The creep behavior of the Ti3Al alloy Ti24Al11Nb

Scripta METALLURGICA

Vol.

23, pp. 1931-1956, 1989 Printed in the U.S.A.

Pergamon Press plc All rights reserved

THE CREEP BEHAVIOR OF THE Ti3AI ALLOY Ti-24AI-IINb R. W. Hayes Metals Technology, Inc. 19801 Nordhoff Street Northridge, California 91324 (Received June 8, 1989) (Revised September 12, 1989) Introduction The intermetallic compound Ti3AI (referred to as ~2) has the DO19 ordered HCP crystal structure and is an excellent candidate high temperature material for structural components in advanced aircraft engines. Its detracting feature is its lack of ductility at temperatures below about 760°C. In order for this intermetallic compound to become a useful engineering material, the ductility below 760°C must be improved without a significant reduction in the high temperature mechanical properties such as creep resistance. In order to achieve this, alloys based upon the Ti3AI compound have been developed. One such alloy currently being evaluated by the air breathing jet engine industry is the alloy Ti-24AI-IINb (Ti-13 wt% AI-21 wt% Nb). Within the temperature regime of potential engineering application, the Ti-24AI-IINb alloy consists of a two phase mixture of the ordered HCP ~ 2 and the BCC ~ with the primary phase being the ~2The purpose of the present study is to evaluate the steady-state creep behavior of the Ti-24Al-llNb alloy and compare these results to the results previously obtained by Mendiratta and Lipsitt on two near stoichiometric Ti3AI compounds (I). The intermetallic compounds studied by Mendiratta and Lipsitt consisted oJ Ti-16 wt% A1 and Ti-16 wt% AI-10 wt% Nb and were referred to as Ti3AI and Ti3AI-10 wt% Nb respectively (1). Experimental The material used in this study was obtained from Timet Corporation in the form of hot rolled plate. The complete chemical composition in weight percent is given in Table I. The precise melting and rolling history is not known. However, examination of the microstructure prior to heat treatment indicated that the rolling was performed in the two phase ~2 plus ~ region. The beta transus of Ti-24-Ii is suggested to be about I130°C. Prior to specimen fabrication, the as received plate was given a 1 hour heat treatment at 1000°C followe~ by an air cool in order to stabilize the material prior to testing. This heat treat is similar to the one employed by Mendiratta and Lipsitt on their alloys prior to testing (i). After heat treat, the microstructure of the Ti-24-Ii was examined optically. The microstructure consisted of a two phase ~2 plus ~ mixture with a grain size on the order of 6~m. After heat treating, threaded round creep specimens were fabricated by stress free crush grinding. The specimens consisted of shoulders llmm in diameter with a 6.35mm diameter by 31.75mm long reduced section. A shallow groove was machined onto the shoulders for placement of the creep extensometer. All creep tests were conducted in Satec systems constant load creep machines with a 16:1 lever arm ratio. All testin~ was performed in an air atmosphere with one reference test being conducted in argon. The creep machines were equipped with Satec systems three zone resistance heating furnaces. Specimen

1931 0036-9748/89 $3.00 + .00 Copyright (c) 1989 Pergamon Press plc

1932

CREEP OF Ti-24AI-IINb

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temperatures were all monitored by attachment of two chromel-alumel type K thermocouples directly to the specimen reduced section. All temperatures were held to within ±2°C of the desired test temperature. The creep strain was measured with mechanical dual dial gage averaging type extensometers. The tests were conducted over the temperature range of 650"C to 760"C at constant applied initial stress levels of 69 MPa to 172.35 MPa. All tests were terminated upon firm establishment of steady-state creep. The specimens, upon completion of testing, were allowed to cool to room temperature under load to minimize microstructural changes taking place as a result of recovery processes. Results Figure I shows the typical creep curves exhibited by Ti-24AI-IINb at 650"C, 704°C and 760°C at a constant initial stress of 103.4 MPa. At all temperatures, a primary followed by a steady-state regime is observed. The curves displayed in Figure 1 are typical of the curves obtained at all temperatures and stresses examined. The steady-state creep rate was determined by calculating the slope of the linear portion of the creep strain vs. time curves such as those presented in Figure i. A complete summary of the creep results obtained in this study are presented in Table 2. A significant observation made in the present study is the absence of surface microcracking on the specimens tested in air. Several researchers have observed extensive microcracking on the surface of Ti-24AI-IINb after creep testing in air (2). The origin of the microcracking as well as its influence on the steady state creep behavior of the alloy is not well established and indeed, more work is needed to address this issue. As mentioned previously, the creep specimens employed in this study were prepared by low stress crush ~rinding with no subsequent surface preparation being employed. A s indicated in Table 2, the steadystate creep rate of Ti-24AI-IINb tested in alr and in argon at 704°C and a stress of 103.4 MPa are virtually identical. From the data in Table 2, both temperature and stress dependence of the steady-state creep rate of Ti-24AI-IINb can be obtained. Figure 2 shows an Arrhen~us plot of log steady-state creep rate (i s ) versus the reciprocal of the absolute temperature (l/T) at 103.4 MPa and 137.8 MPa. This data can be described by the equation: \

I/T

yo

(Eqn. I)

where - A H is the apparent activation energy for the rate controlling creep mechanism, and o is the applied stress. From figure 2 the - ~ H was found to be 2.59Xi05 J/Mole. As shown in Figure 2 the - ~ H appears to be independent of stress level. These results are in good agreement with the well established fact that the - ~ H is independent of stress within the temperature regime where a single thermally activated process controls the creep rate. A second fundamental measurement of the process controlling the creep response is obtained by the evaluation of the stress dependence of the steady-state creep rate. In Figure 3 the log ~ plotted against the log o at 650"C and 704°C is presented. Calculation of the slopes of Figure 3 through the relationship: ~s = Ko n

(Eqn. 2)

(where K is a material constant) gives the stress exponent n. From Figure 3 the stress exponent n was found to be 4.32. As shown in Figure 3 the stress exponent is independent of temperature thus indicating a single rate controlling process governing the steady-state creep rate of Ti-24AI-IINb. The linear plots obtained for both the - ~ H and n indicate that a single thermally activated process is controlling the steady-state creep behavior of the Ti-24AI-IINb within the temperature and stress regime of the present study. Discussions In their previous studies of the creep behavior of Ti-16 wt% A1 and Ti-16 wt% AI-10 wt% Nb, Mendiratta and Lipsitt (I) found two distinct creep regimes. At temperatures above 7000C and stresses above 138 MPa, steady-state creep was sug-

Vol.

23, No.

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CREEP OF T i - 2 4 A I - I I N b

1933

gested to be c o n t r o l l e d by d i s l o c a t i o n climb. The a c t i v a t i o n energy for creep of Ti-16 wt% A1 at 700°C and above was found to be 2.06XI05 J / M o l e and the stress e x p o n e n t at a>138 MPa was b e t w e e n 4.3 and 5.0. A d d i n g i0 wt% Nb (Ti-16 wt% AI-10 wt% Nb) i n c r e a s e d the a c t i v a t i o n energy to 2.85XI05 J / M o l e at temperature above 650°C w i t h a stress e x p o n e n t n equal to 6.0 at s t r e s s e s above 172.5 MPa. The above results found by M e n d i r a t t a and L i p s i t t (i) are in good agreement w i t h expected values for n in the p o w e r law creep regime (3). It was e m p h a s i z e d (i) t h a t the - ~ H c o u l d not be c o r r e l a t e d w i t h the lattice self diffusion v a l u e - ~ H s D d u e to the lack of d i f f u s i o n data for these materials. The data o b t a i n e d by M e n d i r a t t a and L i p s i t t (i) was found to c o n f o r m to the power law e q u a t i o n of the form: ~(a/E)

n

EXP

(-~H

/RT)

(Eqn. 3)

The t r a n s i t i o n in creep m e c h a n i s m o b s e r v e d by M e n d i r a t t a and Lipsitt (i) was e v i d e n c e d by the change in stress e x p o n e n t n and the a c t i v a t i o n energy as the applied stress a n d / o r t e m p e r a t u r e s d e c r e a s e d b e l o w c e r t a i n m i n i m u m values. This was o b s e r v e d in b o t h of t h e i r alloys. It was s u g g e s t e d that the t r a n s i t i o n might r e p r e s e n t a c h a n g e from d i s l o c a t i o n climb to a g r a i n b o u n d a r y sliding mechanism. However, no firm c o n c l u s i o n s were drawn again due to the lack of the a p p r o p r i a t e d i f f u s i o n data (i). In c o m p a r i n g the results of the p r e s e n t study to those of M e n d i r a t t a and Lipsitt (i), the m o s t s i g n i f i c a n t o b s e r v a t i o n is in the shift t o w a r d lower stress values where t y p i c a l p o w e r law creep appears to operate in Ti-24AI-IINb. The stress e x p o n e n t and apparent activation energy o b t a i n e d in the p r e s e n t study are in reasonable agreement w i t h the results o b t a i n e d by M e n d i r a t t a and L i p s i t t (i) in their p o w e r law regime. The data o b t a i n e d in the p r e s e n t study was found to conform to the p o w e r law e q u a t i o n of the form: ~(a)

n

EXP

(-~H

/RT)

(Eqn.

4)

which is the u n n o r m a l i z e d v e r s i o n of e q u a t i o n 3. It is well e s t a b l i s h e d that when p o w e r law c r e e p operates, the rate g o v e r n i n g p r o c e s s is the climb of edge d i s l o c a t i o n s over o b s t a c l e s w i t h i n the grain interior with the - ~ H being equal to the - ~ H _ . It has also been shown that the p o w e r law ~ p o s s e s s e s a strong subgrain o r b ~ r a i n size dependence. In a recent publication, Ruano et al (4) shows that at a c o n s t a n t n o r m a l i z e d stress (a/E) the n o r m a l i z e d p o w e r law creep rate (~s/D A p ~ ) w h e r e A p L is the p o w e r law creep constant, d e c r e a s e s in direct p r o p o r £ ~ o n w i t h d e c r e a s i n g s u b g r a i n or grain size w i t h a r e l a t i o n s h i p equal to a p o w e r of 4. Thus the p o w e r law Es can be taken as a direct m e a s u r e of the rate in w h i c h a d i s l o c a t i o n can climb over its obstacle. Figure 4 shows an A r r h e n i u s p l o t of ~ v e r s u s the r e c i p r o c a l of the a b s o l u t e t e m p e r a t u r e for Ti-16 wt% A I - I O wt% Nb r e p r o d u c e d from M e n d i r a t t a and L i p s i t t (i). In comparing the r e s u l t s of F i g u r e 4 to the results p r e s e n t e d in Figure 2 of this study, one readily o b s e r v e s the s u b s t a n t i a l increase in creep rate of the T i - 2 4 A I - I I N b alloy. It is also s i g n i f i c a n t that the stress level e m p l o y e d in the p r e s e n t study is lower t h a n t h o s e e m p l o y e d by M e n d i r a t t a and L i p s i t t (i). The rationale for the d i f f e r e n c e in c r e e p rate in the p o w e r law regime of the two Ti3AI base i n t e r m e t a l l i c s s t u d i e d by M e n d i r a t t a and Lipsitt il) and that of the T i - 2 4 A I - I I N b of the p r e s e n t study can be found in the m i c r o s t r u c t u r e . The Ti3AI i n t e r m e t a l l i c s e m p l o y e d in the study by M e n d i r a t t a and L i p s i t t (I) c o n s i s t e d only of the o r d e r e d HCP ~2 p h a s e w h e r e a s the T i - 2 4 A I - I I N b alloy c o n s i s t s of a t w o - p h a s e m i x t u r e of o r d e r e d ~2 plus the BCC ~ p h a s e (due to the h i g h e r Nb additions). The i n t e r m e t a l l i c m a t r i c e s are e x p e c t e d to p o s s e s s lower d i f f u s i o n c o e f f i c i e n t s due to t h e i r long range o r d e r e d nature. Thus it is e x p e c t e d that the rate of d i s l o c a t i o n climb will be d e c r e a s e d in the i n t e r m e t a l l i c matrix. Another v a r i a b l e e x p e c t e d to i n f l u e n c e the rate of d i s l o c a t i o n m o t i o n in the i n t e r m e t a l l i c m a t r i x is the a n t i - p h a s e b o u n d a r y energy (APB). D i s l o c a t i o n s moving on s p e c i f i c slip p l a n e s m u s t travel in p a i r s in order to p r e s e r v e the ordered n a t u r e of the matrix. The r e s u l t a n t super d i s l o c a t i o n s m u s t c o n s t r i c t prior to c l i m b thus r e s e m b l i n g e x t e n d e d d i s l o c a t i o n s in low s t a c k i n g fault ener@y materials. G i v e n that the i n t e r m e t a l l i c c o m p o u n d s e v a l u a t e d by M e n d l r a t t a and L i p s i t t (i) w e r e single p h a s e it is not s u r p r i s i n g that the es observed is low c o m p a r e d to the ~s o b s e r v e d in the two p h a s e T ~ A I b a s e alloy Ti-24AI-IINb. The i n c r e a s e in ~s o b s e r v e d in the T i - 2 4 A I - I I N b is s u g g e s t e d to

1934

CREEP OF Ti-24AI-IINb

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23, No. II

arise in part from the contribution of the BCC @ phase to the overall matrix deformation. The similar - ~ H for the Ti-24AI-IINb and the Ti-16 wt% AI-10 wt% Nb (i.) is suggested to arise from the fact that the - ~ H for the Ti-24AI-IINb is expected to be controlled by the most difficult process which is the climb of dislocation in the ordered ~ 2 matrix. It appears from the results presented in this study that alloy modifications to the single phase intermetallic compound matrix to promote enhanced lower temperature ductility will have a profound effect on the high temperature behavior of the material. Conclusions i. No observation of surface microcracking of specimens tested in air was obtained in the present study. More work is needed in order to establish the origin of surface microcracking and its influence on the steady-state creep behavior of Ti-24AI-IINb. 2. The activation energy and stress exponent measured for Ti-24AI-IINb indicate that power law creep operates within the temperature and stress regime evaluated. 3. Without the appropriate diffusion data available it is not possible to define a specific rate controlling mechanism to the creep of the intermetallic compound matrix. 4. The stress level needed to effect power law creep in Ti-24AI-IINb is substantially lower than the stress level needed in the single phase Ti3AI base intermetallics. 5. The presence of the BCC @ phase in the Ti-24AI-IINb alloy is thought to be responsible for the lower stress initiation of power law creep and the accelerated steady-state creep rates observed when compared to the slngle phase Ti3AI base intermetallics. Acknowledqments This research was performed under Metals Technology, Inc. IR&D funding. The material for this program was provided by Dr. J.A. Graves of the Rockwell International Science Center. The author would also like to thank Mr. C.G. Rhodes of the Rockwell International science Center for helpful discussions regarding the issues of surface microcracking of this alloy. References i.

M.G. Mendiratta and H.A.

Lipsitt, J. Mat. Sci., Vol.

15, pp 2985-2990,

(1980) 2. C.G. Rhodes, Rockwell International Science Center, Thousand Oaks, CA., Private Communication, (1989) 3. A.K. Mukherjee, J.E. Bird and J.E. Dorn, Trans ASM, Vol. 62, pp 155-179,

,(1969) 4.

O.A. Ruano, J. Wadsworth and O.D. Sherby, Acta Met., Vol. 36, pp 1117-1128,

(1988)

Vol.

23,

TABLE

No.

I:

ii

CREEP

CHEMICAL

A1 12.6%

Nb 21.2% TABLE

O 0.061% 2:

TEMPERATURE "C 650 704 704* 704 704 760 650 650

.05

COMPOSITION

SUMMARY

OF T i - 2 4 A I - I I N b

OF T i - 2 4 - I i H 42 p p m OF CREEP

I

I

ALLOY

N 60 p p m DATA

INITIAL STRESS (MPa) 103.4 69.4 103.4 103.4 137.8 103.4 137.9 172.35 * T e s t e d in A r g o n

I

1935

I

PERCENT

C 90 p p m

Ti Remainder

FOR Ti-24-ii

STEADY

l

IN W E I G H T

STATE CREEP (HR ) 4 4.80XI09 . 5 0 X I 0 -4 3.33XI0-3 4.08XI0-3 7 . 0 0 X I 0 -3 1.32XI0-3 1.19XI0-3 2.28XI0-3

I

I

RATE

I

.04 760°C a 103.4 MPa .03

704oc

W

650°C .02

.01

.oo 0

I

I

I

I

I

I

I

I

5

10

15

20

25

30

35

40

HOURS

FIGURE I. CrccpcurvesofTi-24Al-llNballoy.

45

CREEP

1936

OF T i - 2 4 A I - I I N b

Vol.

-2.0 - • .1.5

. . . . . . . . . . . .

=

-'

,

,

23,

,

No.

ii

,

"

704" C I

-2.5 -2.0

/

650" C

Ti. 24 AI- 11 Nb &H==259.4 KJ • Mole "1 Q 103.4 MPa A 137.8 MPa

-3.0

.,,# ,2.5 -3.5

............................... i 1.5 2.o 2.5 LOG G HGURE 3. Sums dependenceof the steadystatecreep rate of Ti-24Al-I 1Nb.

3.0

.3.0

1 0"1 850 ]

-3.-~..1 . . . . . . . . . . . . . . . . . . . . . . 11

-~x 104K

I ........ 10

TEMPERATURE (°C) 750 700 650 I' I I

800 I ~

=

600 I

550

ATi~AI + 10 wt% Nb

9 1 0 -2

FIGURE 2, Arrh¢nius plot of the temperaturedependenceof the steady .4ate creep rate of Ti.24A l - 11Nb.

1 0 .3

1 0 .4

1 0 "5

1 0 "6

i 9

I 10

....

I 11

12

I-.x104 T FIGURE 4. Arrhenius plot of the temperature dependence of the steady state creep ofTi-16 wt% AI-10 wt% Nb (reproduced from Mendiratta and Lipsit4