The effects of γ′ precipitate coarsening during isothermal aging and creep of the nickel-base superalloy IN-738

Materials Science and Engineering, 37 ( 1 9 7 9 ) 237 - 247

237

© Elsevier Sequoia S.A., Lausanne -- Printed in the Netherlands

The Effects of 7' Precipitate Coarsening During Isothermal Aging and Creep of the Nickel-Base Superalloy IN-738 R. A. STEVENS and P. E. J. FLEWITT

Central Electricity Generating Board, S.E. Region, Scientific Services Department, Scientific Services Centre, Canal Road, Gravesend, Kent DA12 2SW (Gt. Britain) (Received J u n e 12, 1978; in revised form August 24, 1978)

SUMMARY

The effects of isothermal aging in the temperature range 1023 - 1123 K on the 7' phase in the nickel-base superalloy IN-738 have been studied both with and without an applied creep stress. The initial heat treatment produces a bimodal distribution of small spheroidal precipitates (about 0.05 pm radius) and larger cuboidal precipitates (about 0.5 pm radius). During aging both types of precipitate obey diffusion-controlled coarsening kinetics (Ostwald ripening) but the cuboids coarsen at the expense of the spheroids until the latter eventually disappear. These microstructural changes are reflected by a decrease in room temperature hardness and a loss of creep resistance. It is demonstrated that at 1123 K the creep strength falls rapidly during about the first 5 × 106 s of aging. At 1023 K the coarsening of the 7' precipitates is much slower, but their importance is illustrated by the loss in creep resistance resulting from intermediate overaging heat treatments. However, it is shown that the creep strength at this temperature can subsequently be restored using a regenerative heat treatment. These results are discussed in the light of classical theories of creep of dispersion-hardened alloys.

1. INTRODUCTION

Nickel-base superalloys are widely used in service conditions which require high strength at elevated temperatures, for instance in the manufacture of gas turbine components. These alloys have generally evolved by the empirical addition of a variety of elements to the original Ni-Cr matrix. Some elements (e.g. cobalt, chromium, molybdenum and tungsten) form a solid solution with the face-centred

cubic matrix (7 phase) either to provide strength and/or oxidation resistance or to decrease the solubility of precipitate-forming elements. Other elements (e.g. aluminium and titanium) are added to form the ordered (L12) intermetallic 7' phase Ni3(A1, Ti) [1]. It is from the presence of a coherent dispersion of precipitates of this 7' phase that the superalloys derive much of their mechanical strength and creep resistance. The cast superalloy IN-738 is used as a blade material in the first row high pressure stage of both Avon and Olympus engines at gas turbine electrical power generation installations of the Central Electricity Generating Board [2]. During operation these blades experience a surface temperature of up to a maximum of about 1160 K, which is likely to cause considerable coarsening of the 7' precipitates generated by the original heat treatment. Indeed, previous studies of nickel-base superalloys [3 - 7] have confirmed that the 7' precipitates obey the usual (time) 1/3 coarsening law at elevated temperatures (above 0.5Tin ) consistent with the theories of diffusion-controlled growth at a constant volume fraction [8, 9]. Furthermore, since the 7' precipitates provide creep strength by acting as barriers to the movement of dislocations, then increase of precipitate size and spacing during service can adversely affect the creep strength and the overall service life. The dislocation models of the plastic deformation of metals at elevated temperatures have been reviewed by Blum [10]. Ansell and Weertman [11] developed a general theory which describes the steady state creep of dispersion-hardened alloys and Rowe and Freeman [12] demonstrated the validity of this theory for 7' strengthened nickel-base superalloys. Sherby and Burke [13] have also analysed this model and criticize it on the

238

grounds that in dispersion-hardened alloys the creep rate is frequently more strongly dependent upon the stress. However, there is general agreement for the particle size dependence such that when 7' precipitates of a constant total volume fraction coarsen an initial decrease in creep rate ~ is predicted since # ~ F - 2 , where F is the mean precipitate radius. This is due to the extra work which must be done by the dislocations in climbing over the precipitates. McLean [14] has suggested that, at low applied stress, dislocation movement results from alternate climb and glide processes, leading to a creep rate which is independent of precipitate size for a given volume fraction provided the precipitate spacing is not sufficient to permit dislocation bowing. However, it is generally agreed that the minimum creep rate occurs when the edge-to-edge spacing of precipitates lying in the dislocation slip plane is ~ = p b / o , where p is the shear modulus, b the Burgers vector and o the applied stress. Any further increase in precipitate size, and hence spacing, permits the dislocations to b o w between the precipitates, leading to a decrease in creep resistance as bowing becomes progressively easier (~ ~ F) [15]. The heat treatments usually recommended for nickel-base superalloys are intended primarily to produce volume fractions and size distributions of 7' precipitates which give optimum tensile strength and creep resistance [1]. In particular, the heat treatment developed for the cast superalloy IN-738 [16] produces a bimodal distribution of 7' precipitates (cuboids and spheroids) which has been shown empirically to give better creep resistance than a variety of other distributions. More recently Stevens and Flewitt [17] have demonstrated that such a heat treatment produces a total 7' weight fraction (which is approximately equivalent to the volume fraction [18] ) of about 0.45 which is divided into roughly equal proportions of 7' cuboids (0.5 1 p m diameter) and 7' spheroids (0.05 - 0.1 u m diameter). In this paper the growth kinetics during isothermal aging of the 7' precipitates produced by the recommended heat treatment in the cast superalloy IN-738 are investigated. The range of temperatures experienced by gas turbine blades during service is studied and the contribution of creep strain to these kinetics is examined. R o o m temperature hardness is

used to evaluate the effect of isothermal aging on the tensile strength and the role of 7' precipitate coarsening on the steady state creep rate is analysed and discussed in the light of existing theories.

2. EXPERIMENTAL TECHNIQUES

The IN-738 alloy was supplied by Deritend Precision Castings Ltd., in the form of ingots about 100 mm long and 20 mm in diameter at the top, tapering to 10 mm at the base. Tapering the ingot helped to reduce casting porosity to less than 0.05%, with no micropore exceeding 0.1 mm diameter. The nominal composition of this alloy is shown in Table 1. R o u n d section creep specimens were machined from each ingot (5.05 mm diameter X 27.5 mm gauge length;-~ in B.S.F. threads) and were given the standard heat treatment, under a vacuum of 1.3 mPa, of 7.2 X 10as at 1393 K followed by aging for 5.8 X 104s at 1118 K, cooling in air after each stage. High sensitivity constant load creep tests were performed in air using Mayes TC/20 creep machines at 1023 K and 1123 K with the temperature controlled to + 2 deg. Some tests were continued to rupture whilst others were interrupted at given fractions of the total creep life. The maximum testing time was 8.64 X 106s. Plain cylinders (about 10 mm diameter X 10 mm length) were also machined from the ingots, heat treated and isothermally aged in the absence of stress alongside the creep samples (at 1023 and 1123 K) or separately (at 1073 K). The creep samples and the isothermally aged plain cylinders were sectioned longitudinally and Vickers hardness measurements (HVso) were taken along their lengths. The size of the spheroidal 7' precipitates was determined using an extraction replication technique which has been described previously [19]. These replicas were examined in an AEI EM802 transmission electron microscope iTEM) at 80 kV and precipitate sizes were measured from fixed magnification prints used a Zeiss particle size analyser (TGZ3). The contribution of the large cuboidal 7' precipitates to the total 7' volume fraction after isothermal aging was determined by examination of positive carbon replicas. These replicas were taken from surfaces which had been very

239 TABLE 1 N o m i n a l c o m p o s i t i o n o f IN-738 a n d various phases t h e r e i n a s s u m i n g a 7 ' v o l u m e f r a c t i o n o f 0.45 w i t h t h e c o m p o s i t i o n given in S e c t i o n 3 Ni

C

Co

Cr

Mo

W

Ta

Nb

Al

Ti

B

Zr

Atomic weight IN-738 (wt.%) IN-738 (at.%) 7 ' (wt.%) ,yt (at.%)

58.7 61.42 59.33 72.38 69.15

12.0 0.17 0.79 -

58.9 8.50 8.17 4.57 4.35

52.0 16.00 17.47 1.58 1.70

95.9 1.75 1.02 0.26 0.15

183.9 2.60 0.79 2.07 0.63

180.9 1.75 0.57 3.71 1.15

92.9 0.90 0.57 1.70 1.03

27.0 3.40 7.15 6.24 12.95

47.9 3.40 4.03 7.52 8.80

10.8 0.01 0.06 -

91.2 0.10 0.06 -

Matrix a n d carbides (wt.%)

52.45

0.31

11.72

27.80

2.97

3.04

0.15

0.25

1.08

0.03

0.02

0.18

M a t r i x a n d carbides (at.%)

50.71

1.48

11.29

30.35

1.75

0.91

0.45

0.17

2.27

0.40

0.10

0.11

lightly etched, thereby reducing the Holmes effect [20], in a reagent containing 60% HC1 and 10% HNO3 in water, with 1.5 X 104 kg m -3 of both FeC1 a and C u C I 2. For examination of the creep-induced dislocation structures, discs 3 mm in diameter were spark-machined from both the gauge and thread of creep specimens. These discs were polished at 20 V, using a jetting technique, in an electrolyte containing 5% HC10s in ethyl alcohol maintained at about 263 K. Final thinning was carried out using a procedure described previously [17]. These foils were examined at 100 kV in the TEM, which was equipped with a +25 ° double-axis tilt stage.

3. R E S U L T S

Figure 1 shows the distribution of the two morphologies of ~/' precipitates in an IN-738 specimen after the standard heat treatment (Section 2). We shall refer to the large 7' precipitates as cuboids, although in this system they frequently assume a more degenerate shape. The " d i a m e t e r " of these cuboids was measured in terms of the diameter of a disc of equal projected area and was established as 0.88 gm. This is larger than t h a t previously observed in IN-738 {about 0.5 pm) [17] and is probably attributable to overaging in the long cooling time used to minimize porosity during casting. Small spheroidal ~/' precipitates are finely distributed between the cuboids, and these have a mean radius rso of 44.5 nm. The results of the creep rupture tests are summarized in Table 2 and the creep curves 1123 K are shown in Fig. 2. These curves frequently fail to show a pronounced primary

Fig. 1. C a r b o n surface replica s h o w i n g m i c r o s t r u c t u r e o f IN-738 p r o d u c e d b y initial h e a t t r e a t m e n t : C, c u b o i d s , m e a n " r a d i u s " , 0.44 p m ; S, s p h e r o i d s , m e a n radius, 4 4 . 5 n m .

£ .a 8

2 ~

266~21,~

l,

Time. t [810 s) Fig. 2. C o n s t a n t l o a d c r e e p r u p t u r e curves for IN-738 at 1 1 2 3 K, w i t h initial stresses (in MN m - 2 ) as m a r k e d .

creep range and are characterized by a continuously increasing creep rate, rather than the classical steady state or secondary range where strain is proportional to time. The increase in mean radius of the spheroidal 7' precipitates at the three aging tempera-

240 TABLE 2 IN-738 Creep rupture tests Temperature (K)

Initial stress (MN m -2)

1123 1123

266 244

1123 1123 1023 1023

217 194 457 429

Rupture life Extension (s x 106) (%) 1.74 I 2.77 ( 3.96 6.27 8.45 5.54 7.83

'E 0"20

(Ni0.922, C00.058, Cro.ol7, Moo.o02, W0.002)3 (Alo.51s, Ti0.352, Tao.o46, Nbo.o41, Wo.mT, Cro.o27 )

11231(

-=

--= 0 , , -

6.8 7.9 7.2 10.3 10.9 4.4 3.1

J

,.

0.,0

J+/+,0,,,

50

100

150

200

Fig. 3. Increase in 7' spheroid radius with (time) 1/3 during isothermal aging on ]N-738. • 1023 K, no stress; = 1073 K, no stress; o 1123 K, stress, 206 MN m-2; • 1023 K, stress (in MN m -2) as marked; o 1123 K, no stress; z~ 1123 K, stress (in MN m -2) as marked.

tures, 1023, 1073 and 1123 K, both with and without creep strain is shown as a function of (time) m in Fig. 3. At 1023 K the coarsening r a t e is very slow; the mean radius increases from 44.5 to 67 nm after 1.8 X 106s but shows little or no increase thereafter. Generally there is considerable scatter in the data points which is probably due to local segregation within these cast specimens [19]. However, at 1123 K a good least squares straight line fit is observed, which confirms diffusioncontrolled coarsening kinetics [6, 8, 9] : (77

--

Fsao)l/a = g t 113

tween the creep-deformed and the undeformed specimens. The chemical composition of the 7' phase in IN-738 has been proposed by Bieber and Mihalisin [16] to be

(1)

where t is the time and K is a temperaturedependent rate constant. Although the correlation with experimental data is n o t so good at the lower temperatures there is no reason to assume that similar coarsening kinetics are n o t obeyed. Thus within a given cluster or grouping the 7' spheroid coarsening is independent of the coexisting more widely distributed 7' cuboids. Furthermore, Fig. 3 shows no obvious difference in coarsening rate be-

On the basis of this composition, Table 1 shows the relative phase compositions for the volume fraction of 7' (0.45) produced by the initial heat treatment. For this total fraction virtually all the titanium in the alloy is required for the formation of the 7' phase; each 102g of IN-738 contains 3.40 g titanium, of which 3.38 g partitions to the 7'. Thus the maximum fraction of 7' of this composition that can be produced by heat treatment is a b o u t 0.45. For various aging times at 1023 K and 1123 K, the volume fraction of ~/' cuboids was determined from area fraction measurements on surface replicas (Fig. 4). Prolonged aging (about 8 X 106s) at 1023 K caused only a slight increase in volume fraction, but aging at 1123 K resulted in rapid coarsening of the cuboids at the expense of the spheroids such that dissolution of the spheroids was virtually complete after about 8 X 106 s. After this time the volume fraction of cuboids was about 0.45, so the total 7' volume fraction was virtually constant during aging. Hence the corresponding decrease in the volume fraction of the spheroids during aging could be deduced by difference (Fig. 4).

= 0'4 .9 o

P 0'3 LL

/

Cuboids

a

~

1023K

2 0'2

o.1

sph.

oid,

1123 K

0

I

I

I

I

I

I

1

2

3

4

5

6

O

I

7

~

I

~-.L

8

9 Time, t (s xl0 6}

Fig. 4. Change of relative fractions of 7 ' spheroids and cuboids during isothermal aging of IN-738: • 1023 K, 7 ' spheroids; o 1123 K, v' spheroids; • 1023 K, ~" cuboids; a 1123 K, 7 ' cuboids.

241

~60

440 I

••2•h

~ 1.20 c"

!~ ~ -

~ ~00 m

380

1073K-

360 1123K

340

0

I 1

I 2

I 3

I z,

I 5

I 6

I I I 7 8 9 Time, t (sxlO 6)

(a)

Fig. 5. Change of r o o m t e m p e r a t u r e hardness (HV30) during isothermal aging of IN-738: • 1023 K, no stress; * 1073 K, no stress; D 1123 K, stress, 206 MN m - 2 ; • 1023 K, stress (in MN m - 2 ) as m a r k e d ; o 1123 K, no stress; ~ 1123 K, stress (in MN m - 2 ) as marked.

Figure 5 shows the change in room temperature hardness associated with aging at 1023, 1073 and 1123 K. Again, aging at 1023 K has little effect, b u t after 7 X 108 s at 1073 K the hardness HV30 has decreased from the original value of 432 to 409. Aging at 1123 K in the absence of creep stress causes a very rapid decrease in hardness, AHV ~ 60, within the first 4 X 106 s followed by little change, suggesting that the 7' precipitates have achieved a critical size or spacing. Hardness measurements from the gauges of specimens crept at 1123 K show a reduction in the rate of hardness decrease, which becomes more pronounced as the stress is increased. This is unlikely to be due to differences in the 7' precipitate structures; creep strain generally has little effect on coarsening rate in nickel-base superalloys which contain 7' volume fractions up to 0.33 [21, 22] and Fig. 2 shows little difference, outside experimental error, between the size of 7' precipitates in stressed and unstressed specimens. Furthermore, at 1023 K specimens increase in hardness during creep. Thin foils were prepared from longitudinal sections of both the gauge and thread of a creep specimen which had ruptured in 1.74 X 10 s s under an initial stress of 266 MN m -2 at 1123 K. Figure 6(a) shows the bright field image of the two morphologies of 7' precipitates in the thread of the specimen. Within the

(b) Fig. 6. (a) L o w dislocation density in threaded end of creep sample aged for 1.74 x 106 s at 1123 K. (b) High dislocation density in gauge of same sample (stress, 266 MN m - 2 ) .

7 matrix there is an equilibrium number of dislocations (about 105 lines mm-2), which is confirmed by imaging under a range of twobeam diffraction conditions such that g ' b --/=0 for the Burgers vectors of all possible dislocations [ 2 3 ] . Figure 6(b) shows the corresponding microstructure within the gauge of this specimen. Consistent with the coarsening measurements in Fig. 3, the distribution and size of 7' spheroids and cuboids is similar to that within the unstressed thread. However, there is a significant increase in the number of dislocations (108 - 109 lines mm-~). Clearly, the 7' precipitates have formed obstacles to the dislocation m o v e m e n t which have been overcome by looping around the precipitates. The greater room temperature hardness of the

242 TABLE 3 Effects o f intermediate overaging heat t r e a t m e n t (7•2 x 105 s at 1123 K) on creep properties of IN-738 at 1023 K Time o f interruption (s × 106)

Creep rate before overaging 9 ~l(S - 1 × 1 0 - )

Minimum creep rate after overaging ~2(s - 1 x 10 - 9 )

-

4.20 a

-

0 1.80 2.75 3.60

3.47 6.06 5.86

e2/el

e2/el (predicted)

2.40

2.39

2.75 3.33 3.12

2.04 1.91 1.89

Total creep life (s x 106)

)

5.54

10.10 9.56 20.20 18.30

3.31 5.24 3.55 4.39

a M i n i m u m secondary creep rate.

gauges of the creep specimens when compared with the similarly aged unstressed specimens (Fig. 5) can therefore be simply attributed to this higher density of dislocations produced by the creep strain (Fig. 6(b)). To investigate the effect of the 7' precipitate distribution on the creep resistance of IN738, creep tests were performed at 1023 K at which temperature the precipitate-coarsening rate is low. An initial applied stress of 457 MN m -2 was used, which gives rupture in 5.4 X 106 s. Tests were interrupted at various stages in the secondary creep range (1.8 X 106 , 2.75 X 106 and 3.6 X 106 s) and the specimens were subjected to an overaging heat treatment for about 7 X 105 s at 1123 K. A further specimen was similarly treated before the creep test. The effect of such an overaging treatment is to cause an acceleration in the coarsening of the 7' precipitates. The specimens were then retested under the original conditions of temperature and load. Figure 7 (a) shows the subsequent creep curves, and the results are summarized in Table 3; at all points, the overaging treatment causes a very long primary creep range followed b y a greatly accelerated secondary creep rate. The rupture lives are also appreciably reduced, indicating that the creep resistance of IN-738 is strongly dependent on the 7' precipitate distribution at this temperature (1023 K). This was emphasized b y a further test subjected to the same overaging heat treatment after 2.75 X 106 s (Fig. 7(b)). On retesting, the creep rate increased from 3.1 X 10 -9 to 9.0 X 1 0 -9 s - i . After a further period of 6.7 X 105 s the test was again interrupted, and the specimen was subjected to the initial solution and aging heat

'A

'

'

B

5

c

;3 2

1

1

2

3

/* 5 Time, t (sxlO 6]

(a)

.E o m

0 Time, t ( s ~ O )

(b) Fig. 7. (a) E f f e c t of intermediate overaging heat treatm e n t (7.2 × 105 s at 1123 K) on creep strength of IN-738 at 1023 K: curve A, overaging t r e a t m e n t perf o r m e d before creep tests; curve B, overaging treatm e n t p e r f o r m e d after 1.8 × 106 s ; c u r v e C, after 2.7 × 106 s ; c u r v e D, after 3.6 × 106 s ; c u r v e E, no overaging treatment• (b) Creep curve of IN-738 specimen subjected to overaging heat t r e a t m e n t (as above} after 2.75 X 106 s and then reheat treated (7.2 × 105 s at 1393 K, 5.76 × 104 s at 1118 K) after a further 6.7 × 105 s.

243 treatment (Section 2) in an attempt to regenerate the original distribution of 7' precipitates. This had the effect of decreasing the creep rate to 3.3 × 10 -9 s-z, which is only slightly higher than the rate before the overaging treatment. The creep life after the reheat treatment was 3.5 × 106 s, giving a total life of 6.9 X 106 s; this compares with the total life of 5.4 × 106 s obtained from the original uninterrupted creep test.

4. Discussion The fact that the 7' spheroids approximately o b e y t z/a coarsening kinetics (Fig. 3) is in some ways surprising. If we consider the 7' phase precipitation as a whole, then classically the larger particles (cuboids) should coarsen at the expense of the smaller particles (spheroids), resulting in a decrease in the average size of the spheroids. The number of 7' cuboids remains essentially constant during isothermal aging, and as shown in Fig. 4 a constant total volume fraction is maintained by dissolution of the spheroids. However, the majority of 7' spheroids within the matrix are influenced by the diffusion field associated with neighbouring spheroids, rather than that of cuboids, and thus o b e y conventional coarsening kinetics as an independent precipitate system. Lifshitz and Slyozov [8] and Wagner [9] give the exact form of the growth equation for a distribution of spherical precipitates as 8 (D7V m Ce

Ft3 - F 0 a = ~

~

)t

(2)

where D is the composite coefficient of diffusion for the various elements, 7 the free energy of the precipitate/matrix interface, Vm the molar volume of the precipitate, Ce the concentration of 7'-forming elements in equilibrium with a precipitate of infinite radius, R the gas constant and T the absolute temperature. This theory assumes that the volume fraction of precipitate is small so that the spacing ~, between the precipitates is much greater than their radius. Ardell [6] has proposed a modification to eqn.-(2) to account for the finite precipitate volume fraction:

]¢ 8 [DTV m Ce) "F2--~g = m - ~ " ~ t = kmKat

(3)

where 10 > kr. > 1 for volume fractions f

such that 0.25 > f > 0. Thus an acceleration of coarsening is predicted for larger volume fractions of precipitate. However, no such effect has been observed in other nickel-base superalloys [5, 6] even for cases where the 7' volume fraction is about 0.6; in fact the phen o m e n o n has been positively identified in only one system. Although the volume fraction of 7' spheroids in IN-738 decreases monotonically during aging (Fig. 4), there is no corresponding decrease in coarsening rate (Fig. 3) as predicted b y eqn. (3). The temperature-dependent terms in eqn. (3) are D and Ce. The diffusion coefficient D has a dependence of the form D = Do e -O/RT where Q is the activation energy for the diffusion of the 7'-forming elements (principally aluminium and titanium) within the matrix. Stevens and Flewitt [17] have previously established that the equilibrium volume fraction of 7' phase is constant (0.45) over the temperature range studied (1023 1123 K). If it is assumed that the composition of the 7' phase does not change significantly during isothermal aging [ 1 8 ] , then Ce may also be assumed to be constant. The relation defining K can therefore be written as In (K 3 T) = constant -- Q / R T

(4,

The gradient of the straight line fit obtained by plotting In K a T versus 1 / T (Fig. 8) yields a -- 10:) ~e

I k-

10z

101

9~0

9!2

9j'/*

916 98 f l (~fl x 10-4)

Fig. 8. Plot of In K3T vs. l/T, where K is the rate constant for 7' spheroid coarsening in IN-738. Activation energy Q = 2.69 × 105 J mo1-1. value for the activation energy for 7' spheroid coarsening of Q = 2.69 × 105 J mol -z. This value compares favourably with the activation energies for diffusion of aluminium (2.70 X 105 J tool -z) and titanium (2.57 X 105 J

244

mol - t ) in nickel [24] and is in agreement with previously observed values for 7' coarsening in other nickel-base superalloys, e.g. 2.70 × 105 J mo1-1 for Udimet 700 [18]. The creep resistance of nickel-base superalloys depends critically on the size and spacing of the 7' precipitates. During service, first row high pressure gas turbine blades experience maximum surface temperatures in the range 1023 - 1160 K, and as shown in Figs. 3 and 4 coarsening and changes in the relative fraction of the two morphologies of 7' precipitates occurs at an appreciable rate in this temperature range. Hence a continuously changing creep resistance during service is to be expected. As described in Section 1, the Ansell-Weertman model [ 11 ] can be used to predict the change in the secondary creep rate with time. At low stresses, i.e. when o < p b/k, no pile-up or bowing of the dislocations at the precipitates is predicted and the secondary creep rate is governed by the rate at which dislocations can climb over the precipitates: ~s

-

rrobaD ' kTF2

(5)

where D' = D~ exp (-- Q / R T ) is the coefficient of vacancy or interstitial diffusion regulating the velocity of dislocation climb. Hence for a given stress and temperature, 7' precipitate coarsening at a constant volume fraction will initially reduce the creep rate since ~ r- - 2 When the applied stress exceeds the Orowan stress, o > pb/X, dislocations move past precipitates by bowing and this can result in "pinching o f f " loops around the precipitates. This process continues until the back stress exerted by the loops is sufficient to prevent further bowing. The creep rate is then determined by the rate at which the loop nearest the precipitate climbs over that precipitate and is annihilated; the other loops move inwards and a new loop is formed by "pinching o f f " an arrested dislocation. The secondary creep rate is then given by rfo4k2D '

is

p3kT F

(6)

The edge-to-edge spacing of precipitates within a dislocation slip plane is defined as [25] X = 0.82 F {(n/f) 1/2 -- 2}

(7)

where f is the volume fraction of precipitates. Thus for a constant volume fraction, precipi-

tate coarsening will increase the creep rate as ~ F. Although as described by Sherby and Burke [13] the stress
o.8L

0

I

I

I

I

I

I

2

3

4

5

I

I

I

I|

B 7 8 9 Time, t (sxlO 6)

Fig. 9. Variation with time of m e a n "/' spheroid radius Fs, mean 7' cuboid radius Fc and the respective mean precipitate spacings X s and k c within a dislocation slip plane during isothermal aging (1123 K) of IN-738.

the same time the mean spacing Xs of the spheroids (calculated from eqn. (7) after compensating for the volume occupied by the cuboids) shows a corresponding rapid increase and eventually approaches infinity as the

245 500 'z E

,

-~ ~oo ~S 300 200

~

D

100

......._ 1

2

3

C Time,t (s.lO6) 4

5

6

Fig. 10. Modes of dislocation movement during creep of IN-738 at 1123 K: A, bowing between spheroids and cuboids, ks < )'c; B, climb over spheroids, bowing between cuboids; C, climb over both spheroids and cuboids; D, bowing between spheroids and cuboids, X s > X¢.

spheroids disappear (about 8 × 106 s). Figure 9 shows that after a b o u t 5.5 × 106 s the mean spacing of spheroids and cuboids is equal, so after this time the cuboids become the principal barriers to dislocation movement. From the spheroid and cuboid spacings at various aging times, the associated Orowan stresses can be derived and the dominant creep mechanism determined as a function of the applied stress. Figure 10 maps the dislocation mechanisms which can operate at 1123 K. In region A the creep stress is sufficiently high to cause dislocations to b o w between b o t h spheroids and cuboids, whilst in region B such bowing is possible between cuboids but not between spheroids and dislocation climb over spheroids is the rate-controlling process. At very low stresses (region C) no dislocation bowing is possible and creep is regulated by dislocation climb over both spheroids and cuboids. After about 5.5. × 106 s (region D) it is only the cuboids which can provide a significant contribution to creep resistance since their spacing is less than that of the spheroids. However, it is likely that once the precipitates have achieved such a size different modes of dislocation impedance may operate. For example, as pointed out by Oblak and Kear [ 2 3 ] , since the 7' cuboids have an ability to store dislocations at the interface their actual contribution to creep resistance may be greater than that suggested by consideration of a simple dislocation bowing model. Nevertheless, it must be concluded that during creep at 1123 K a rapid loss in creep resistance is to be expected once the 7' spheroids have achieved a

size such that the applied stress exceeds the Orowan stress. When operating at 1123 K, the first row high pressure gas turbine blades are subject to longitudinal stresses of the order of 120 MN m -2. This stress exceeds the Orowan stress after about 2.2 × 106 s (600 h) (Fig. 10); further thermal exposure for times up to 5.5 × 106 s (about 1500 h) will then lead to a monotonic increase in creep rate as defined by eqn. (6). The effect of aging on the creep rate can be seen in Fig. 2; at all stresses studied the creep curves fail to show the classical secondary range (e ~ t) but instead show a continuously increasing creep rate throughout such that the onset of tertiary creep cannot be identified. It is worth pointing o u t that the observed variations in the extent of primary creep strain may well be the result of casting texture; these differences could then be explained by the orientation dependence observed by Oblak and Kear [ 2 3 ] . At 1023 K the coarsening rate of the 7' precipitates is much slower; thus the loss of creep resistance is also correspondingly slower. The contribution of 7' spheroids to overall creep resistance is lost after about 5.5 × 1 0 6 s at 1123 K, b u t extrapolation from Fig. 3 shows that the spheroids take about 8 X 107 s to coarsen to the same extent at 1023 K. The creep behaviour of IN-738 at 1023 K was studied using an initial applied stress of 457 M N m -2. Since the O r o w a n stress associated with the original 7' spheroid distribution is 446 M N m -2, a bowing model is envisaged at all points on the creep curve. The creep rate istherefore determined by the size and spacing of the 7' spheroids: o4 ~2

(8) This dependence is illustrated by the dramatic effect on the creep rate of an intermediate heat treatment (7 × 105 s at 1123 K) which overages and thus coarsens the 3" spheroids. On retesting these specimens, increased creep rates are observed (Fig. 7(a)). At each interruption point (before creep and at various points in the pseudo-secondary creep range), Figs. 3 and 4 provide the appropriate values of Fs and X and the corresponding values resuiting from the overaging treatment. Hence the ratio ~2/~1 of the secondary creep rate after the overaging treatment to that before can be predicted from eqn. (8). In doing this,

246

the change in stress due to the change in net section during primary creep was taken into account. In Table 3 these predicted ratios are compared with the experimentally determined values. The good agreement between the measured and predicted changes in creep rate emphasizes the large contribution the 7' spheroids make to the creep resistance of IN-738 at this temperature (1023 K). The creep strength will decrease only very slowly since the coarsening rate for 7' spheroids is low, but clearly it is important during service at this temperature to avoid operational temperature excursions if the strength level is to be preserved. These results point to the potential use of intermediate heat treatments for regenerating creep properties. As shown in Fig. 7(b), the resistance to deformation of specimens containing overaged 7' precipitate distributions can be recovered with considerable success if they are subjected to a heat treatment similar to that originally applied which regenerates a microstructure that approximates to the original. The successful extension of this procedure to isothermally crept specimens is likely to depend on the creep temperature; for instance, at 1123 K the 7' precipitates will overage again in a fairly short period of time (4 X 106 - 5 X 106 s), whereas at 1023 K overaging is more than an order of magnitude slower and correspondingly longer life extensions should be possible. However, a cautionary note has to be added since this procedure fails to consider any contribution from creep cavitation, which is frequently the most important life-limiting parameter [ 2 7 ] .

5. CONCLUSIONS

(1) The recommended heat treatment for IN-738 establishes a bimodal distribution of ~,' precipitates (cuboids and spheroids), the total volume fraction of which is the maxim u m permitted by the titanium content of the alloy. (2) The 7' spheroids obey conventional (time) 1/3 diffusion-controlled coarsening kinetics, with an activation energy of 2.69 X 105 J mo1-1. This occurs at a constant total volume fraction, the cuboids coarsening at the expense of and eventually to the exclusion of the spheroids.

(3) The 7' spheroid coarsening rate is essentially independent of the precipitate volume fraction and of applied creep stresses up to a b o u t 460 MN m - 2 . (4) Coarsening of the 7' precipitates during isothermal aging is reflected by a decrease in room temperature hardness and an associated loss in creep resistance. Changes in the pseudosecondary creep rate as the 7' precipitates coarsen can be explained by existing dislocation creep theories. 7' spheroids cease to offer resistance to dislocation bowing after about 5.5 X 108 s at 1123 K. However, at 1023 K their contribution is more important and creep life at this temperature is considerably reduced by a prior overaging heat treatment.

ACKNOWLEDGMENTS

This paper is published with the permission of the Director General, CEGB SE Region.

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