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Creep and discontinuous precipitation in a Ti3Al-Nb alloy at 923°K
Scripta M E T A L L U R G I C A et M A T E R I A L I A
Vol.
CREEP
24, pp. 1209-121¢, 1990 Printed in the U.S.A.
AND DISCONTINUOUS
P e r g a m o n Press plc All rights r e s e r v e d
PRECIPITATION
IN
A Ti3A1-Nb ALLOY AT 923°K
GE Corporate
R. G. Rowe and M. F. X. Gigliotti Research and Development, Schenectady,
NY 12301
and GE Aircraft
B. J. Marquardt Engines, Cincinnati,
(Received
April
17,
OH 45215
1990)
Introduction Creep r e s i s t a n c e is c r i t i c a l for the high t e m p e r a t u r e a p p l i c a t i o n s in which Ti3Al-base alloys are envisioned. The creep rate below 0.2% is g e n e r a l l y high for these alloys b e c a u s e t r a n s i e n t or p r i m a r y creep can e x t e n d to szrains of several tenths p e r c e n t (I) . This is much g r e a t e r than the amount cf primary creep o b s e r v e d in nickel-base superalloys of comparable tensile strength. (2) The alloy Ti3AI-12.5Nb has a m i c r o s t r u ct u r e which consists of alpha-2 (D0!~) and beta (bcc) phases which are formed by transformation from an e l e v a t e d temperature beta phase field. O r t h o r h o m b i c Ti2AINb has also been i d e n t i f i e d in this alloy (3). T r a n s f o r m a t i o n from beta to alpha-2 plus beta results in laths of a!pha-2 phase which vary in size a c c o r d i n g to the c o o l i n g rate t h r o u g h the t r a n s f o r m a t i o n temperature. We have shown that m i c r o s t r u c t u r e s t r o n g l y a f f e c t s %he creep behavior of this alloy. (i) In this paper, we r e p o r t our o b s e r v a t i o n of discontinuous precipitation at grain boundaries during 923°K creep testing. Experimental
Procedure
The alloy Ti3AI-12.5Nb at.% was prepared by n o n - c o n s u m a b l e arc mei~ing. The compositions of the two heats are t a b u l a t e d in Table I. The alloys were cast into copper chill molds and extruded at 1313°K-1355°K to 12:1 reduction in area. Samples were heat treated in argon and m a c h i n e d into b u t t o n h e a d tensile specimens with a gage section 2 mm d i a m e t e r by 11.7 mm long. Three m i c r o s t r u c ~ u r e s were studied. These were: alpha-2 plus beta with 60% prior alpha-2, and zransformed beta m i c r o s t r u c t u r e s with two different t r a n s f o r m a t i o n lath d i m e n s i c n s . Heat treatment conditions are shown in Table 2. TABLE SAMPLE 932,934 938
#
i.
Alloy Compositions Ti Bal. Bal.
(at.%)
A1
Nb
24.7 26.5
12.9 12.6
Creep samples were d e a d weight l o a d e d at a s t r e s s of 175 M2a and a temperature of 923°K in a r g o n gas. Sample extension was m e a s u r e d by an e x t e n s o m e t e r on the b u t t o n h e a d of the sample. E x t e n s i o n was d i g i t a l l y recorded c o n t i n u o u s l y f r o m the onset of loading. The e x t e n s i o n of sample 934-B3 was recorded manually.
1209 0 0 3 6 - 9 7 4 8 / 9 0 $3.00 + .00 C o p y r i g h t (c) 1990 P e r g a m o n Press
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1210
CREEP
TABLE 2.
IN Ti3AI-Nb
Vol.
24, No.
7
Heat Treatment Conditions Heat Treatment
Grain Size (~m)
Sample #
Notation
932-AB
1313/He
1313°K/lhr/He+1033°K/lhr
6-8
938-OS,OT
1408/SQ
1408°K/lhr/Salt Quench to 1088°K,+1033°K/lhr
300
934-B3
1448/He
1448°K/lhr/He+1033°K/lhr
750
Experimental Results The m i n i m u m creep rate and total primary creep strain of samples with three different m i c r o s t r u c t u r e s are tabulated in Table 3. Figure l(a) shows the creep curve of alpha-2 plus beta heat treated sample 932-AB. The m i n i m u m creep rate extension is shown superimposed as a dotted line. The total primary creep strain of this sample was 2.0%. The creep curves for beta heat treated samples 934-B3 (helium cooled) and 938-OS (salt quenched) are shown in Figures l(b)and l(c), respectively. The total primary creep strain for salt quenched sample 938-OS was 1.5%. Sample 934-B3 which was heat treated at 1448°K and helium cooled had only 0.3% total primary creep strain. TABLE 3.
Creep Test Results,
923°K, 175 MPa
Primary Creep Strain (%)
Minimum Creep Rate (sec -I)
Time to Failure or Unload (hrs)
934-AB (1313/He)
2.0
3.9E-7
68
934-B3 (1448/He)
0.3
2.4E-9
236
(unloaded)
938-OS (1408/SQ)
1.5
8.7E-9
811
(unloaded)
Sample Number
The microstructures of these samples are shown in Figures 2-5. Figures 2(a) and 2(b) show the e q u i a x e d grain structure of sample 932-AB. The grain size was 6-8 ~m. This sample c o n s i s t e d of e q u i a x e d prior alpha-2 grains s u r r o u n d e d by transformed alpha-2 plus beta. The beta phase showed up as dark areas between light alpha-2 laths in this b a c k s c a t t e r e d electron image. The fraction of prior alpha phase was approximately 60%. The gage section had a different distribution of beta phase. Figures 3(a) and 3(b) show this. The gage section of this sample consisted of e q u i a x e d alpha-2 grains s u r r o u n d e d by adjacent alpha-2 plus beta regions in which the density contrast of the two phases was reduced. Grain boundaries p e r p e n d i c u l a r to the load axis e x h i b i t e d d i s c o n t i n u o u s high and low d e n s i t y p h a s e s w h i c h had a g r e a t e r d e n s i t y d i f f e r e n c e and hence g r e a t e r compositional difference than the transformed alpha-2 plus beta regions. This is shown in Figure 3(b) . The b a c k s c a t t e r e d e l e c t r o n image d e n s i t y for the dark discontinuous phase was darker than the initial beta phase image density and the image density of the light discontinuous phase was less than that of the initial alpha-2 p h a s e image density. Rupture, which o c c u r r e d by s e p a r a t i o n at the interfaces between the discontinuous phases and the prior alpha-2 grains, limited the total thermal exposure of this sample to 68 hours. It is r e a s o n a b l e to assume that the d i s c o n t i n u o u s phases are alpha-2 and beta phases of c o m p o s i t i o n s c o r r e s p o n d i n g to the creep t e m p e r a t u r e , 923°K. Banerjee (4) has shown that the niobium content in the beta phase of Ti-24AI-IINb changed rapidly from 1373°K to 1315°K. Our results suggest that the composition
Vol.
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7
CREEP
IN Ti3AI-Nb
1211
of the alpha-2 phase must also change below this temperature. It has been shown that o r t h o r h o m b i c Ti2AINb forms in this alloy (3), and it will take quantitative a n a l y s i s to unambiguously establish the identity of the discontinuous phases Figure 4 shows the m i c r o s t r u c t u r e of the gage section of beta heat treated and salt q u e n c h e d sample 938-OT (1448/SQ) . This sample had a two phase W i d m a n s t ~ t t e n m i c r o s t r u c t u r e with coarse 1 X 8 um laths interspersed with finer 0.5 x 3 um laths. Beta phase was distributed between alpha-2 laths. This sample also e x h i b i t e d d i s c o n t i n u o u s p r e c i p i t a t i o n at grain b o u n d a r i e s in the gage section and not in the buttonhead.. The d i s t r i b u t i o n of these pha~es at the grain b o u n d a r y is shown in Figure 4(b) . Discontinuous phase growth was by the g r o w t h of p a r a l l e l high and low d e n s i t y laths p e r p e n d i c u l a r to the grain boundary. In this case, the diffusion distance was equal to the lath spacing. Because of this, there appears to be no change in the c h a r a c t e r i s t i c diffusion distance with the extent of discontinuous phase growth. Beta heat t r e a t e d s a m p l e 934-B3 (1448/He), had a coarse W i d n a n s t ~ t t e n m i c r o s t r u c t u r e , with coarse 5 x 30 u m laths m i x e d with finer 1 x 3 um laths, Figures 5(a) and 5(b). All laths had a fine low density backbone as did the grain boundaries. Beta phase was present as thin layers between alpha-2 laths. There was little evidence of discontinuous grain boundary precipitation in this sample. Primary creep lasted only ii0 hours and the total extent of primary creep was 0.32%. Discussion of Results D i s c o n t i n u o u s p r e c i p i t a t i o n and dissolution of gamma-prime has been observed during tensile and c o m p r e s s i v e creep of the gamma + g a m m a - p r i m e superalloy Ni16Cr-5AI-4Ta (5) and in IN-100 (6). It was shown that the stress sense on the grain b o u n d a r i e s a f f e c t e d c h r o m i u m d e p l e t i o n and enhancement and d i s c o n t i n u o u s g a m m a - p r i m e precipitation. S t r e s s - a s s i s t e d discontinuous p r e c i p i t a t i o n has also been o b s e r v e d in copper-cadmium, and the stress sense, c o m p r e s s i o n or tension, d e t e r m i n e d which grain boundaries e x h i b i t e d discontinuous precipitation.(7) The magnitude of the stress directly affected the rate of discontinuous phase growth. Our o b s e r v a t i o n of d i s c o n t i n u o u s p r e c i p i t a t i o n in T i 3 A I - 1 2 . 5 N b is the first o b s e r v a t i o n of this phenomenon in t i t a n i u m aluminide alloys. It is reasonable to assume that the high density of the dark discontinuous phase is due to a niobium content which is similar to that of the beta phase. S t r e s s - r e l a t e d migration of niobium from transformed alpha-2 plus beta grains to the discontinuous phase would then appear to be necessary. Dissolution of beta laths in the alpha-2 plus beta regions is evident from the comparison of Figures 2 and 3. The p r i m a r y and s t e a d y state c r e e p rate of fine g r a i n e d sample 934-AB (1313/He) was high, c o n s i s t e n t with creep d o m i n a t e d by grain b o u n d a r y sliding. The m a g n i t u d e of the p r i m a r y creep strain was high, also. A 0.6 ~m thick two phase layer grew at grain b o u n d a r i e s which were p e r p e n d i c u l a r to the load axis. Since d i s c o n t i n u o u s phase growth did not occur in the b u t t o n h e a d region, one can conclude that its growth was stress or grain boundary deformation-enhanced. The extent of discontinuous phase growth may have been limited by the orientation of the two phases relative to the stress axis. D i s c o n t i n u o u s growth occurred with both p h a s e s g r o w i n g in the plane of the grain boundary. At this orientation, further growth would appear to require i n c r e a s i n g l y longer diffusion distances. The time p e r i o d over which transient creep occurred in this sample was 33 hours, consistent with limited discontinuous phase growth. D i s c o n t i n u o u s p r e c i p i t a t i o n o c c u r r e d in b e t a heat t r e a t e d T i 3 A l - 1 2 . 5 N b samples except those with coarse W i d m a n s t ~ t t e n lath microstructures, it has also been o b s e r v e d in Ti3AI-7.5Nb samples at 923°K.(8) Discontinuous precipitation at the grain boundaries of beta heat treated and salt quenched sample 938-CS took the form of p a r a l l e l two phase lath growth p e r p e n d i c u l a r to the plane of the grain boundary. This could be d e s c r i b e d as d i s c o n t i n u o u s c o a r s e n i n g rather than d i s c o n t i n u o u s p r e c i p i t a t i o n except that the c o m p o s i t i o n of the d i s c o n t i n u o u s phases a p p e a r e d to be d i f f e r e n t than that of the parent phases. Prinary creep
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CREEP IN Ti3AI-Nb
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lasted 420 hours in this sample, consistent with c o n t i n u e d discontinuous precipitation throughout the creep test. Bending of the discontinuous a!pha-2 and beta laths in Figure 4(b) is an indication of deformation of the discontinuous !a%hs during their growth. Discontinuous precipitation only occurred in the gage section of this sample, not in the buttonhead. This again suggests stress or grain boundary sliding-enhanced growth of the discontinuous phases. An additional driving force for d i s c o n t i n u o u s precipitation in T i 3 A I - 1 2 . 5 N b may be the difference between alpha-2 and beta compositions at the creep test temperature of 923~K and the initial compositions established during heat treatment, it could also be driven by the precipitation of alpha-2 plus ordered orthorhombic Ti2AiNb phase. Coarse Widmanst~tten sample 934-B3, which was produced by cooling at a slower rate than sample 93S-OS, did not exhibit discontinuous precipitation. This sample differed from salt quenched sample 938-0S in several ways: it had a lower aluminum content, it transformed to alpha-2 at a slower cooling rate, and it had a coarse Widmanst~tten lath structure. This may have resulted in alpha-2 and beta compositions closer to 923°K equilibrium compositions, reducing the driving force for d i s c o n t i n u o u s precipitation. Detailed analysis of the composition and identity of the initial and discontinuous phases will have to be performed to confirm which factor contributed most significantly to this difference. Sample 934-B3 had the lowest total primary creep strain of all three samples. The observation of high primary creep strains in samples in which discontinuous precipitation occurred and low primary creep strain in sample 934-B3 suggests that the extent of primary creep in fine grained sample 934-AB and salt quenched sample 938-0S was determined by the growth of discontinuous phases. It is probable that the measured minimum creep rate of these samples was also influenced by discontinuous precipitation. Primary or transient creep is usually associated with dislocation hardening. In the case of discontinuous precipitation another process appears to occur. New undeformed grains of alpha-2 and beta or alpha-2 and orthorhombic phase at the grain boundaries could permit additional dislocation generation at the high rates characteristic of the initial stages of primary creep. The rate of primary creep thus may thus depend upon the rate of growth of the discontinuous phases. Continued discontinuous phase growth during the steady state creep regime would also increase the minimum creep rate. Conclusions i. Discontinuous precipitation, which was observed in the gage section and not in the buttonhead of 923°K, 175 MPa creep samples of Ti3AI-12.SNb, was associated either with applied stress or grain boundary deformation. 2. D i s c o n t i n u o u s precipitation led to high primary phenomenon may contribute to the high primary creep observed in other Ti3AI-Nb alloys.
creep strains. strain which has
This been
3. Different forms of discontinuous precipitation were observed in Ti3Al-12.5Nb samples having fine grained equiaxed or coarse g r a i n e d t r a n s f o r m e d beta microstructures. 4. Slowly cooled, coarse discontinuous precipitation.
Widmanst~tten
5. Measured m i n i m u m precipitation.
rates
creep
may
have
Ti-24.7Al-12.9Nb been
influenced
did by
not
exhibit
discontinuous
References I. R.G. Rowe, in Hiah Temoerature Aluminides and Intermetallics, Proceedings, Symposium, Indianapolis, IA Oct. 1989, TMS-AIME, Warrendale, PA, to be published. 2. Atlas of Creed and S~ress Rupture Curves. Ed. H. Boyer, ASM International,
Vol.
:.,~' No
CREEP
7 •
IN T i . A I - N b
~,'~13
D
~ : ~ a = = Park, OH 1988 3. ~ . B a n.e r j e.e , . A.K. . Gogia, . . T. K. . " ~. .a n d . V A. Joshi, . . .~a'~ ~,'~, , 36, 871 (1989) . 4. K. M u r a l e e d h a r a n a n d D. B a n e r j e e , Met. Trans. A, 20A, 1139 (i989) . 5. J.K. T i e n a n d R.P. Gamble, Met. Trans., 2, 1663 (!97!). 6 M. ~.orlsh~ta, F. K o s h i i s h i , u N a g a i and K Shoi'-, J. Jpn. Soc. Po;-:der P o w d e r Y.eta!!., 32, 232 (1985) 7. M.S. Su!onen, A c t a Met., 12, 749 (i964). 8. R.G. Rowe, M.F.X. G i g ! i o t t i and B.J. M a r q u a r d t unp~:b!ished r e s e a r c h . Acknowledgements The a u t h o r s w o u l d like to a c k n o w l e d g e GE A i r c r a f t E n g i n e s s u p p o r t for this w o r k u n d e r IR&D f u n d i n g . We 'wou.~' ~ a l s o like to t h a n k ~rof. Duk ~. Yoon of the Korea Advanced I n s t i t u t e of S c i e n c e and ~'~"~-,~±~x for h~!o~,,~_ .~_ d i s c u s s i o n s on discontinuous precipitation and the effezt of stress on discontinuous precipitation.
12 i
923K/175MPa
~ o
/
0.6
5 t 923KI175MPa
/
923K/175MPa
0.5 £ 0.4 u~ 0.3
sl
== &-
0.2
0
0.1
O~
0
0
. . . . . .
80
20 40 60 Time, hrs. (a)
0
0.0
200400600800 Time, hrs.
0
50
100
150
200
Time, hrs.
(b)
(c)
FIG. !. C r e e p c u r v e s at 923°K a n d 175 M P a of T i 3 A ! - ! 2 . 5 N b samples with various heat t r e a t m e n t s , (a) S a m p l e 9 3 2 - A B w h i c h h a d a fine g r a i n e d equiaxed alpha-2 m i c r o s t r u c t u r e , a n d s a m p l e s (b) 9 3 8 - O S and (c) 934-B w h i c h were beta heat t r e a t e d and c o o l e d at d i f f e r e n t q u e n c h rates .
I- - I
~
I
(a) FIG. 2. Micrographs of T i 3 A ! - I 2 . S N b creep sample B.~K~av~_~_d e l e c t r o n i m a g e s of the b u t t o n h e a d
-
-
-I
(b) 932-AB
(1313/He)
(a),
(b)
1214
CREEP IN Ti3AI-Nb
I ~ FIG. 3. (1313/He) represent
~:_~
-~
Vol.
24, No.
7
rj
(a) (b) Micrographs of the gage section of Ti3Al-12.SNb creep sample 932-AB (a), (b) Backscattered electron images of the gage section. Dark areas regions of higher density.
~
(a) (b) FIG. 4. (a) Optical micrograph of the gage section of Ti3AI-12.5Nb creep sample 938-OT (1408/SQ) (b) Backscattered scanning electron image of creep sample 938-0T (1408/SQ). Dark areas represent regions of higher density
(a) (b) FIG. 5 (a) Optical micrograph of the gage section of Ti3AI-12.5Nb creep sample 934-B3 (1448/He) (b) Backscattered scanning electron image. Dark areas represent regions of higher density
Vol.
CREEP
24, pp. 1209-121¢, 1990 Printed in the U.S.A.
AND DISCONTINUOUS
P e r g a m o n Press plc All rights r e s e r v e d
PRECIPITATION
IN
A Ti3A1-Nb ALLOY AT 923°K
GE Corporate
R. G. Rowe and M. F. X. Gigliotti Research and Development, Schenectady,
NY 12301
and GE Aircraft
B. J. Marquardt Engines, Cincinnati,
(Received
April
17,
OH 45215
1990)
Introduction Creep r e s i s t a n c e is c r i t i c a l for the high t e m p e r a t u r e a p p l i c a t i o n s in which Ti3Al-base alloys are envisioned. The creep rate below 0.2% is g e n e r a l l y high for these alloys b e c a u s e t r a n s i e n t or p r i m a r y creep can e x t e n d to szrains of several tenths p e r c e n t (I) . This is much g r e a t e r than the amount cf primary creep o b s e r v e d in nickel-base superalloys of comparable tensile strength. (2) The alloy Ti3AI-12.5Nb has a m i c r o s t r u ct u r e which consists of alpha-2 (D0!~) and beta (bcc) phases which are formed by transformation from an e l e v a t e d temperature beta phase field. O r t h o r h o m b i c Ti2AINb has also been i d e n t i f i e d in this alloy (3). T r a n s f o r m a t i o n from beta to alpha-2 plus beta results in laths of a!pha-2 phase which vary in size a c c o r d i n g to the c o o l i n g rate t h r o u g h the t r a n s f o r m a t i o n temperature. We have shown that m i c r o s t r u c t u r e s t r o n g l y a f f e c t s %he creep behavior of this alloy. (i) In this paper, we r e p o r t our o b s e r v a t i o n of discontinuous precipitation at grain boundaries during 923°K creep testing. Experimental
Procedure
The alloy Ti3AI-12.5Nb at.% was prepared by n o n - c o n s u m a b l e arc mei~ing. The compositions of the two heats are t a b u l a t e d in Table I. The alloys were cast into copper chill molds and extruded at 1313°K-1355°K to 12:1 reduction in area. Samples were heat treated in argon and m a c h i n e d into b u t t o n h e a d tensile specimens with a gage section 2 mm d i a m e t e r by 11.7 mm long. Three m i c r o s t r u c ~ u r e s were studied. These were: alpha-2 plus beta with 60% prior alpha-2, and zransformed beta m i c r o s t r u c t u r e s with two different t r a n s f o r m a t i o n lath d i m e n s i c n s . Heat treatment conditions are shown in Table 2. TABLE SAMPLE 932,934 938
#
i.
Alloy Compositions Ti Bal. Bal.
(at.%)
A1
Nb
24.7 26.5
12.9 12.6
Creep samples were d e a d weight l o a d e d at a s t r e s s of 175 M2a and a temperature of 923°K in a r g o n gas. Sample extension was m e a s u r e d by an e x t e n s o m e t e r on the b u t t o n h e a d of the sample. E x t e n s i o n was d i g i t a l l y recorded c o n t i n u o u s l y f r o m the onset of loading. The e x t e n s i o n of sample 934-B3 was recorded manually.
1209 0 0 3 6 - 9 7 4 8 / 9 0 $3.00 + .00 C o p y r i g h t (c) 1990 P e r g a m o n Press
plc
1210
CREEP
TABLE 2.
IN Ti3AI-Nb
Vol.
24, No.
7
Heat Treatment Conditions Heat Treatment
Grain Size (~m)
Sample #
Notation
932-AB
1313/He
1313°K/lhr/He+1033°K/lhr
6-8
938-OS,OT
1408/SQ
1408°K/lhr/Salt Quench to 1088°K,+1033°K/lhr
300
934-B3
1448/He
1448°K/lhr/He+1033°K/lhr
750
Experimental Results The m i n i m u m creep rate and total primary creep strain of samples with three different m i c r o s t r u c t u r e s are tabulated in Table 3. Figure l(a) shows the creep curve of alpha-2 plus beta heat treated sample 932-AB. The m i n i m u m creep rate extension is shown superimposed as a dotted line. The total primary creep strain of this sample was 2.0%. The creep curves for beta heat treated samples 934-B3 (helium cooled) and 938-OS (salt quenched) are shown in Figures l(b)and l(c), respectively. The total primary creep strain for salt quenched sample 938-OS was 1.5%. Sample 934-B3 which was heat treated at 1448°K and helium cooled had only 0.3% total primary creep strain. TABLE 3.
Creep Test Results,
923°K, 175 MPa
Primary Creep Strain (%)
Minimum Creep Rate (sec -I)
Time to Failure or Unload (hrs)
934-AB (1313/He)
2.0
3.9E-7
68
934-B3 (1448/He)
0.3
2.4E-9
236
(unloaded)
938-OS (1408/SQ)
1.5
8.7E-9
811
(unloaded)
Sample Number
The microstructures of these samples are shown in Figures 2-5. Figures 2(a) and 2(b) show the e q u i a x e d grain structure of sample 932-AB. The grain size was 6-8 ~m. This sample c o n s i s t e d of e q u i a x e d prior alpha-2 grains s u r r o u n d e d by transformed alpha-2 plus beta. The beta phase showed up as dark areas between light alpha-2 laths in this b a c k s c a t t e r e d electron image. The fraction of prior alpha phase was approximately 60%. The gage section had a different distribution of beta phase. Figures 3(a) and 3(b) show this. The gage section of this sample consisted of e q u i a x e d alpha-2 grains s u r r o u n d e d by adjacent alpha-2 plus beta regions in which the density contrast of the two phases was reduced. Grain boundaries p e r p e n d i c u l a r to the load axis e x h i b i t e d d i s c o n t i n u o u s high and low d e n s i t y p h a s e s w h i c h had a g r e a t e r d e n s i t y d i f f e r e n c e and hence g r e a t e r compositional difference than the transformed alpha-2 plus beta regions. This is shown in Figure 3(b) . The b a c k s c a t t e r e d e l e c t r o n image d e n s i t y for the dark discontinuous phase was darker than the initial beta phase image density and the image density of the light discontinuous phase was less than that of the initial alpha-2 p h a s e image density. Rupture, which o c c u r r e d by s e p a r a t i o n at the interfaces between the discontinuous phases and the prior alpha-2 grains, limited the total thermal exposure of this sample to 68 hours. It is r e a s o n a b l e to assume that the d i s c o n t i n u o u s phases are alpha-2 and beta phases of c o m p o s i t i o n s c o r r e s p o n d i n g to the creep t e m p e r a t u r e , 923°K. Banerjee (4) has shown that the niobium content in the beta phase of Ti-24AI-IINb changed rapidly from 1373°K to 1315°K. Our results suggest that the composition
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of the alpha-2 phase must also change below this temperature. It has been shown that o r t h o r h o m b i c Ti2AINb forms in this alloy (3), and it will take quantitative a n a l y s i s to unambiguously establish the identity of the discontinuous phases Figure 4 shows the m i c r o s t r u c t u r e of the gage section of beta heat treated and salt q u e n c h e d sample 938-OT (1448/SQ) . This sample had a two phase W i d m a n s t ~ t t e n m i c r o s t r u c t u r e with coarse 1 X 8 um laths interspersed with finer 0.5 x 3 um laths. Beta phase was distributed between alpha-2 laths. This sample also e x h i b i t e d d i s c o n t i n u o u s p r e c i p i t a t i o n at grain b o u n d a r i e s in the gage section and not in the buttonhead.. The d i s t r i b u t i o n of these pha~es at the grain b o u n d a r y is shown in Figure 4(b) . Discontinuous phase growth was by the g r o w t h of p a r a l l e l high and low d e n s i t y laths p e r p e n d i c u l a r to the grain boundary. In this case, the diffusion distance was equal to the lath spacing. Because of this, there appears to be no change in the c h a r a c t e r i s t i c diffusion distance with the extent of discontinuous phase growth. Beta heat t r e a t e d s a m p l e 934-B3 (1448/He), had a coarse W i d n a n s t ~ t t e n m i c r o s t r u c t u r e , with coarse 5 x 30 u m laths m i x e d with finer 1 x 3 um laths, Figures 5(a) and 5(b). All laths had a fine low density backbone as did the grain boundaries. Beta phase was present as thin layers between alpha-2 laths. There was little evidence of discontinuous grain boundary precipitation in this sample. Primary creep lasted only ii0 hours and the total extent of primary creep was 0.32%. Discussion of Results D i s c o n t i n u o u s p r e c i p i t a t i o n and dissolution of gamma-prime has been observed during tensile and c o m p r e s s i v e creep of the gamma + g a m m a - p r i m e superalloy Ni16Cr-5AI-4Ta (5) and in IN-100 (6). It was shown that the stress sense on the grain b o u n d a r i e s a f f e c t e d c h r o m i u m d e p l e t i o n and enhancement and d i s c o n t i n u o u s g a m m a - p r i m e precipitation. S t r e s s - a s s i s t e d discontinuous p r e c i p i t a t i o n has also been o b s e r v e d in copper-cadmium, and the stress sense, c o m p r e s s i o n or tension, d e t e r m i n e d which grain boundaries e x h i b i t e d discontinuous precipitation.(7) The magnitude of the stress directly affected the rate of discontinuous phase growth. Our o b s e r v a t i o n of d i s c o n t i n u o u s p r e c i p i t a t i o n in T i 3 A I - 1 2 . 5 N b is the first o b s e r v a t i o n of this phenomenon in t i t a n i u m aluminide alloys. It is reasonable to assume that the high density of the dark discontinuous phase is due to a niobium content which is similar to that of the beta phase. S t r e s s - r e l a t e d migration of niobium from transformed alpha-2 plus beta grains to the discontinuous phase would then appear to be necessary. Dissolution of beta laths in the alpha-2 plus beta regions is evident from the comparison of Figures 2 and 3. The p r i m a r y and s t e a d y state c r e e p rate of fine g r a i n e d sample 934-AB (1313/He) was high, c o n s i s t e n t with creep d o m i n a t e d by grain b o u n d a r y sliding. The m a g n i t u d e of the p r i m a r y creep strain was high, also. A 0.6 ~m thick two phase layer grew at grain b o u n d a r i e s which were p e r p e n d i c u l a r to the load axis. Since d i s c o n t i n u o u s phase growth did not occur in the b u t t o n h e a d region, one can conclude that its growth was stress or grain boundary deformation-enhanced. The extent of discontinuous phase growth may have been limited by the orientation of the two phases relative to the stress axis. D i s c o n t i n u o u s growth occurred with both p h a s e s g r o w i n g in the plane of the grain boundary. At this orientation, further growth would appear to require i n c r e a s i n g l y longer diffusion distances. The time p e r i o d over which transient creep occurred in this sample was 33 hours, consistent with limited discontinuous phase growth. D i s c o n t i n u o u s p r e c i p i t a t i o n o c c u r r e d in b e t a heat t r e a t e d T i 3 A l - 1 2 . 5 N b samples except those with coarse W i d m a n s t ~ t t e n lath microstructures, it has also been o b s e r v e d in Ti3AI-7.5Nb samples at 923°K.(8) Discontinuous precipitation at the grain boundaries of beta heat treated and salt quenched sample 938-CS took the form of p a r a l l e l two phase lath growth p e r p e n d i c u l a r to the plane of the grain boundary. This could be d e s c r i b e d as d i s c o n t i n u o u s c o a r s e n i n g rather than d i s c o n t i n u o u s p r e c i p i t a t i o n except that the c o m p o s i t i o n of the d i s c o n t i n u o u s phases a p p e a r e d to be d i f f e r e n t than that of the parent phases. Prinary creep
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lasted 420 hours in this sample, consistent with c o n t i n u e d discontinuous precipitation throughout the creep test. Bending of the discontinuous a!pha-2 and beta laths in Figure 4(b) is an indication of deformation of the discontinuous !a%hs during their growth. Discontinuous precipitation only occurred in the gage section of this sample, not in the buttonhead. This again suggests stress or grain boundary sliding-enhanced growth of the discontinuous phases. An additional driving force for d i s c o n t i n u o u s precipitation in T i 3 A I - 1 2 . 5 N b may be the difference between alpha-2 and beta compositions at the creep test temperature of 923~K and the initial compositions established during heat treatment, it could also be driven by the precipitation of alpha-2 plus ordered orthorhombic Ti2AiNb phase. Coarse Widmanst~tten sample 934-B3, which was produced by cooling at a slower rate than sample 93S-OS, did not exhibit discontinuous precipitation. This sample differed from salt quenched sample 938-0S in several ways: it had a lower aluminum content, it transformed to alpha-2 at a slower cooling rate, and it had a coarse Widmanst~tten lath structure. This may have resulted in alpha-2 and beta compositions closer to 923°K equilibrium compositions, reducing the driving force for d i s c o n t i n u o u s precipitation. Detailed analysis of the composition and identity of the initial and discontinuous phases will have to be performed to confirm which factor contributed most significantly to this difference. Sample 934-B3 had the lowest total primary creep strain of all three samples. The observation of high primary creep strains in samples in which discontinuous precipitation occurred and low primary creep strain in sample 934-B3 suggests that the extent of primary creep in fine grained sample 934-AB and salt quenched sample 938-0S was determined by the growth of discontinuous phases. It is probable that the measured minimum creep rate of these samples was also influenced by discontinuous precipitation. Primary or transient creep is usually associated with dislocation hardening. In the case of discontinuous precipitation another process appears to occur. New undeformed grains of alpha-2 and beta or alpha-2 and orthorhombic phase at the grain boundaries could permit additional dislocation generation at the high rates characteristic of the initial stages of primary creep. The rate of primary creep thus may thus depend upon the rate of growth of the discontinuous phases. Continued discontinuous phase growth during the steady state creep regime would also increase the minimum creep rate. Conclusions i. Discontinuous precipitation, which was observed in the gage section and not in the buttonhead of 923°K, 175 MPa creep samples of Ti3AI-12.SNb, was associated either with applied stress or grain boundary deformation. 2. D i s c o n t i n u o u s precipitation led to high primary phenomenon may contribute to the high primary creep observed in other Ti3AI-Nb alloys.
creep strains. strain which has
This been
3. Different forms of discontinuous precipitation were observed in Ti3Al-12.5Nb samples having fine grained equiaxed or coarse g r a i n e d t r a n s f o r m e d beta microstructures. 4. Slowly cooled, coarse discontinuous precipitation.
Widmanst~tten
5. Measured m i n i m u m precipitation.
rates
creep
may
have
Ti-24.7Al-12.9Nb been
influenced
did by
not
exhibit
discontinuous
References I. R.G. Rowe, in Hiah Temoerature Aluminides and Intermetallics, Proceedings, Symposium, Indianapolis, IA Oct. 1989, TMS-AIME, Warrendale, PA, to be published. 2. Atlas of Creed and S~ress Rupture Curves. Ed. H. Boyer, ASM International,
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~,'~13
D
~ : ~ a = = Park, OH 1988 3. ~ . B a n.e r j e.e , . A.K. . Gogia, . . T. K. . " ~. .a n d . V A. Joshi, . . .~a'~ ~,'~, , 36, 871 (1989) . 4. K. M u r a l e e d h a r a n a n d D. B a n e r j e e , Met. Trans. A, 20A, 1139 (i989) . 5. J.K. T i e n a n d R.P. Gamble, Met. Trans., 2, 1663 (!97!). 6 M. ~.orlsh~ta, F. K o s h i i s h i , u N a g a i and K Shoi'-, J. Jpn. Soc. Po;-:der P o w d e r Y.eta!!., 32, 232 (1985) 7. M.S. Su!onen, A c t a Met., 12, 749 (i964). 8. R.G. Rowe, M.F.X. G i g ! i o t t i and B.J. M a r q u a r d t unp~:b!ished r e s e a r c h . Acknowledgements The a u t h o r s w o u l d like to a c k n o w l e d g e GE A i r c r a f t E n g i n e s s u p p o r t for this w o r k u n d e r IR&D f u n d i n g . We 'wou.~' ~ a l s o like to t h a n k ~rof. Duk ~. Yoon of the Korea Advanced I n s t i t u t e of S c i e n c e and ~'~"~-,~±~x for h~!o~,,~_ .~_ d i s c u s s i o n s on discontinuous precipitation and the effezt of stress on discontinuous precipitation.
12 i
923K/175MPa
~ o
/
0.6
5 t 923KI175MPa
/
923K/175MPa
0.5 £ 0.4 u~ 0.3
sl
== &-
0.2
0
0.1
O~
0
0
. . . . . .
80
20 40 60 Time, hrs. (a)
0
0.0
200400600800 Time, hrs.
0
50
100
150
200
Time, hrs.
(b)
(c)
FIG. !. C r e e p c u r v e s at 923°K a n d 175 M P a of T i 3 A ! - ! 2 . 5 N b samples with various heat t r e a t m e n t s , (a) S a m p l e 9 3 2 - A B w h i c h h a d a fine g r a i n e d equiaxed alpha-2 m i c r o s t r u c t u r e , a n d s a m p l e s (b) 9 3 8 - O S and (c) 934-B w h i c h were beta heat t r e a t e d and c o o l e d at d i f f e r e n t q u e n c h rates .
I- - I
~
I
(a) FIG. 2. Micrographs of T i 3 A ! - I 2 . S N b creep sample B.~K~av~_~_d e l e c t r o n i m a g e s of the b u t t o n h e a d
-
-
-I
(b) 932-AB
(1313/He)
(a),
(b)
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I ~ FIG. 3. (1313/He) represent
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-~
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rj
(a) (b) Micrographs of the gage section of Ti3Al-12.SNb creep sample 932-AB (a), (b) Backscattered electron images of the gage section. Dark areas regions of higher density.
~
(a) (b) FIG. 4. (a) Optical micrograph of the gage section of Ti3AI-12.5Nb creep sample 938-OT (1408/SQ) (b) Backscattered scanning electron image of creep sample 938-0T (1408/SQ). Dark areas represent regions of higher density
(a) (b) FIG. 5 (a) Optical micrograph of the gage section of Ti3AI-12.5Nb creep sample 934-B3 (1448/He) (b) Backscattered scanning electron image. Dark areas represent regions of higher density