;

i.e

SIAin[0s an SESF twinning mechanism is smaller than that r an SISF mechanis d high temperatures, ggesting that twinn through an SESF echanism at both temperatures. The critical size generally in-

be-, at

50 °C and 350 MPa, the critical size of a twin nucleus in com-n a

-

and ompression are completely different; i.e., the dominant deforma-

n;

win formation. At 750 °C, however, tension anisms were both associated with <112>

S-s

PWA1480, at temperatures of 750 °C and 900 °C, the dominant i-e e

ught to be the reason why there is no difference in the reep deformation mechanism for tensile and compressive stresses,

and why the asymmetry was smaller in PWA1480 than in the other single-crystal superalloys.

SISF [24]. This superlattice-stacking-fault formation mechanism clearly indicates that the applied stress along [001] will favor

SF formation in tension and SESF formation in compression. s shown in Table III, the applied stress favored SISF formation tension and SESF formation in compression when testing along 01]. At a given temperature and applied stress, the critical size sociated with a

fo m at intermediate aning is likely to formsu

m

creases at higher temperatures, which indicates that twinningcomes less likely to occur at higher temperatures. However9pression reaches a minimum at n (number of atomic layers inucleus) =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

rm than an SESF. It should be clear from the above considera-fotion that, in this study, the compressive stress favored twin nucleation and SESF formation even at 900 °C. To sum up, at 900 °C, the deformation mechanisms in tensionction mode in tension occurs by an a/2<110> dislocation motiohowever, the creep mechanism in compression was associatedwith <112> slip and tnd compression mecha

slip. The tension/compression asymmetry of CMSX-4 and TM75 was larger at 900 °C than at 750 °C because the <112> viscouslip does not occur under tensile stress at 900 °C. Indeformation mode in both tension and compression are a combnation of a/2<110> dislocation bowing on the {111} plane in thmatrix and climbing along the γ/γ′ interface; that is to say, thcreep rates are controlled by the dislocation climbing process.This is thoc

(c)

B

A

200 nm

g=1_1_1_

(b)

B

A

200 nm

g=1_ _11

400 nm

g=111 _ _

(a)

Fig. 10. Dislocation stru pted compressive crept specimins in 011.

ctures of interru ens: (a) PWA1480 (ε = 2.1 %), (b) CMSX-4 (ε = 1.1 %), and he TMS-75 specimen (ε = 5.4 %), Beam Direction ~(c) TMS-75 (ε = 1.2 %), Beam Direction ~112, (d) mechanical tw t

400 nm

(d) g=200

FM

ig. 11. Ratio of compressive to tensile creep strain at 20 h. Mar-247LC is DS alloy.

0700 750 800 850 900 950

Temperature, T (℃)

5

10ssiv

e t

oensi

le S

trai

n

CMSX-4

TMS-75

15

Ra

om

pre

tio o

f C

T PWA1480

Mar-M247LC

440

Yield deformation mechanisms For yield strength, tension/compression asymmetry has generally

as

－－ －

aids the cross-]

echa-le and compressive tests at 750 °C, ten-

m-"

gth. However, since the lattice image of a microtwin .

y metry of yield

ths

-

e hanical twin formation. Grain boundaries

the islocations.

t s

is the sm under tensile stress at the

temperature.

Acknowledgements This research was partially supported by Grant-in-Aid for Scien-tific Research (C), 19560706, 2007. The authors would like to gratefully acknowledge to Dr. Harada for the provision of materi-als.

References

1. M. Yamashita and K. Kakehi, “Tension/compression asym-metry in yield and creep strengths of Ni-based superalloy with a high amount of tantalum,” Scripta Materialia, 55 (2006), 139-142. 2. K. Kakehi, “Tension/compression asymmetry in creep be-havior of a Ni-based superalloy,” Scripta Materialia, 41 (1999), 461-465. 3. D.M. Shah and D.N. Duhl, “The effect of orientation, tem-perature and gamma prime size on the yield strength of a single crystal nickel base superalloy,” Superalloy 1984, M. Gell, C.S. Kortovich, R.H. Bricknell, W.B. Kent, and J.F. Radavich, eds. (Warrendale, PA: The Metallurgical Society of AIME, 1984), 105-114. 4. F.E. Heredia and D.P. Pope, “The tension/compression flow asymmetry in a high γ′ volume fraction nickel base alloy,” Acta Metallurgica, 34 (1986), 279-285. 5. R.V. Miner et al., “Orientation and temperature dependence of some mechanical properties of the single-crystal nickel-base superalloy Rene N4: PART III. Tension-compression anisotropy,” Metallurgical transactions A, 17A (1986), 507-512. 6. S. Takeuchi and E. Kuramoto, “Temperature and orientation dependence of the yield stress in Ni3Ga single crystals,” Acta Metallurgica, 21 (1973), 415-425. 7. C. Lall, S. Chin and D.P. Pope, “The orientation and tem-perature dependence of the yield stress of Ni3 (Al, Nb) single crystals,” Metallurgical Transactions A, 10A (1979), 1323-1332. 8. H. Zhou et al., “Thermomechanical fatigue behavior of the third-generation single-crystal superalloy TMS-75: Deformation structure,” Metallurgical and Materials Transactions A, 35A (2004), 1779-1787. 9. D.B. Williams and C.B. Carter, Transmission Electron Mi-croscopy III (New York, NY: Kluwer Academic/Plenum Publish-ers, 1996), 386-408. 10. W.W. Milligan, N. Jayaraman and R.C. Hill, “Low cycle fatigue of MAR-M 200 single crystals with a bimodal γ′ distribu-tion at 760 and 870°C,” Materials Science and Engineering, 82 (1986), 127-139. 11. G.B. Viswanathan et al., “Deformation mechanisms at in-termediate creep temperatures in Rene88 DT,” Superalloy 2004, K.A. Green, T.M. Pollock, H. Harada, T.E. Howson, R.C. Reed, J.J. Schirra, and S. Walston, eds., (Warrendale, PA: TMS, 2004), 173-178. 12. M. Kolbe, “The high temperature decrease of the critical resolved shear stress in nickel-base superalloys,” Materials Sci-ence and Engineering A, A319-321 (2001), 383-387. 13. A. Guimier and J.L. Strudel, Proceedings of the Second International Conference on the Strength of Metals & Alloys, ASM, 3 (1970), 1145-1149.

been explained in terms of the “core width effect”, which waccounted for by the following reaction.

－

a/6[211]+S.F.+a/6[112] →a/2[101]

Since constriction of the superpartial dislocationslip process, tension is stronger than compression in the [001orientation in which the extended superpartial dislocation retardscross-slip activity. In CMSX-4 and TMS-75, while the γ′ shearingby a/2<110> dislocations was the dominant deformation mnism in both tensisile/compressive yield strengths were comparable at every teperature in both alloys. From these results, the "core width effectwould not be a primary factor of tension/compression asymmetry in yield stren(Fig. 4(d)) and an a/3<112> dislocation on the {111} plane (Fig4(b)) were clearly observed in PWA1480, the asymmetry of ten-sile/compressive yield strengths is considered to be caused bmechanical twin formation. Therefore, the asymstrengths at 750 °C would be primarily due to the microtwin.

Conclusions

The tension/compression asymmetries of yield and creep strengof three-generations single-crystal superalloys and a DS superal-loy have been investigated; the following results were obtained. 1. Tension/compression asymmetry of both yield and creep

strengths in the [001] orientation are attributed to the direc-tional-formation characteristic of a mechanical twin; i.e., twin formation will occur only under the [001] compressive stress condition, and the formation of mechanical twins results in weakness.

2. Massive microtwins were observed at the same time as SISF/ SESF stacking fault pairs. The nucleation of the microtwin is associated with the nature of superlattice stacking faults.

3. In PWA1480 and DS Mar-M247LC, the extent of creep asymmetry was small. A high level of tantalum will prohibit theformation of an SISF/ SESF stacking fault pair and increase thcritical stress of mecalso inhibit {111}<112> slip operation and the formation ofmechanical twins.

4. The tension/compression asymmetry of the strengths was not observed when the dominant deformation mechanism was motion of the a/2<110> d

5. The extent of creep asymmetry of CMSX-4 and TMS-75 was found to be larger at 900 °C than at 750 °C. The reason is thathe <112> viscous slip does occur under compressive streseven at 900 °C; whereas, a/2<110> dislocation's motion dominant deformation mechani

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14. W.F. Hosford, The Mechanics of Crystals and Textured Polycrystals (Oxford Univ. Press, 1993), 169-72. 15. Y. Mishima, S. Ochiai and T. Suzuki, “Lattice parameters of Ni(γ), Ni3Al(γ′) and Ni3Ga(γ′) solid solutions with additions of transition and B-subgroup elements,” Acta Metallurgica, 33 (1985), 1161-1169. 16. B.H. Kear and D.P. Pope, Superalloys, Supercomposites and Superceramics, J.K. Tien, ed., (San Diego, Academic Press, 1989), 545. 17. F. Diologent and P. Caron, “On the creep behavior at 1033 K of new generation single-crystal superalloys,” Materials Sci-ence and Engineering A, A385 (2004), 245–257. 18. P. Veyssi`ere, G. Saada, Microscopy and plasticity of the L12 phase, in: F. Nabarro, M. Duesbery, eds., Dislocations in Solids, Vol. 10, Elsevier Sciences, 1996, 314-317. 19. C.M. Rae, K. Kakehi and R.C. Reed, Proc. of 7th Liege Conference of Materials for Advanced Power Engineering, Liege, Belgium, J. Lecomte-Beckers, M. Carton, F. Schubert and P.J. Ennis, eds. (Forschungszentrum Julich GmbH, 2002), 207-216. 20. R.A. Mackay and R.D. Maier, “The influence of orientation on the stress rupture properties of nickel-base superalloy single crystals,” Metallurgical Transactions A, 13A (1982), 1747-1754. 21. G.R. Leverant, B.H. Kear, and J.M. Oblak, “Creep of pre-cipitation-hardened nickel-base alloy single crystal at high tem-peratures,” Metallurgical Transactions, 4 (1973), 355-362. 22. N. Matan et al., “Creep of CMSX-4 superalloy single crys-tals: Effect of misorientation and temperature,” Acta materialia, 47 (1999), 1549-1563. 23. D.M. Knowles and Q.Z. Chen, “Superlattice stacking fault formation and twinning during creep in γ/γ′ single crystal superal-loy CMSX-4,” Materials Science and Engineering A, A340 (2003), 88-102. 24. B. Decamps, A.J. Morton and M. Condat, “On the mecha-nism of shear of γ′ precipitates by single (a/2)<110> dissociated matrix dislocations in Ni-based superalloys,” Philosophical Magazine A, 64 (1991), 641-668.

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