Intermediate temperature creep properties of gamma TiAl

Pergamon PII: S1359-6454(97)00063-3

INTERMEDIATE

TEMPERATURE GAMMA

Acta mater. Vol. 45, No. 9, pp. 3573-3585, 1997 0 1997 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved Printed in Great Britain 1359-6454/97 $17.00 + 0.00

CREEP TiAl

PROPERTIES

OF

MIN LU and K. J. HEMKER Department

of Mechanical

Engineering,

The Johns

(Received 4 December

Hopkins

University,

Baltimore,

MD 21218, U.S.A.

1996; accepted 27 January 1997)

Abstract-Compressive creep tests of single phase y Ti47AlSIMn2 have been conducted at stresses from 280 to 400 MPa in a temperature range of 5O(r6OO”C. The creep curves exhibited a primary region in which the creep rate decreased rapidly, a secondary region corresponding to a minimum creep rate, and an extended tertiary region in which the creep rate increased steadily with time. Unlike yielding, the creep behavior of this alloy was found to be normal; creep strength decreased with increasing temperature. Temperature change experiments and TEM observations have clearly elucidated the fact that a microstructural steady-state is not attained in the creep of this alloy. Instead, the overall creep performance is dominated by the tertiary creep that results from internal changes in the deformation microstructure. Although several dislocation and twinning mechanisms have been associated with the creep of TiAl, the multiplication and increased mobility of ordinary dislocations was found to play a dominant role in determining the creep behavior of this alloy. 0 1997 Acta Metallurgica Inc.

been much less conclusive. The measured values of range from 0.7 to 2.1 times the value of QdlRuslonr Qcreep Intermetallics of the titanium-aluminide system are see Table 1 and Refs [13-171. Attempts to relate the frequently cited as promising high temperature observed creep behavior to diffusion controlled structural materials, mainly due to their low density, processes appear tentative at best. Reported values of good oxidation resistance, relatively high modulus of the stress exponent (n) derived from minimum creep elasticity and yield stress [l]. Although the first rates are in the range of 3-8 [6, 9, 13-201. These engineering alloys will most likely have a lamellar numbers are in reasonable agreement with those structure of y and ~1~phases, with a titanium-rich measured in pure metals and solid solution alloys, composition, systematic investigations of singlewhich are generally associated with power law creep phase y are needed to develop a basic understanding and dislocation related processes. However, power of the high temperature deformation behavior of this law creep is a rather broad phenomenon, and further promising group of alloys. Several authors have observations are needed to identify the rate identified the deformation mechanisms associated controlling mechanisms. with the anomalous yielding behavior of polycrysIn the literature, creep of near y TiAl has been talline y TiAl [2%5], but relatively few mechanistic shown to lead to: dislocation activity [8, 9, 13, 171, studies have been forthcoming on the creep behavior twinning [21,22], grain boundary sliding [ 151, grain of this alloy. boundary migration [23] and dynamic recrystallizaA review of the literature suggests that it may be tion [7, 8, 11, 23, 241. Of these, dislocation motion, useful to separate studies associated with the creep twinning and recrystallization have been most behavior of single phase y alloys from those of duplex commonly reported, but their effects on creep are not or fully lamellar y-a2 two-phase alloys. As can be seen completely understood. The occurrence of dynamic in Table 1, the activation energy for creep (QcreeP)in recrystallization appears to be both alloy and the two-phase alloys generally lies between 300 and temperature dependent and may not be appropriate 350 kJ/mol [6-l 11. These values are higher, but not for a universal description of creep in TiAl. The much higher, than the value of the activation energy for the diffusion of Ti in ordered TiAl (Qdlffu_ observation of twinning, an athermal process, during time dependent high temperature deformation is sion = 291 kJ/mol) that has been reported by Mehrer completely unexpected and deserves further study. et al. [12]. However, the creep behavior of these Dislocation motion is the most common occurrence two-phase alloys has not yet been ascribed to a during power law creep, but the ordered structure of singular deformation process, and comparisons with TiAl raises interesting questions as to the type and single phase y TiAl are warranted. arrangement of dislocations in the deformation Attempts to measure and understand the temperamicrostructure. Both ordinary [20, 21, 251 and super ture dependence of creep in single phase y TiAl have 1. INTRODUCTION

3573

LU and HEMKER:

3514

CREEP OF GAMMA TiAl

[17,21,26] dislocations have been identified in crept specimens. Although the formation of weak subgrain walls has been reported in several studies [20, 271, the formation of a distinct subgrain substructure similar to those associated with the creep of pure metals and class II solid solution alloys has not been shown. The wide variations in the measured values of Qcreep and the lack of a well-defined dislocation substructure bring into question the existence of steady-state creep and the usefulness of traditional creep models for describing creep in single phase y TiAl. In the 1950s and 1960s Sherby and Dorn and their colleagues [28-311 showed that Qcreep is equal to QdlKuslon in a large number of metals and solid-solution alloys. The observation that high temperature creep in these metals and alloys is controlled by diffusion processes resulted in a phenomenological equation for the steady-state creep rate that is written as:

coefficient, 0 is where D,. is the lattice self-diffusion the stress, n is the stress exponent ( _ 5 for most pure metals), p is the shear modulus, b is the Burgers vector, and A is a constant of the order of 10’. The widespread acceptance of this description has led to the common practice of characterizing creep behavior by determining Qcreepand n from a series of creep curves. This phenomenological description of high temperature creep has been used to successfully predict the creep behavior of a relatively large number of pure metals and solid solution alloys, but there is evidence to suggest that this may not be true for a number of intermetallic alloys. The most basic assumption in the Sherby-Dorn work, constant dislocation structure during steady-state creep, is not necessarily true in ordered alloys. Hemker and Nix [32] have shown that, unlike pure metals, intermediate temperature steady-state creep does not exist in

Table

I. Reported

2. EXPERIMENTAL 2.1. Material

Takahashi

ef al. [I31

Takahashi

ef al. [14]

single phase T&Also

phase TtiAIz~ TiAl TiAl + W phase Ti4,A1S,

Kampe et al. [7] Wheeler ef ul. [8]

duplex TiS1A14, duplex Ti4XAl48NbEr~

Hayes and London [9] Hayes and Martin [15]

duplex TizIA14XNbl single phase Ti,oAls,Nb,

and Vasudevan

Bartels et al. [l I] Takahashi and Oikawa Ishikawa et al. [17]

[lo]

duplex

Tis2Alrs

duplex TiMA14RCrz [I61

and sample preparation

values for Q~,A~,,~~ and QcreePin TiAl

single duplex duplex single

Mehrer et al. [IZ] Martin er ul. [6]

PROCEDURES

A polycrystalline ingot of nominal composition Ti4,A15,MnZ (at.%) was provided for this study by P. L. Martin of Rockwell International Corporation and R. W. Hayes of Metals Technology (California, U.S.A.). The thermomechanical processing of the material, namely isothermal forging, resulted in an equiaxed fine grain structure with a grain size of approximately 10 /*m. A subsequent heat treatment (48 h at 1lOOC) was performed in order to (i) increase the grain size to approximately 150 pm; (ii) stabilize the initial structure prior to deformation; and (iii) obtain a low initial dislocation density to ensure that the dislocations observed after testing were produced during the deformation process (the latter was confirmed by TEM). After the heat

Alloy studied

Reference

Viswanathan

N&Al. Their TEM observations indicate that the minimum creep rate is actually related to a change in dislocation mechanisms; the minimum creep rate is associated with a transition from the formation of the Kear-Wilsdorf locks to the multiplication of mobile dislocations on the cube cross slip planes. This paper combines monotonic and transient mechanical testing with quantitative transmission electron microscopy (TEM) in an attempt to elucidate the fundamental mechanisms that control intermediate temperature creep in y TiAl. The results from this study are compared and contrasted with the fundamental description of creep in pure metals and solid-solution alloys. Where appropriate, parallels to previous work on N&Al have been highlighted, and a general understanding of creep in intermetallic alloys is being pursued. Comparisons between anomalous yielding and creep have also been made, and a description of creep that is in general agreement with the currently accepted view of yielding in y TiAl is presented.

single phase Ti-53.4%Al single phase Ti&(50 - 53)%Al

Q (kJ/mole) Qd+ru.,,.= 29 I = 300 Qcrcsp 350 360600 35& 600 340 300 419 317:;40 192S 560 278 350 270 460 340 300 600

Notes Ti diffusion

high low high low

in TiAl

D v (r (i

low rJ high 0 low T high T HT’d (u; 900°C HT’d @ 1350 + 900°C recovery GB recrystallization high d intermediate

0

LU and HEMKER:

CREEP OF GAMMA TiAl

treatment, X-ray diffraction experiments revealed some texture parallel to a (loll? direction [4]. Optical microscopy confirmed the presence of a single-phase microstructure with a fairly uniform grain size of approximately 150 pm. Compressive creep test specimens (3.5 mm x 3.5 mm x 10 mm) were cut by wire electrical discharge machining (EDM), with a loading axis along the (1011 texture direction. The ends of the EDM machined specimens were lapped to ensure that the ends of the specimen were parallel and all four sides of the specimen were polished using 400-600 grit grinding paper followed by diamond pastes (6, 1 and 0.1 /*m) to achieve a near mirror finish.

3575

0.8

0.6

0.4

0.2

0 0

50

Tie

2.2. Experimental

techniques

The temperatures used in this study (50&6OO”C) were all within the region of the yield strength anomaly and are approximately half of the melting temperature of this material. The applied stresses were set at 7&100% of the yield stress. These relatively high stresses were required to overcome the extremely high work hardening of this material at these intermediate temperatures. Monotonic compressive creep tests were conducted using an MTS machine. The applied stress of the monotonic creep tests was kept constant within 1 MPa by adjusting the load at each 0.25% increment of strain. Strain was measured with a clip gage extensometer and recorded with a precision of 4 x 10e5. The temperature of the specimen was measured by a ChromelAlumel (type K) thermocouple spot welded on the surface of the specimen and was controlled within f 1 K of the set value by an optical furnace. To eliminate the effects of environmental degradation, pure argon was used for all creep tests conducted in this study. Temperature change tests provide the opportunity to measure the kinetics of creep without changing the internal structure, provided that the temperature change is executed rapidly. Welding the thermocouple onto the specimen and heating with an optical furnace allowed for very rapid changes in temperature. During a 50°C temperature jump test, the temperature jumps from 550 to 600°C in less than 20 s, overshoots a maximum of 2°C for approximately 80 s, and stabilizes within 2 min. This time is relatively short compared to the time associated with creep and has been neglected in this study. TEM foils of the crept samples were cut 45” off the compression axis. These foils were mechanically thinned to 150 pm using 600 grit grinding paper and electropolished using a Struers Tenupol 2 with a solution of 6% perchloric acid, 35% 2tThe brackets used in this paper follow the convention introduced by Hug et al. [2], in which the first two indices are interchangable

but the last one is not.

150

100

ZCKI

(hrs)

Fig. I. Creep strain vs time for a creep test conducted at 500°C and at an applied stress of 280 MPa. The unusually high strain hardening rate of this alloy is compared with other metals and alloys in Table 2.

butoxyethanol and 59% methanol at 24 V and -30°C. TEM observations were carried out with a Philips EM420 transmission electron microscope operated at 120 kV. Conventional bright field and weak beam diffraction conditions were used to characterize the dislocation structure. In this work, special attention was given to the existence and density of: ordinary dislocations (bordlnary= l/2( 1 lo]), superdislocations (bsuper = (loll), faulted dipoles, mechanical twinning and recrystallization in the crept specimens. Evolution of the deformation microstructure was documented as a function of creep strain.

3. EXPERIMENTAL

3.1. Monotonic

RESULTS

creep behavior

All of the creep curves obtained in this study exhibited a “normal” primary creep response in which the creep rate was initially high but decreased quickly with increasing creep strain. The rapid exhaustion of the creep rate by strain hardening is evidenced in the creep curve presented in Fig. 1. This specimen, which was loaded to 70% of its yield stress at half of its melting temperature, exhibited 0.4% creep strain after 20 h (7 x lo4 s) and only 0.1% more after 1 week (6 x 10’s). These amounts are compared with similar values for cc-Ti, Al and Ni,Al in Table 2, and this table has been included here to demonstrate the unusually high degree of strain hardening that occurs during the primary creep of y-TiAl. The temperature dependence of primary creep in this alloy has been measured by conducting short term creep tests at 500, 562 and 597°C. The results of these tests are plotted in Fig. 2 and indicate that, unlike N&Al, y-TiAl has a normal temperature

3576

LU and HEMKER:

CREEP OF GAMMA TiAl

Table 2. The amount of creep strain obtained for Ti, Al, Ni,AI and TiAl Creep strain (24 h)

Material

~I~,

T/T*

Pure Al [40]

0.63

0.6

1%

Ti ASTM grade 3 [41]

0.40

0.4

5%

Ni,Al [32] TiAl (this studvj

0.70 0.71

0.5 0.5

0.4% 0.4%

dependence. This is reflected by the fact that the amount of primary creep increases with increasing temperature and does not show the anomalous strength dependence that has been reported for constant strain-rate experiments. The rapid exhaustion observed during primary creep necessitated the use of higher stresses for studies involving the entire creep curve. Applied stresses of the order of the yield stress (a,,%) were employed in a series of tests that were conducted at temperatures corresponding to the anomalous yielding behavior that has been recorded in constant strain-rate experiments. Creep curves that were obtained at 550, 515 and 597°C with an applied stress of 400 MPa are shown in Fig. 3(a). The overall shape of these creep curves is remarkably similar in nature to the intermediate temperature curves that have been reported for N&Al [32]. As with N&Al, the creep rate decreases with strain in the initial portion of the curve, reaches a minimum after the primary region and then gives way to an extended region where the creep rate continually increases with time. The overall amount of creep strain is observed to increase with temperature. To date, the majority of studies on creep in TiAl have characterized the creep behavior of TiAl by applying the constitutive laws that were derived for pure metals and class II solid solution alloys. In this approach, the stress exponent (n) is measured by

Creep strain (1 week) Rupture (@ > 16%) after 56 h Rupture after 100 h > 10% <0.6%

holding the temperature constant and measuring the minimum creep rate for tests conducted at a variety of stresses and Qcrcepis measured by holding the stress constant, assuming an Arrhenius relationship and plotting the log of the minimum creep rate vs reciprocal temperature. The creep rates from the creep curves in Fig. 3(a) are plotted in Fig. 3(b). In these curves, the duration of the region of minimum creep rate was always observed to be very limited as compared to the life of the entire creep curve. This short duration is inconsistent with the

(a) Ul”““““““““‘] d = 400 MPa

E .= i? a t

3

0.6

o = 280 MPa

__ 0

0

4x105

2x1@

6x105

Time (s) 5

10

15

20

25

Time (hts) Fig. 2. Strain vs time plots for a constant stress of 280 MPa. Note the normal primary creep behavior of T&A151Mn2; the primary creep strength decreases with increasing temperature.

Fig. 3. (a) Plots of creep strain vs time for creep tests conducted at an applied stress of 400 MPa and 550, 575 and 597°C. Normal temperature dependence, creep strength decreasing as temperature increases and an extended region of tertiary creep are evident in these curves. In (b) the creep-rate is plotted as a function of time, and the limited nature of secondary creep is evident.

LU and HEMKER:

CREEP

OF GAMMA

TiAl

3577

increase of 47°C resulted in a near two-fold increase in the creep rate; the creep rate was measured to be 1.7 x lo-* and 3.5 x lo-’ s-’ immediately before and after the change in temperature, respectively. Measuring the creep rate before and after the temperature change in Fig. 5 yields a creep activation energy of 380 + 5 kJ/molt, which was found to be significantly lower than was obtained by the traditional method that employs a series of monotonic experiments.

Q = 440 f 15kJhnol

3.3. Optical metallography 0.00114

0.00118

0.00122

lJ+WM

Fig. 4. A semi-logarithmic plot of Jog(&) vs the reciprocal temperature for a series of monotonic creep experiments.

Fitting an Arrhenius relationship to this data yields an activation energy for creep that is 440 + 15 kJ/mol.

notion of steady-state creep and will therefore be referred to as “secondary” creep in the remainder of this paper. Plotting the minimum creep rates of six different creep tests against reciprocal temperature in the normal manner, see Fig. 4, gives a measured value of the apparent activation energy for secondary creep of 440 + 15 kJ/mol. However, the extended region of continually increasing creep rates, which is commonly referred to as tertiary creep, dominates the creep performance of y-TiAl and appears to be more important than secondary creep. Tertiary creep is generally associated with the accumulation of damage, but the fact that this behavior has been measured during controlled atmosphere compressive creep testing precludes a number of the most common tertiary mechanisms (i.e. necking, void formation, environmental degradation) and suggests that a change in the internal microstructure of the specimens (i.e. grain growth, recrystallization, particle coarsening, dislocation or twin multiplication) is responsible for the accelerated behavior. 3.2.

Creep testing thermally etched the TiAl specimens used in this study, and optical micrographs of crept specimens show the grain sizes and the presence of linear features in the interior of the grains of the deformed material [see Fig. 6(a)]. Detailed observations of the specimens crept in this study failed to produce any clear evidence of the recrystallization, grain boundary sliding, grain boundary bulging or crack formation at the grain boundaries, that has been observed in earlier works [7, 8, 11, 15, 23, 241. This absence may be related to differences in the stress and temperature regimes employed in these studies. The linear features have been noted by other authors [15, 251 and interpreted to be the result of twinning. This interpretation is consistent with the observations made in the present study. The number of grains that contain these apparent twins was observed to increase with increasing creep strain. In order to obtain a better measure of the extent of twinning during creep, the fraction of grains containing twins was measured as a function of creep strain and plotted in Fig. 6(b). This figure shows that twinning activity increases throughout the life of the creep experiment and suggests that the number of grains with twins asymptotically approaches 100% with creep strain.

2.1L

I

55O’C

597°C &3.5x10-7

Temperature change tests

Rapidly executed temperature change tests have been used to experimentally measure Qcreepwithout having to assume the existence of a microstructural steady state during secondary creep. An example of a temperature change test is shown in Fig. 5. In this experiment, the specimen was loaded to 400 MPa and crept to 2.0% creep strain after 144 h at 550°C. Comparisons with other monotonic curves indicate that this specimen is in the secondary creep region. At this point, the temperature was increased from 550 to 597°C in the manner described in Section 2.2. This

tThe variation in this measurement has been determined using the data of a single test that was collected within the first hour after the temperature change.

-2.0

C?AT

:

bl.7x10-8

v

1.9 ~

-um

-2lm

0

zoca

4m

Time (s) Fig. 5. A transient temperature change experiment conducted after 144 h of creep at an applied stress of 400 MPa. Changing the temperature from 550 to 597°C resulted in a two-fold increase in the creep-rate and a value for Qcrcepof 380 f 5 kJ/mol. aAT indicates the thermal expansion of both the specimen and ceramic platens.

3578

CREEP

LU and HEMKER:

OF GAMMA

TiAl

)Y--

0 0.

80 -

6Q/

/

/

/

0

1’

40’ ;

/

/

/

0

/

Y

2o ;# I O;il 0

’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’: 20 25 15 5 10 Creep Strain (%)

09 Fig. 6. (a) An optical micrograph of a specimen stress of 350 MPa, and (b) the fraction of grains

3.4.

TEA4 observations

Monotonic creep tests have been interrupted at various stages along the creep curves and TEM foils

crept to 1.5% creep strain at 550°C under an applied containing twins plotted as a function of creep strain.

containing varying amounts of deformation have been prepared. These foils have been used to document the development and importance of various microstructural deformation mechanisms, as

LU and HEMKER:

CREEP

well as to confirm or disprove the existence of a microstructural steady state. The following observations were made on specimens crept to the primary, secondary and tertiary stages of creep behavior. 3.4.1. Primary creep. Figure 7(a) is a micrograph of the microstructure associated with a creep test that was interrupted during primary creep (6 = 0.6%). This specimen, which had been crept for 24 h at 597°C and an applied stress of 280 MPa, was found to contain ordinary dislocations, superdislocations and faulted dipoles. These three different defects can be seen in Fig. 7(a), where they are labeled OD, SD and FD, respectively. Observations of 10-1.5 areas indicated that, under these test conditions, the densities of superdislocations and faulted dipoles are significantly higher than the density of ordinary dislocations. The micrograph shown in Fig. 7(b) is another example of primary creep; this specimen was loaded to a higher stress (400 MPa) at the same temperature and achieved the same amount of creep strain in a shorter period (3 h). The major observed difference between this specimen and the previous one

6)

(b)

Fig. 7. TEM micrographs of the ordinary dislocations. superdislocations and faulted dipoles associated with primary creep. The specimen in (a) has been crept to 0.6% strain at 597°C under an applied stress of 280 MPa for 24 h. and the one in (b) has been crept to 0.6% creep strain at 597’C under an applied stress of 400 MPa.

OF GAMMA

TiAl

3579

is the fact that the one tested at the higher stress does not contain faulted dipoles. The number of superdislocations remained approximately the same, but the relative density of ordinary dislocations was found to be higher in the latter example. Primary creep specimens were examined at a variety of stresses and temperatures, and the disappearance of faulted dipoles and increased activity of ordinary dislocations appeared to be most closely related to the increase in the applied stress. In Fig. 7, the Burgers’ vectors of the ordinary dislocation (bordlnary = 1/2[ liO]) and superdislocations vector of the &per< = [ 10I]) and the translation faulted dipoles (R,,,, = 1/3[1ilJ) were found using conventional g ??b and g ??Rf analyses. Projections of the [I TO], [ IOI] and [Oii] directions are overlaid on Fig. 7(a), and tilting experiments were used to confirm that the ordinary dislocations and superdislocation are aligned along their screw orientations - (uordlnaiy= [ilO] and II,,~~~= [loll). Although both types of dislocations are aligned, the ordinary dislocations appear to bow out locally between pinning points while the superdislocations are much more linear and appear to be locked along the entire length of the screw segments. The faulted dipoles, which are aligned along the [Oii] direction but are not screw in character, appear to be the same as those characterized by Hug et al. [2] and Viguier and Hemker [33]. The ordinary dislocations, superdislocations and faulted dipoles that were associated with corroborate the primary creep in this study observations from yielding studies of the same alloy [4]. Very few twins have been observed in any of the primary creep specimens. 3.4.2. Secondary creep. At longer creep times, the region of minimum creep rate yielded the following changes in the microstructure. Faulted dipoles were not observed in any of the regions that have been studied. Ordinary dislocations and superdislocations remained in the microstructure. The density of aligned superdislocations remained fairly constant as compared with the density associated with primary creep. However, the density of ordinary dislocations increased significantly. In fact, the most striking observation made in this TEM study was that the amount and degree of bowing of ordinary dislocations is much greater than that associated with primary creep. Another pronounced change is that the number of twins increased significantly. Figure 8 shows the TEM micrographs of a region of a specimen after 168 h of creep at 550°C and at an applied stress of 400 MPa with a creep strain of 2.2%. Two families of aligned superdislocations and one family of bowed ordinary dislocations can be seen in Fig. 8(a). The bowing of the ordinary dislocations can be seen more clearly in Fig. 8(b). As in primary creep, the superdislocations are all aligned along the screw direction. A weak-beam image of one of these superdislocations is shown in Fig. 9. The highly linear structure and screw orientation suggest that the

3580

LU and HEMKER:

CREEP OF GAMMA TiAl

(4

Fig. 9. A weak-beam close-up of a superdislocatior 1 from Fig. 8. These superdislocations were found to be disscjciated screws and are believed to be in a locked configt uation similar to a KW or roof-like barrier.

type twins. The dislocations surrounding the twins were found to be ordinary dislocations; note their invisibility in Fig. 10(b) when the diffraction vet:tor is g = [002].

(b)

Fig. 8. TEM micrographs, (a) bright field (BF) image with g = [l 1 I] and (b) weak-beam (g, 5g) image with g = [Iii], associated with secondary creep of a specimen crept at 550-C and 400 MPa for 168 h. The microstructure for secondary creep contains long straight screw superdislocations and aligned but slightly bowed ordinary dislocations.

superdislocations have cross-slipped and dissociated into a non-planar “Kear-Wilsdorf’ or “roof-like” configuration. The ordinary dislocations still exhibit a slight preference for the screw orientation but the amount of bowing is progressed and is no longer localized. A closer look at these defects shows that the motion of ordinary dislocations is still inhibited by the cusps that have been associated with primary creep and yielding. However, these cusps or pinning points are more widely spaced and most of the ordinary dislocations can be seen to bow and glide past their pinning points. The increase in mobility of ordinary dislocations has led to dislocation tangles, but no clear evidence of subgrain formation has been observed. General observations of the crept specimens indicated that a number of grains contained twins as well as dislocations. Figures 10(a) and (b) show the TEM micrographs of another grain of the specimen deformed to 2.2% strain at 55O’C and 400 MPa. The twins that are visible have been identified to be { Ill’,

(b)

Fig. 10. Micrographs, (a) g = [lli] and (b) g =: [002], showing twins in the microstructure associated with secondary creep of a specimen crept at 550°C and 4010 MPa for 168 h. The twins were found to be lying on the ~{lll} plane and the invisibility of the dislocations in (b) in dicates that they are ordinary dislocations.

LU and HEMKER:

CREEP

OF GAMMA

TiAl

3581

points and are themselves a means for increasing the dislocation density. The density of microtwins was also observed to increase during tertiary creep, which is consistent with the fact that more linear features were found under optical microscopy at larger creep strains. The generation of ordinary dislocation has been shown to occur at twin intersections [34, 351, but this mechanism does not appear to account for the dramatic increase in ordinary dislocation density that occurs during tertiary creep. In this study, all grains were found to contain a large number of dislocations, but only a fraction of these grains contained microtwins. The creation of ordinary dislocations at microtwin intersections may occur, but it cannot be used to explain the dramatic increase in ordinary dislocation density because of the fact that a large amount of ordinary dislocations were found in grains without microtwins.

4. DISCUSSION

4.1. General

Fig. 11, The deformation microstructure associated with tertiary creep of a specimen crept at 597°C and 400 MPd for 70 h; (a) bright field (BF) image with g = [Iii] and (b) a (g. 5g) weak-beam image with g = [Iii].The degree of bowing and density of ordinary dislocations increase dramatically during tertiary creep.

Tertiary

3.4.3. micrographs nearly

creep.

Figures

of a specimen

70 h at 597’C,

that

at an applied

II(a)

and

had

been

stress

(b) show crept

for

of 400 MPa,

to a creep strain of - 23%. Compared to the microstructures observed in the primary and secthese micrographs illustrate a ondary regions, dramatic increase in the density and the degree of bowing of ordinary dislocations. The ordinary dislocations associated with tertiary creep do not exhibit a strongly preferred orientation. They were also observed to be relatively homogeneously distributed inside the grains; no subgrains were evident. The dramatic increase in dislocation density could be related to several different dislocation mechanisms. At various locations, these ordinary dislocations were observed to be spiraling around the pinning points in a manner that is consistent with a Z-mill dislocation source. Continued motion of ordinary dislocations was also found to occur by the pinching off of dislocation loops at the pinning points. The relatively high number of dislocation loops observed in Fig. 11 appear to be related to the expansion of the smaller loops formed at the pinning

creep

behavior

The creep curves obtained in this study exhibited the same basic features: primary creep where the creep rate decreases with time, secondary creep in which the creep rate goes through a minimum, and finally tertiary creep involving a continuous acceleration of the creep rate. Although the general shape of these curves is consistent with the majority of pure metals and solid solution alloys. the limited nature of primary and secondary creep and the increased significance of tertiary creep are not. For this reason, fundamental descriptions of creep that focus on steady-state creep behavior and disregard the importance of the transients must be expanded to include primary and especially tertiary creep. Recent studies by Hayes and Martin [15] and Maruyama et al. [23] have incorporated the importance of tertiary creep by using the time spent in tertiary creep phenomenological to calculate Qcreep and through descriptions such as the theta parameter [36]. Additional studies and a mechanistic description of tertiary creep are still needed. The creep tests conducted in this study all revealed a normal temperature dependence for both primary and overall creep behavior; both creep strengths decreased with increasing temperature. The effect of temperature on the creep properties of y TiAl is most clearly illustrated by the fact that the lives of the creep tests are significantly shorter at higher temperatures. This normal temperature dependence suggests that creep is controlled by a thermally activated process, and many investigators have used minimum creep rates and the empirical Sherby-Dorn relations to calculate an activation energy for this process. The results from this exercise have been fortuitous at best. The large scatter in the measured values of QcraP and the limited nature of secondary creep raise very

LU and HEMKER:

3582

CREEP OF GAMMA TiAl

serious questions about the existence of a constant dislocation substructure during secondary or “steady-state” creep in this alloy. 4.2. Absence of steady-state

creep in y TiAl

Transient temperature change creep experiments offer an alternative way of determining Qcreepwithout having to presume the formation of a stable steady-state microstructure. Traditional creep tests, which were conducted at an applied stress of 400 MPa in the temperature region where the yield strength anomaly is observed, resulted in a creep activation energy of 440 + 15 kJ/mol. However, the temperature change tests conducted under the same conditions yielded a creep activation energy of the order of 380 k 5 kJ/mol. The fact that the creep activation energy obtained from temperature change tests is significantly less than that obtained from the traditional monotonic constant stress tests suggests that there is a fundamental difference between the two experiments. This difference, the repeatability of the temperature change tests, the large reported scatter for Qcreepand the shape of the creep curves, all point to the absence of constant structure steady-state creep in 7 TiAl. The most direct method of determining the existence or absence of a microstructural steady state is to conduct TEM studies at various points along the creep curve. In the present study, the deformation microstructure has been characterized at a number of pre-selected creep strains. Unlike what is generally observed in the creep of pure metals and alloys, no evidence of subgrain formation was observed in y TiAl creep. In fact, the microstructure was observed to evolve throughout the creep test, which is in agreement with the work of Ishikawa and Oikawa [25]. Faulted dipoles, superdislocations and ordinary dislocations all contributed to primary creep. The density and degree of bowing of ordinary dislocations was observed to increase with increasing creep strain, and the density of twinning was also found to increase throughout the life of the creep test. These observations clearly indicate that intermediate temperature creep of single-phase y TiAl does not produce a constant microstructure, and we have concluded that a steady state does not exist in TiAl creep at intermediate temperatures. 4.3. Deformation

mechanisms

In the absence of steady-state creep, TEM observations of the creep microstructure have been used to characterize the deformation processes that govern the creep behavior of y TiAl. 4.3.1. Superdislocations. The generation and motion of superdislocations during creep has been observed in two forms; the formation of faulted dipoles during primary creep and the creation of rectilinear KW or roof type barriers can both be attributed to superdislocation motion. Inverse creep in N&Al has been correlated with the bowing out and

subsequent motion of superdislocations on the cube cross-slip plane [32]. Parallels between these two alloys and the similar shapes of their inverse creep curves suggest that the same mechanism may be operative in TiAl. However, the disappearance of faulted dipoles in the later stages of creep and the observation that the density and linear shape of locked screw superdislocations remains constant throughout the creep experiment are strong indications that superdislocation activity is limited to primary creep. The dramatic decrease in the creep rate that is observed during primary creep may also be explained in terms of the exhaustion of superdislocation motion. As in yielding, the motion of superdislocations during creep has been found to be inhibited by two competing processes; localized pinning and the formation of faulted dipoles and more global cross-slip and the formation of rectilinear KW or roof barriers. The formation of faulted dipoles at lower temperatures has been ascribed to the localized pinning of a superdislocation, bypassing of the pinned segment and the drawing out of a dipole whose energy is subsequently reduced by the passage of partial dislocations and the formation of an extrinsic stacking fault [2, 331. The fact that faulted dipoles are unstable at higher temperatures could be used to explain the normal temperature dependence of primary creep. However, the formation of KW or roof barriers has been observed to dominate at higher temperatures and longer times. Moreover, HREM [37] observations have shown these barriers to be dissociated in a highly nonplanar configuration, and as such, very formidable obstacles to further superdislocation motion. Further testimony to the immobile nature of cross-slip locked superdislocations can be seen in the large number of superdislocation pairs or dipoles that have been observed in this study. The observation of edge dislocation dipoles is generally related to the interaction of long range stress fields, but the screw dipoles observed in this study are more likely related to the efficient cross-slip locking of the screw segments of an expanding loop. The close proximity of the screw dipoles are a strong indication that super dislocations do not significantly contribute to tertiary creep. The macroscopic shape of the creep curves also suggest that superdislocation motion does not play an important role in tertiary creep. The temperature dependence of tertiary creep has been found to be much less in TiAl than it is in N&AI. The usually high temperature dependence of tertiary creep in N&Al has been associated with the Peierls-like motion of superdislocations on the cube cross-slip plane. Cross-slipped superpartial dislocations have been observed to be more widely dissociated in TiAl than in N&Al [37], and the Peierls stress should be proportionally higher. In this light, the more moderate temperature dependence exhibited by TiAl

LU and HEMKER:

CREEP

indicates that cube glide of superdislocations cannot be used to explain tertiary creep in this alloy. 4.3.2. Ordinary dislocations. The extended nature of tertiary creep that has been shown in this study appears to be related to the observed increased activity of ordinary dislocations. Although present during primary creep, the importance of ordinary dislocation motion has been found to take place after the motion of superdislocations has been exhausted. The dramatic increase in the density and mobility of ordinary dislocations shown in the TEM micrographs associated with creep are a clear indication that they play a dominant role in controlling the intermediate temperature creep strength of gamma TiAl. The observed morphology of these dislocations provides some clue as to the link between yielding and creep in this alloy. The observation that ordinary dislocations are pinned and aligned along the screw direction during primary creep suggests that their motion is initially inhibited by the same intrinsic process that has been used to describe the yield strength anomaly [38, 391. In these models the intrinsic pinning process involves a double cross-slip process by which the dislocation cross-slips to a secondary octahedral slip plane and then back to the primary slip system as is illustrated in Figs 12(a)(c). When the dislocation arrives in the configuration drawn in Fig. 12(c), the jogged segment

2 7

\

(b)

z

I/

OF GAMMA

TiAl

is unable to glide forward with the rest of the dislocation line and continued motion required one of several non-conservative processes: (1) the jog can zip in the lateral direction and allow one segment to grow at the expense of the other as shown in Fig. 12(d); (2) the jog can climb forward as in Fig. 12(e); (3) the leading dislocation segments can pinch off leaving a small loop in their wake as seen in Fig. 12(f); and (4) if the jog is sufficiently high the segments can spiral past each other and operate as a Z-mill source as drawn in Fig. 12(g). The increasingly bowed nature of the ordinary dislocation associated with the latter stages of creep suggests that one or more of these processes does occur if the dislocation is given a long enough time at a high enough temperature. Evidence for the last two mechanisms can be seen in Fig. 11. Moreover, the large number of dislocation loops visible in this figure can be attributed to the growth of the small loops that are created when the original dislocation pinches off. Both the expansion of dislocation loops and the spiraling of a Z-mill source will lead to an increase in the dislocation density. This increase is the key component for describing and modeling the extended nature of tertiary creep that controls the overall creep behavior of TiAI. The importance of diffusion in this process is not altogether clear. The best measured value for Qcreep

(111)

%

\

(4 /

3583

motion

Fig. 12. Schematic of ordinary dislocation motion in TiAl. A local segment of a screw dislocation (a), which cross-slips to a secondary octahedral plane (b) and then back onto the primary slip plane (c), creates two edge segment jogs that cannot glide with the rest of the dislocation. In this configuration, the dislocation can move forward by: (d) lateral motion of the jog and the expansion of one loop with respect to its neighbor, (e) climb of the edge component jog, (f) pinching off and leaving behind small dislocation loops or debris, or (e) the by-passing and spiraling of the dislocation segments on parallel slip planes.

LU and HEMKER:

3584

CREEP

[12] Ti in TiAl (Qurwon = 291 kJ/mol). The activation energy for the diffusion of Al in TiAl has not been measured but is believed to be somewhat higher than for Ti. Since dislocation climb would require the diffusion of both species, the activation energy for creep may be related to the climb of the jogs that form at pinning points. However, the importance of thermal activation on the motion of ordinary dislocations cannot be neglected. More detailed observations of the dislocation structure and comparison with yielding and transient experiments will be needed to fully describe this relation. 4.3.3. TiGzning. It is widely accepted that twinning is an athermal process, and the increased activity of twinning during creep cannot be attributed to thermal activation. In this light, the observation that the density of twins increases during creep of TiAl is rather surprising. However, a number of studies have shown that mechanical twinning does occur during creep deformation of TiAl, both in single-phase and two-phase alloys, see for examples Refs [21, 221. Linear features resembling twins, similar to those shown in Fig. 6, were also reported for creep in single phase y TiAl [15, 251. When considering the importance of twins, it is important to emphasize the fact that the dislocations are not replaced by twins. In fact, the density of dislocations also increases with increasing creep strain. Twin intersections have been observed to produce ordinary dislocations [34, 351, but the fact that a high density of ordinary dislocations exists in grains that do not contain twins indicates that dislocation activity is not dependent on twinning. Microstructural evidence found in this study suggests that twins may be the result of stress concentrations that develop during the accumulation of creep strain in polycrystalline specimens. Grains with unfavorable orientations for dislocation motion might be more favorable sites for the formation of twins. Moreover, geometric constraints may lead to compatibility stresses that promote twinning. This idea is supported by the fact that the number of grains that undergo twinning varies with the amount of creep strain, see Fig. 6 and with the yield stress [4]. (380 kJ/mol) activation

is

energy

higher for

than

the

the

diffusion

OF GAMMA

reported

of

5. CONCLUSIONS Creep behavior of single phase y Ti,,AIS,Mn, has been investigated in this study. Results of monotonic and transient experiments have been correlated with TEM observations and the following conclusions have been drawn. (1) The intermediate temperature creep response of y TiAl can be separated into three stages: primary, secondary and tertiary creep. (2) The extremely high amount of strain hardening that occurs during primary creep can be

(3)

(4)

(5)

(6)

TiAl

explained by the exhaustion of superdislocations and the intrinsic pinning of ordinary dislocations. The minimum creep rate is reached during secondary creep, but this stage is short lived and is not associated with a microstructural steady state. The overall creep performance of this alloy is dominated by tertiary creep, which is a direct result of the increased activity of ordinary dislocations. In the absence of steady-state creep, measured values of the activation energy have been obtained using temperature change experiments. The measured value of Qcreepcan be ascribed to either diffusion-related climb or the thermally activated motion of ordinary dislocations. Unlike yielding, the creep strength of TiAl exhibits a normal temperature dependence. However, the intrinsic pinning and subsequent motion of ordinary dislocation motion is the controlling process in both cases, and creep and yielding are inter-related processes.

Acknowledgements-This work was supported by the Mechanics and Materials Program in the Division of Mechanical and Structural Systems of the Engineering Directorate of the National Science Foundation under Grant No. CMS-9409538.

REFERENCES I. Kim,

Y. W. and Dimiduk, D. M., JOM, 1991, 8, 40. A., Phil. Mug. 2. Hug, G., Loiseau, A. and Lasalmonie, A, 1986, 54, 47. 3. Court, S. A., Vasudevan, V. K. and Fraser, H. L., Phil. Msg. A, 1990, 61, 141. 4. Viguier, B., Hemker, K. J., Bonneville, J., Louchet, F. and Martin, J-L., Phil. Mug. A., 1995, 71, 1295. 5. Farenc, S. and Couret, A., in High-Temperature Ordered Intermetallic Alloys V, ed. I. Baker, R. Daviola, J. D. Whittenberger and M. H. Yoo. Mater. Res. Sot. Proc. 288, Pittsburgh, PA, 1993, p. 965. M. G. and Lipsitt, H. A., 6. Martin, P. L., Mendiratta, Metall. Trans. A, 1983, 14, 2170. I. Kampe, S. L., Bryant, J. D. and Christodoulou, L., Metall. Trans. A, 1991, 22, 447. 8. Wheeler, D. A., London, B. and Larsen, Jr D. E.. Scripta metall. mater., 1992, 26, 939. B., Acta metall. mater.. 9. Hayes, R. W. and London, 1992, 40, 2167. 10. Viswanathan, G. B. and Vasudevan, V. K., in Higher-Temperature Ordered Intermetallic Alloys V, ed. I. Baker, R. Daviola, J. D. Whittenberger and M. H. Yoo. Mater. Res. Sot. Proc. 288, Pittsburgh, PA, 1993, p. 1155. 11. Bartels, A., Seeger, J. and Mecking, H., in HighTemperature Ordered Intermetallic Alloys V, ed- I. Baker, R. Daviola, J. D. Whittenberaer and M. H. Yoo. Mater. Res. Sot. Proc. 288, Pittsburgh, PA, 1993, p. 1179. 12. Mehrer, H., Sprengel, W. and Denkinger, M., Diffusion in Ordered Alloys. TMS, Warrendale, PA, 1993, p. 51. 13. Takahashi, T., Nagai, H. and Oikawa, H., Mater. Sci. Engng A, 1989, 114, 13.

LU and HEMKER:

CREEP

14. Takahashi.

T.. Nagai, H. and Oikawa, H., Mcrter. 1989. 30, 1044. R. W. and Martin, P. L., Acta metall. mater.,

Trans., JIM,

15. Hayes,

1995, 43, 2761.

16. Takahashi, Pruc.,

T. and Oikawa,

H.. Marer.

Rex. Sot. Symp.

1989. 133, 699.

17. Ishikawa,

18. 19. 20. 21. 22.

Y., Maruyama, K. and Oikawa. H., in Intermetallics. R. Darolia, Structural ed. J. J. Lewandowski, C. T. Liu, P. L. Martin, D. B. Miracle and M. V. Nathal. The Minerals, Metals & Materials Society. Warrendale, PA. 1993, p. 345. Takahashi, T., Nagai, H. and Oikawa, H.. Marer. Sri. Engng A, 1990, 128, 195. Nagai, H., Takahashi. T. and Oikawa. H.. J. Mater. Sci.. 1990. 25, 629. Oikawa, H.. Mater. Sci. Engng A, 1992, 153, 427. Loiseau, A. and Lasalmonie, A., Mater. Sci. Engrrg A, 1984. 67, 163. Jin, 2. and Bieler, T. R., Scripta mctall. mater., 1992,

27, 1301. K., Takahashi, T. and Oikawa, H., Mater. 23. Maruyama, Sci. Engng A, 1992, 153, 433. 24. Mitao, S., Kohsaka. Y. and Ouchi, C., Nippon Kinzoku 1986, 50, 840. Gakkai-shi, 25. Ishikawa, Y. and Oikawa. H., Mater. Trans., JIM,

1994, 35, 336. 26. Huang, J. S. and 1991, 25, 1901.

OF GAMMA

Y-W..

Scripta

metall.

mater.,

3585

27. Lipsitt,

H. A., Schechtman, D. and Schafrik, R. E., Trans. A, 1975, 6, 1991. 28. Sherbv. 0. D.. Orr. R. L. and Dorn. J. E.. Trans. AIMI? 1954, 200, 71. 29. Orr, R. L., Sherby, 0. D. and Dorn. J. E., Trans. ASM. Metall.

1954, 46, 113. 30. Sherby, 0. D.. Actcr metall.,

31. Sherby,

1962, 10, 135. D. and Miller. A. K., ASME J. Ma/u. 1979, 101, 387. K. J. and Nix, W. D., Actcr metall. mater.,

0.

Tech&., 32. Hemker, 1991, 39, 1901. 33. Viguier, B. and Hemker, K. J., Phil. Mag. A, 1996, 73, 575. 34. Wardle, S., Phan, I. and Hug, G., Phil. Mug. A, 1993, 68, 471. 35. Morris, M. A., Phil. Mug. A, 1994, 69, 129. 36. Evans, R. W. and Wilshire, B., Creep qf Metals and Alloys. The Institute of Metals, London, 1985, Chapter

6. 37. Hemker, Engng A.

38. Louchet,

K. J., Viguier, B. and Mills, M. J., Muter. Sci. 1993, 164. 163. F. and Viguier, B.. Phil. Mag. A, 1995, 71,

1313. 39. Sriram,

S., Dimiduk, D. M., Hazzledine, P. M. and Vasudevan, V. K., Phil. Msg. A, 1997, in press. 40. Servi, I. S. and Grant, N. J., Trans. AIME, 1951, 191, 909.

41. Metal Kim.

TiAl

Metals,

p. 537.

Vol. 1, Properties and Selection of 8th edn, American Society for Metals. 1961.

Handbook,