Characteristics of elevated temperature deformation in several L12- modified Al3Ti-based alloys

Intermetdlics 3 (1995) 221-232 0 I995 Elsevier Science Limited Printed in Great Britain. All rights resewed 0966-9795/95/$09.50

Characteristics of elevated temperature deformation in several L12- modified A&Ti-based alloys J. Daniel Whittenberger,‘* K. S. Kumar,bt M. S. DiPietrob & S. A. Brown’ ‘NASA Lewis Research Center, Cleveland, OH, USA bMartin Marietta Laboratories, Baltimore, MD, USA (Received 8 April 1994; accepted 19 June 1994) Replacement of approximately 9 at% of the Al in the DO,, crystal structure compound Al,Ti with Cr, Fe and/or Mn leads to the formation of a potentially more ductile Ll, phase. Utilizing room temperature hardness as a measure of ductility, a series of quaternary alloys containing a total of 9 at% of Cr, Fe and/or Mn in Al-25Ti was produced. Based on the presence of a single phase and minimums in hardness-composition profiles, a quinary and four quaternary compositions were selected for further study, cast as 2 kg heats and forged into pancakes. Small diameter compression test samples were taken from each forging and tested under constant velocity and constant load conditions in air between 900 and 1200 K. The results from this testing indicated that the quinary alloy, A1,,Ti,,Mn,5Cr3Fe,.5, had the best strength; however results from other work (K. S. Kumar and S. A. Brown, Scripta Metall. et Mater., 26 (1992) 197-202) indicates that grain size might be more important than composition in determining elevated temperature deformation resistance. Comparison of the existing creep results for Ll,-modified Al,Ti indicates that larger grain sizes are necessary for strength and that the stress-independent activation energies for deformation are close to that measured for volume diffusion even when grain boundary deformation mechanisms are active. Contrasting the 1100 K creep behavior of the quinary composition to other Ti-Al intermetallics reveals that the current Ll, forms of Al,Ti do not possess a strength advantage over single phase ~1~or y alloys or a lamellar o2 + y mixture. Kev words: Al,Ti, Al,Ti-based,

creep, mechanical properties

behavior when conducting three point bend tests of cast and forged Al,,Ti,,Cr, at 298 and 473 K, subsequent heat treatment of forged A16,Ti25Cr97lead to some 473 K ductility. Lastly, as bend testing of Fe-modified Al,Ti failed to show ductility at room temperature8 or below 923 K,9 little evidence exists which suggests that iron, by itself, will provide low temperature tensile plasticity. Because single element ternary additions of Cr and Mn did result in some ductilization of A13Tibased materials, Kumar et al. ‘O-” reasoned that combinations of these elements with and without iron might be successful. Therefore they melted, cast and homogenized a number of quaternary alloys containing a total of -9 at% of Cr, Fe and/or Mn in a nominal Al-25Ti base. By utilizing the presence of a single phase and room temperature hardness minimums as criteria, a quinary and four quaternary compositions were selected for further study. Results in terms of the microstructures,

INTRODUCTION Replacement of approximately 9 at% of the Al in the D022crystal structure compound AI,Ti with Cr, Fe and/or Mn leads to the formation of a potentially more ductile Ll, phase. While room temperature tensile testing of Al,Ti modified with 8 or 9 at% Mn has revealed some elongation,‘-4 testing of Al,,Ti,,Cr, below 623 K did not result in any ductility. 2,5Zhang et al. ‘, however, have reported limited room temperature bending during four point testing of cast and hot isostatically pressed (HIP’ed) A 166Ti25Cr9, and Schneibel et ~l.,~ recorded room temperature ductility when bend testing polished Al,,Ti,,Cr, samples with low interstitial content. Although Kumar et a1.2did not observe any deviations from linear loaddisplacement *Currently at Max-Planck-Institut fur Metallforschung, Institut fur Werkstoffwis, 70174 Stuttgart, Germany. tNow at Brown University, Providence, Rhode Island, USA. 221

222

J. D. Whittenberger, K. S. Kumar, M. S. DiPietro, S. A. Brown

compressive yield strengths, and influence of heat treatment for these five alloys are reported in Ref. 10, and data from bend and tension testing of these materials can be found in Ref. 11. Unfortunately the hoped for improvements in lower temperature (< 400 K) tensile ductility through quaternary and/ or higher alloying additions were not realized. Concurrent with the bend, tensile and compressive yield strength measurements,““’ the elevated temperature, time dependent flow strength properties of the five quaternary/quinary L12 modified Al,Ti-based materials were also being determined. The purpose of this paper is to document the slow plastic compression characteristics of these materials and to compare their properties to published results for other LIZ modified titanium trialuminides.

EXPERIMENTAL

PROCEDURES

Sufficient master alloys to produce 2 kg ingots of five different compositions (Table 1) were induction melted under an argon atmosphere and cast into 60 mm diameter graphite molds, cropped and homogenized in a two step process under flowing argon: 24 h at 1323 K followed by 96 h at 1473 K. After heat treatment the billets were isothermally forged -70 % at 1323 K into pancakes. Light optical and X-ray examination of each forging revealed that the materials were essentially single phase and had equiaxed grain sizes of the order of 20 prn.‘O,” Further analysis by transmission electron microscopy revealed the presence of occasional rod- or globular-shaped precipitates in several of the materials (for instance A1,,Ti,,Mn,,,Fe,.S as shown in Fig. 3(a) of Ref. 10); however such econd phases were a very minor constituent.

Specimens for chemical analysis were taken from each cast/forged material and examined at both Martin Marietta and the Lewis Research Center. Table l(a) presents the chemistries as originally determined at Martin Marietta, “J’ where Al, Cr, Fe, Mn and Ti for all five alloys were estimated by inductively coupled plasma methods, while the interstitials in three of the alloys were measured by instrumental techniques. Because of the unexpectedly high interstitials levels found in the three alloys, all the materials were completely reanalyzed at the Lewis Research Center (Table l(b)) to verify such results. In this case Al content and the concentrations of all substitutional elements were also determined by inductively coupled plasma. Quantitative chemistries were established through calibration with matrix matched (Al-based) solutions made by mixing single element standards which were formulated to yield a slightly Ti-, Cr-, Fe-, Mn-, etc. rich mixture and a slightly Ti-, Cr-, Fe-, Mn-, etc. poor mixture. Carbon was measured via an infrared detector in combination with thermal analysis. Both oxygen and nitrogen were estimated by an inert gas fusion method with an infrared detector for oxygen and a thermal conductivity cell for nitrogen. The instruments used for all three interstitial elements were calibrated with standards traceable to NBS. At both laboratories the measured compositions for each alloy were normalized to yield a total of 100%; it is these normalized data which are reported in Table 1. Right cylindrical compression samples approximately 5 mm in diameter by 11 mm tall were electrodischarge machined from each forging with the long specimen axis parallel to the forging direction. The slow strain rate deformation behavior of the

Table 1. Chemical compositions (at%) of the cast and forged LIZ-basedAl,Ti alloys Material

Al

C

Cr

Fe

Mn

N

0

Ti

Ni

Si

Zr

_ 0.40 0.19 0.16

24.6 22.9 24.1 23.8 24.7

-

-

-

0.28 0.15 044 0.17 0.11

24.3 23.8 24.3 24.2 24.6

0.2 -

0.6 1.0 1.5 0.4 0.4

0.1 0.5 1.0 0.1 0.7

(a) Martin Marietta Laboratories* Al,,Ti,,Cr,.,Fe, 5 Al,,Ti,,Cr,.,Mn, 5*’ 4,%& +fn, 5 Al,,Ti,,Mn, ,Fe, 3** Al,,Ti,,Mn,.,Cr,Fe,.,”

66.0 67.0 66.8 65.1 65.2

0.17 0.68 0.10

4.9 4.5 0.3 0.2 3.2

4.5 0.1 4.3 2.4 1.7

_ 4.6 4.5 6.7 4.8

_ 0.46 0.40 0.19

(b) NASA Lewis Research Center AlJ&Cr, SFe, s Al,,Ti,,Cr,,Mr+ Al,,Ti,,Fe,.,Mn, 5 A166Ti25Mn67Fe2j Al,,Ti,,Mn, ,Cr,Fe, s

65.3 65.5 63.3 65.5 64.9

0.05 0.04 0.05 0.03 0.05

4.6 4.5 0.3 0.1 3.0

4.4 0.1 4.3 2.6 1.8

4.3 4.0 6.5 4.4

0.42 0.30 0.82 0.27 0.12

* Analyses reported in weight percent in Refs 10 and 11. ** Samples were also analyzed for H; however in all cases the H content was
Elevated temperature deformation in modjied Al,Ti-based alloys

Liz modified Al,Ti alloys was characterized by both constant load and constant velocity compression testing in air between 900 and 1200 K. Constant velocity experiments were conducted in a universal test machine at speeds ranging from 2.12 X 10e3to 2.12 X 10” mm/s. The autographically recorded load-time charts were then converted to true compressive stresses, strains, and strain rates via the offset method and the assumption of conservation of volume. I2 Compressive creep testing was undertaken in lever arm test machines, where deformation was determined as a function of time by measuring the relative positions of ceramic push bars applying a constant load to the specimen.13 Contraction-time data were normalized with respect to the final specimen length and converted into true stresses and strains. The microstructures of selected specimens were characterized by standard light optical microscopy techniques in both the unetched and as-etched states, where the etchant was 5HF-5HNO,-90H,O (parts by volume).

RESULTS Alloy chemistry The chemistries reported in Table 1 in terms of atomic percent show that the desired compositions, as reflected by the chemical formulae used for identification, were obtained. However instances of cross contamination did occur: for instance, a small amount of Fe was present in Al,,Ti,,Cr,,Mn,., and Al,,Ti,,Fe,.,Mn., and Al,,Ti,SMn, ,Fe, 3 contained some Cr. Furthermore the limited scope of the initial chemical analysis (Table l(a)) failed to reveal the presence of Si and Zr which existed in significant amounts in all five materials (Table l(b)): in the worst case, for example, Al,,Ti,,Fe,.,Mn,., had 1.5 at% Si and 1.O at% Zr. Such contamination is not unreasonable, nor unexpected, since all the alloys were melted in crucibles fabricated from silica and zirconia. In general, both chemical analyses yielded Cr, Fe, Mn and Ti concentrations which were in good agreement for each alloy, where with one exception (Ti in Al,,Ti,,Cr, SMn4.5) the measured values are within kO.5 at%. Much larger discrepancies were found in the Al results; however it is likely that the higher Al contents in Table l(a) compared to those in Table l(b) are due to inclusion of Si and Zr in the latter study’s overall alloy chemistry. In terms of the interstitial levels in the three alloys which were examined in common,

223

major inconsistencies exited: for instance, Table l(a) reports 0.68 at% C in A166Ti25Mns.7Fe2.3while Table l(b) gives 0.03 at%, and Table I(a) reports 0.4 at% 0 in A166Ti25Cr4,5Mn4.5 compared to 0.15% in Table l(b). Although both studies indicated large concentrations of interstitials; taken as a whole, the interstitial concentrations (C + N + 0) in Table l(a) are about twice that of Table l(b). The reason(s) for the discrepancies in interstitial contents is not known; however both studies rank the total interstitial content for each material in the same order. Mechanical properties Typical stress-strain curves obtained under constant velocity conditions for the five Ll,-modified Al,Ti alloys are illustrated in Fig. 1. Part (a) demonstrates behavior at 900 K, where continuous work hardening occurs to at least 6% strain, and the stress-strain curves are insensitive to the imposed strain rate. As illustrated in Fig. 1 (a) a factor of 50 decrease in the imposed deformation rate at 900 K has no effect on the strength of Al,,Ti,,Mn,.,Fe, 3. Testing at 1000 K (Fig. l(b)) yields somewhat similar behavior at faster strain rates (>2 X lo-’ s-l); however at slower deformation rates the degree of strain hardening lessens, and the materials become weaker as the strain rate decreases. The strengths of the modified Al,Ti alloys at 1100 K (Fig. l(c)) are quite dependent on the imposed deformation rate, and the stress-strain curves change from continuous work hardening to flow at a more or less constant stress as the strain rate deceases. True compressive strain rate-stress results for the three Liz modified Ti,Al alloys, which were basically only tested under constant velocity conditions, are presented in Fig. 2, where flow stresses (0) and deformation rates (E) were evaluated at 1% strain. Irrespective of the material, the flow stresses at 900 and 1000 K for rates greater than 10m6s-l are essentially independent of the strain rate. However at 1100 K and very slow rates at 1000 K (< 10m6ssi), flow strength decreases with a decreasing deformation rate. Based on the limited 1100 K data for Al,,Ti,,Cr,.,Pe,., (Fig. 2(a)) and Al,6Ti&r,.,Mn+, (Fig. 2(b)) it appears that several time dependent deformation mechanisms might be operating. Comparison of the strength levels for the three materials reveals that the 4.5Cr-4.5Fe quaternary (Fig. 2(a)) is slightly stronger than either the 4.5Cr-4.5Mn (Fig. 2(b)) or 6.7Mn-2.3Fe (Fig. 2(c)) versions.

J. D. Whittenberger, K. S. Kumar, M. S. DiPietro, S. A. Brown 400 -

600 r

9OOK

0

00

$

360 -

0

0 A

I zt

300

-

,’

**

8’

al .> 8 Approximate strain rate,

lOOOK

a /’ A’ --7; I

d’

A

00 K

I’

z f 250 -

,’

0

I’ /

a’ I

200 0

2

4

6

6

10-7

104

(a)1 10-6

10”

103

True compressive strain rate, s-1

True compressive strain, percent

400

r

Approximate strain rate, s-1

‘t-

2oc

cl

2.0x1 0”

100

0

400 350

4 6 2 True compressive strain, percent

a True compressive strain rate, s-1

r

Approximate strain rate,

IJ

2.0x10-4

0

2.0x10-5

300 250 200 150 100

50 0

2

4

6

8

TNe compressive strain, percent

100 IO-7

I

I

10-e

10-6

(4 1 10-O

104

TN~ compressive strain rate, s-1

Characteristic true compressive stress-strain curves for the cast and forged Cr/Fe/Mn-modified Al,Ti materials as functions of temperature and imposed strain rate. (a) A1,,Ti,,Mn,.,Fe,.3 at 900 K, (b) A166Ti25Fe4.5Mn4_5 at 1000 K and (c) A1,6Ti,,Cr,.,Mn, 5 at 1100 K.

Fig. 2. True compressive flow stress-strain rate behavior for (a) Al,,Ti,,Cr, 5Fe4.5, (b) Al,Ti&r,.,Mn,., and (c) Als6TizsMn6.7Fe23 from constant velocity testing as a function of temperature. Solid symbol in part (b) represents a constant load creep test result.

Constant load creep curves for A166Ti,5Fe,.,Mn,.5 and A1,,Ti,,Mn,.,Cr,Fel.5 tested between 1000 and 1200 K are presented in Figs 3 and 4, respectively. The observed creep behavior under all temperature-applied stress conditions was normal, where the instantaneous strain rate decreased throughout the primary regime until a steady state creep

rate was reached. Unfortunately, furnace failures occurred during the highest stress tests at 1000 K for both materials, hence the breaks in the curves in Figs 3(a) and 4(a). For the case of the 250 MPa test of A1,6Ti,SFe,.,Mn,., (Fig. 3(a)), restarting the test yielded a more prolonged primary regime than the initial start; however, as a whole, both

225

Elevated temperature deformation in modiJied AI,Ti-based alloys

300 MPa

/

0.4

0.6

0.8

1.0 0

Time, MS

0.2

0.4

0.6

0.8

1.0

1.2

1.4

1.6

Time, MS 7

r

10

El 0

100

200

300

400

500

I

600

40 80

0

Time, ks

I

I

I

I

I@)\

120 160 200 240 280 320

360

Time. ks

10 9

r

E

0

20

40

60

80

100

120

140

160

Time, ks

$

6

4

7

0

20

40

60

80

loo

Time, ks

Fig. 3. True compressive creep curves for Al,,Ti,,Fe, ,Mn4 5 as a function of the applied stress at (a) 1000 K, (b) 1100 K and (c) 1200 K.

Fig. 4. True compressive creep curves for A1,,Ti,,Mn,,Cr3Fe,., as a function of the applied stress at (a) 1000 K, (b) 1100 K and (c) 1200 K.

creep curves are contiguous with one another. Steady state creep appears almost immediately after the restart of the 300 MPa test for Al,,Tiz5Mn4.SCr,Fe, s (Fig. 4(a)), but the rate is about 25 % faster than the steady state rate measured before the furnace failure. Comparison of the 1000 and 1100 IS creep behavior for A1,,Ti,,Mn4.5Fe,., (Figs 3(a) and (b))

and A16,Ti,,Mn4.,Cr,Fe, 5 (Figs 4(a) and (b)) in conjunction with the results from constant velocity testing at these temperatures (Figs l(b) and (c)) indicate a similar trend, where a strong dependence exists between the creep rate and stress. Furthermore, as demonstrated in Figs 3(c) and 4(c), this behavior is continued at 1200 K for both materials.

J. D. Whittenberger, K. S. Kumar, M. S. DiPietro, S. A. Brown

226

determinations, Ri, and relatively low standard deviations for the stress exponents, 8, in Table 2, eqn (1) adequately represents the slow behavior for plastic strain rate deformation these two Ll,-based Al,Ti alloys. Comparison of A1,,Ti,,Mn,.,Fe,.S to Al,,Ti,,Mn,.,Cr,Fe,., (Fig. 5) reveals that they have about equal strengths at 900 K, while at 1000 K the quinary alloy is stronger than the quaternary, especially below lOA s-l. At 1100 K both materials demonstrate equal properties, but A 1,Ti,,Fe,.,Mn,., has an advantage over A1,,Ti,SMn4.,Cr,Fe,.s at 1200 K. Contrasting the behavior of all five forged L&-based Al,Ti alloys (Figs 2 and 5) reveals that (1) 4.5Cr4.5Mn (Fig. 2(b)) and 6.7Mn-2.3Fe (Fig. 2(c)) versions are much weaker than the other three alloys; (2) within the current limits of testing the 4.5Cr45Fe quaternary has about the same properties as A&T&Fe4, Mn,.,; and (3) this latter alloy and the quinary are approximately equal in strength. Because stress exponents for Al,,Ti,,Mn,.,Fe,., and A1,6Ti,,Mn4.,Cr,Fe,.5 (Fig. 5) exhibit significant changes with test temperature (Table 2), use of a temperature compensated power law (eqn (2)), with B as the pre-exponential constant,

Temperature,

v

1200

Solid symbols denote constant load Open symbols denote constant velocity

10-8

10-7

104

106

104

104

True compressive strain rate, s-1

20’@’ 104

’ 10-8

10-7

106

104

10”

E = Ba*exp

104

(2)

RT

True compressive strain rate, s-1

Fig. 5. True compressive flow stress-strain rate behavior for (a) A1,6Ti,5Fe,.5Mn,,, and (b) A1,,Ti,,Mn,.5Cr3Fe,.5 as a function of test temperature. Open symbols signify constant velocity test results, and solid symbols represents constant load creep data.

Flow stress-strain rate data for A1,,Ti,,Fe4.5Mn4., and Al,Ti,,Mn.,Cr,Fe,., are presented in Fig. 5, where the constant velocity results, taken at 1% strain, are given by the open symbols, and the steady state constant load creep data are represented by the filled symbols. As can be seen in the 1100 K data for both materials, both test methods yield similar flow stress-strain rate values; furthermore this conclusion can also be drawn for A166Ti25Cr4.5Mn4.5 (Fig. 2(b)). A power law equation (eqn (1)) was applied E= Aa”

(-Q)

(1)

at each test temperature to the data where a clear dependency of strain rate on stress exists. The results of this analysis are expressed by the lines in Fig. 5 and are given in terms of the power law constant A and stress exponent II in Table 2(a). Based on the visual fit of the power law curves in Fig. 5 in combination with high coefficient of

to determine stress independent activation energies for deformation Q would not be mathematically valid. Results (Table 2(b)) from multiple linear regression fits for Al,,Ti,,Mn,,Fe,, and Al,Ti,,Mn,.,Cr,Fe,., show this to be the case, as eqn (2) poorly represents the data. For example, the standard deviation for the activation energy, a,, is high for these two materials, and the coefficients of determination for eqn (2) are low in comparison to those &* values (Table 2(a)) for the individual fits at each temperature via eqn (1). Use of a temperature compensated exponential creep equation with the flow stress-strain rate data for Al,,Ti,,Mn,.,Fe,., and Al,,Ti,,Mn,.,Cr,Fe,.,, however, can yield reasonable regression fits, where E =Cexp (Da) exp (-Q)

(3)

RT

with C and D being constants. A listing of the appropriate parameters, measure of fit along with the standard deviations for the stress factor, a,, are given in Table 3. Equation (3) was also applied to the stress-strain rate data in Fig. 2 connected by the dashed lines, and the results of these analyses are presented in Table 3. In several

Elevated temperature deformation in modified Al,Ti-based alloys

227

Table 2. Power law and temperature compensated power law fits of the flow stress-strain rate data for several LIZ-basedAI,Ti alloys (a) Power law Material

Temperature

A166Ti&e4Nn4

1000 1100 1200 1000 1100 1200 1200

5

A16,Ti&ln4.,Cr,Fel5 A16,TizsCr86

Number of data points

Material

AI,,Ti,,Fe,.,Mn,, A166T12,Mn4.,Cr,Pe, 5 Al,,Ti,,Cr,.sMn,., Al,,Ti,,Mn, ,Fez 3 A167Ti2SCr,‘S A163T130 Al,

+6.,

2Ti2,b

*I6 8*I6

A163Ti2@6

sNbl

x ‘20

Al,iTboFe,

&fn

I 9*20

Ti-34Al 27 Ti44Al 27 Ti49Al ”

11 9 5 4 19 9 9 8 8 16 16 16

A

2.43 2.66 1.27 5.59 1.54 1.28 6.06

x x x x x x x

lo-*’ lo-l6 lo-” 10-52 10-14 10m9 lo-l6

(b) Temperature Temperature range (K) (s$

1000-1200 1100-1200 1000-l 100 1000-1100 973-1073 1100-1200 1100-1200 110~1200 1100-1200 980-1130 980-l 130 980-1130

1.79 8.39 1.06 4.73 2.02 1.15 1.56 3.73 4.68 2.13 1.23 6.14

x x x x x x x x x x

n

Ri2

557 4.13 2.54 17.70 3.26 1.54 4.65

0.996 0.969 0,994 0.962 0.987 0.988 0.934

4

0.36 0.43 0.19 3.50 0.19 0.17 0.71 ____

compensated power law

102 lo5 lo-l7 10-s 10-l’ 103 lo-’ 105 lo4 10~3

x lo-’

n

&

% (kJZo1)

&lOl)

3.3 2.8 12.5 7.4 14.7 2-8 2.9 2.7 2.6 5.6 5.1 7.3

335.2 393.0 365.6 296.5 464.2 308.6 312.5 362.3 340.7 342.6 391.7 444-4

0.917 0,928 0.986 0.997 0,947 0.983 0.986 0.99 0.98 0.986 0.976 0.973

044 0.34 1.5 0.5 0.96 0.15 0.14 0.13 0.15 0.20 0.27 0.41

_____

36.3 45.4 36‘5 16.4 31.6 27.4 25.0 30.7 37.6 15.4 19.6 24.0

* XDTM synthesized, powder metallurgy materials.

Table 3. Temperature compensated exponential law fits for several L&-based Al,Ti alloys Material

A1,,Ti2,Fe4..$ln4 5 A1,Tiz,Mn,.sCr3Fe, 5 Al,,Tiz,Cr, sMn4 5 A166Ti25Cr4.5Fe4.5 A1,,Ti,,Mn,.,Fez 3 A167Ti2SCr8’5

Number of data points 11 9 7 5 4 19

Temperature range (K)

1000-1200 1100-1200 1000-l 100 100&l 100 1000-1100 973-1073

(kJko1)

2.16 1.52 4.68 1.08 5.72 2.11

x x x x x x

lo* lo7 lo8 10” lo5 10’4

instances a temperature compensated power law (eqn (2)) could fit the data as well as eqn (3), and in these two cases almost identical activation energies were obtained. For example the power law calculation for Al,,Ti,,Mn,.,Fe,., (Fig. 2(c)) yielded a Q value of 296.5 kJ/mol, as compared to 295.5 kJ/mol from eqn (3). Likewise a temperature compensated power law fit (eqn (2)) of the data dashed lines for bounded by the short A1,Ti,SCr4.,Mn,., (Fig. 2(b)) gave an activation energy of 3655 kJ/mol, as opposed to 366.5 kJ/mol via eqn (3). Therefore good multiple regression fits of the experimental data via eqns (2) or (3) yield virtually identical stress independent activation energies.

0.0267 0.0168 0.0405 0.0480 0.0329 0.0595

353.9 316.2 366.5 451.4 295.5 514.2

Ri2

6D

6, (kJ/mol)

0,963 0.98 1 0,965 0.967 0.999 0.981

0.0023 0~0010 0.0041 0.0092 0.0003 0.0023

24.5 18.9 50.0 61.7 2.1 20.0

Microstructure Typical microstructures found in the cast and forged Al,Ti-based alloys after testing are presented in Fig. 6. On a light optical level A 1,,Ti,,Fe,.,Mn,.S has a relatively high density of second phases (Fig. 6(a)), while little second phase matter could be detected in either Al,Ti,,Mn,.,Fe,., or Al,,Ti,,Mn,.,Cr,Fe,., at a magnification of 500 times. Both A166Ti25Cr4.5Fe4.5 and A166Ti25Cr4.5Mn4.5had similar second phase content; however, as can be seen in Fig. 6(b), the degree was much less then that found in Al,,Ti,,Fe,.,Mn,., (Fig. 6(a)). Furthermore, comparison of the unetched microstructures for each

228

J. D. Whittenberger. K. S. Kumur, hf. S. DiPietro, ,!$.A. Brown

Fig. 6. Typical as tested microstructures in the cast and forged Al,Ti-based alloys. (a) A1,,Ti,,Fe,.SMn4., creep tested at 1000 K and 250 MPa for -I MS to 5.7 14, strain; (b) A166Ti25Cr4.5Mn4_5 tested at 1100 K and -2 X 10e5 s-I to 57% strain: (c) A166Ti25Mn4.5Cr3Fe,_5tested at 1000 K and -2 X 1P5 s-l to 6.7% strain; (d) A1,,Ti,,Mn,SCr,Fe,.S creep tested at 1200 K and 75 MPa for -81 ks to 9% strain; (e) Al,,T&,Fe, sMn, 5 creep tested at 1200 K and 60 MPa for -67 ks to 3-2 %I strain; (f) A&,Ti,,Mn,,Cr,Fe,., creep tested at 1000 K and 300 MPa for -16 MS to 5% strain, The compression axis is horizontal.

alloy subjected to various test conditions failed to reveal any obvious change in the distribution, or amount, of second phases on a light optical level. While neither the chemistry nor crystal structure of the second phases have been established in this study, the observation that A166Ti25Fe4.SMn4.5 has the highest concentration of second phases among the five different alloys correlates with the ‘impurity’ chemistries reported in Table 1 (b). This material has a very high interstitial level C+N+O = 1.31 at%, in conjunction with a Ti+ Zr content greater than 25 at%. Additionally, if Si behaves as Ti, then the overall ‘effective Ti’ level for Al,,Ti*,Fe, SMnG, is much higher and reaches about 27%. Extending this reasoning further,

there appears to be a relationship between relative second phase content and the total interstitial and ‘effective Ti’ levels. The primary factor is the interstitial content, where lower C+N+O correlates with less second phase. The ‘effective Ti’ level, on the other hand, seems to play a secondary role, where at a constant interstitial content more ‘effective Ti’ yields a greater amount of second phases: for example Al,Ti,,Cr,.,Mn, (C+N+O = 0.49; ‘Ti’= 25.3) has more second phase than Al,Ti,,Mn,.,Fe,., (C+N+O = 0.47; ‘Ti’= 24.7). Examination of the etched microstructures after elevated temperature deformation did not reveal any noticeable grain growth in the Liz modified Al,Ti alloys (Figs 6 (c) and (d)). As expected from

Elevated temperature deformation in mod$ed AI,Ti-based alloys

the high Al content in these materials, little oxidative attack occurred even at 1200 K (Fig. 6(e)); on the other hand, essentially all surfaces contained shallow cracks (Fig. 6(e)) irrespective of the test conditions. While most of the observed cracks probably resulted from electrodischarge machining and did not grow significantly during testing, examples of gross cracking were also observed (Fig. 6 (f)) in all alloys under all compressive conditions. Even though such fissuring was not catastrophic (Fig. 4(a)), its appearance is indicative of an inherently brittle material.

et al., 1990 +

10-9

10-a

‘1, ,/’

10-7

’

104

229

n

.

104

&qTi$Ta

104

103

True compressive strain rate, s-1

DISCUSSION

The stress-strain rate plots at 1000 and 1100 K in Fig. 1 for the nominally single phase, -20 pm grain size, cast and forged quaternary/quinary Ll, Al,Ti alloys are similar in appearance to those for the simpler, nominally single phase, -40 pm grain size, cast and forged Ll, ternary alloys modified with 8 Cr or 9 Mn.2,14 Additionally, as typified in Figs 7 (a) and (b), the 1000 and 1100 K flow stress plateaus, which exist over large ranges of strain rate for the five alloys investigated in this work, agree with the results for A1,,Cr2,Crs and Al,Cr,,Mn, in Refs 2 and 14. Sawamura and Fine l5 have also undertaken testing of nominally Al&r&r, (arc melted material with columnar grains: 140 X 590 pm) at several’ constant engineering strain rates between 973 and 1073 K, and their results are presented in Fig. 7, where the plotted compression stress values are twice the shear stresses reported by Sawamura and Fine. The most striking difference in Figs 7(a) and (b) between the results of Sawamura and Fine l5 and those of Brown et aL2,14 for nominally Al&r,@-, is the dependency of flow stress on strain rate, particularly at -1000 K. It is likely that this difference is simply due to the definition of stress in the two studies: Brown et al. used the flow stress at 1% true strain, while Sawamura and Fine utilized the steady state stress between 22 and 38% true strain. Since the 1000 K stress-strain plots for as well as those in this study (Fig. Al&r&r,,2 l(b)), indicate that the rate of work hardening is dependent on the imposed strain rate, evaluation of the flow stress at higher strain values (> 1%) would result in a different dependency on the strain rate. With this idea in mind the strength of Al&r,@-, at -1000 and - 1100 K is about the same in both studies and somewhat less than that found for A1,,Ti2,Mn4,,Cr,Fe,.s (Figs 7 (a) and

400 f-

‘1 ‘\ Sawamura and Fine, 1992. at 1073

L Brown et al., 1990 40@) 104

1

’ 10-7

104

104

I 104

103

10-2

True compressive strain rate, s-1

True compressive strain rate, s-l

Fig. 7. True compressive flow stress-strain rate behavior for several Ll, modified AI,Ti alloys as a function of test temperature. (a) 1000 K, (b) 1100 K and (c) 1200 K.

(b)), Al,,Ti,,Fe,,Mn,., (Figs 7(a) and (b)) or A166Ti25Cr4.5Fe4.5 (Fig. 2 (a)). The 1200 K creep results for 15-20 pm grain size A1,Ti2,Mn4.,Cr,Fe, .5 and Al,Ti,,Fe,.,Mn,., are compared to a XDTM synthesized, powder metallurgy processed A1,.2Ti2,Fe,.,‘6 and a much larger grain size (about 60 pm), low interstitial level A1,,Ti2,Cr,6 in Fig. 7(c). The current quaternary and quinary alloys are stronger than the Fe-modified

230

J. D. Whittenberger, K. S. Kumar, M. S. DiPietro, S. A. Brown

Ll, materiaL6 but at lower strain rates they are considerably weaker than Schneibel et al’s6 larger grain size, low interstitial Cr-modified Ll, alloy. The trends demonstrated by the flow stress-strain rate results in Figs 2 and 5 for the five quaternaryjquinary alloys are in agreement with the 0.2% compressive yield strength-temperature data for these same materials. lo In particular between 900 and 1200 K the yield and flow strengths of the quaternaries decrease with temperature; whereas neither the yield nor flow strengths of the quinary (Fig. 5(b)) are very sensitive to temperature at 900 or 1000 K. Kumar and Brown” also found that AIGGTiZ,Mn&r,Fe,.s and A166Ti,SFe&ln,., were the strongest Ll, modified alloys and that they exhibited similar elevated temperature yield strengths; this observation is reinforced by the current temperature-flow stress-strain rate results (Figs 2, 5 and 7). Surprisingly, Al,,Ti,,Mn~.Q-,Fe,., had little second phase content on a light optical scale, while Al,,Ti,,Fe,.,Mn,, (Fig. 6(a)) contained a significant level of precipitates; therefore, micron size second phases do not lead to significant elevated temperature strengthening. Although such large size precipitates are not detrimental in compression between 900 and 1200 K, it is likely that they would act as crack initiation sites and lead to poor low temperature tensile ductility and toughness. While both Kumar and Brown’s investigation” and this study seem to indicate that the choice of alloying additions required to force the DOZ2 to Ll, phase transformation in Al,Ti are important in determining mechanical strength, this conclusion has been drawn into question by the heat treatment study of Kumar and Brown.” They found that the 300-1200 K yield strength of an annealed, 70 firn Al,6Ti1,Mn,.@-,Fe,.5 was at least 100 MPa less than the small grain, -20 pm form. Further, the yield strength of the heat treated quinary was identical to that of a large grain size (200 pm) Al,,Cr,,Mn,, which, itself, was about 50 MPa weaker than the original 40 pm material. The heat treatment study of Kumar and Brown lo strongly suggests: (1) that the grain size of LIZmodified Al,Ti is an important factor in determining the short term yield stress, where strength decreases with increasing grain size; and (2) that grain size effect could override the influence of alloying. The former contention is corroborated by the results for the XDTM synthesized, powder metallurgy processed A1,.,Ti,,Fe8+ in Figs 7(b) and (c), where this alloy consisted of nominally 20 pm islands of matrix cemented together with a rela-

tively thick mantle (-5 pm) of small matrix grains and A1203 particles. At 1100 K and fast strain rates (~2 X 10e4s’)this aluminide has a flow stress of about 600 MPa,16 which far exceeds the values for the cast and forged materials (Fig. 7(b)). However, such an advantage is not maintained at lower strain rates and the flow stress of Al@ ,Ti,,Fe,., rapidly decreases to about 50 MPa when the strain rate is -2 X lo-‘s“. Presumably this behavior indicates that small grains can provide significant short term strengthening at 1100 K, but the improvement can be lost at slower strain rates. Furthermore, the complete loss of strength for this material at a higher temperature (1200 K in Fig. 7(c)) reconfirms the deleterious role of small grains during creep. Therefore, it must be concluded that weakening occurs by time/temperature dependent grain boundary deformation mechanisms in Ll,-modified Al,Ti. Obviously, for high temperature creep strength one would require ‘large’ grains, and this appears to be born out by the 1200 K results for Schneibel et d’s A167Ti&r8 (Fig. 7(c)). Based on the flow stress-strain rate work to date on Ll,-modified Al ,Ti alloys, the primary dislocation mechanism(s) responsible for creep in large grain sized materials is not clear. While convincing evidence for the formation of subgrains in ternary,‘“” quaternary and quinary” alloys exists, no controlled series of experiments to correlate subgrain size and flow stress have been undertaken to date. Additionally, with the exception of Schneibel et al6 measurement of the stress exponents have not yielded values close to 5. For example, studies ,of Fe-modified16 and (Fe + Mn/Nb)-modified2’ Al ,Ti alloys gave stress exponents of about 3 for powder metallurgy, XDTM processed materials tested at 1100 and 1200 K, while cast and forged materials (Table 2) yield n values as high as 18 at 1000 K, and as low as 1.5 at 1200 K. Although high stress exponents at 1000 K can be rationalized as a dominance of low temperature deformation processes (high work hardening and low recovery rates) and very low exponents at 1200 K can be justified through the prevalence of grain boundary deformation mechanisms, a persuasive argument for climb controlled creep in Ll,-modified Al ,Ti at any temperature can not be presently made. While little specific information on the controlling dislocation mechanism for creep in L12modified Al,Ti can be drawn from the existing stress exponents, a more definitive observation regarding the activation energy for deformation is

Elevated temperature deformation in modjied AI,Ti-based alloys

possible. The results from several studies are listed in Table 2 (b) and Table 3, and with two exceptions Q is of the order of 350 KJlmoI, which agrees with the activation energy for interdiffusion between Al 67Cr25Cr8and Al,Cr,,Mn,. 21 Thus it would appear that elevated temperature slow plastic deformation in these materials is a volume diffusion controlled process, even when strength is dependent on grain size. One of the main reasons for interest in Ll,modified Al,Ti is its low density (-3.9 Mg/m3) in comparison to Ni-base superalloys (-8 Mg/m3). However, other Ti-based intermetallics also have relatively low densities and are capable of being used at high temperatures: for example the densities of both Ti,Al (LY*) TiAl (y) are about 4 Mg/m3. As no one Ti+AI intermetallic has a distinct advantage in terms of density, the eventual application would depend on the properties of each phase such as mechanical strength and oxidation resistance. Figure 8 compares the 1100 K strength of the quinary Ll,-modified Al,Ti to that for a single phase a2 alloy (Ti-34Al), and single phase y material (Ti49AI) and a two phase (cu,+y)Ti-44Al composition in a lamellar microstructure*.22 With the possible exception of the (Ye,it appears that the single phase Ll,-modified Al,Ti does not possesses any strength advantage over y or the lamellar a2+y mixture. This behavior, coupled with the existence of high fracture toughness plus some limited room temperature tensile ductility in TiAl-based alloys’3 and the lack of low temperature plasticity in tension for Ll,-modified Al,Ti, seriously diminishes the prospects for this latter intermetallic. However, the high Al content in the Ll,-modified Al,Ti can lead to the formation of stable, protective alumina scales and materials which display good cyclic oxidation resistance at

231

1473 K.24 In comparison TiAl does not form a protective Al203 scale in air above 1023 K.25 Lastly, it should be noted that a very recent paper by Nakayama and Mabuchi 26indicated that Al,,Ti,,Cr, does possess some room temperature bend ductility in agreement with the earlier work More surprisof Zhang et al.’ on Al,,Ti,,Cr,. ingly, Nakayama and Mabuchi reported that Al 6zTi25Cr13and A 16,Ti&r,4 were even more ductile in bending, with permanent tensile strains to 0.9% being recorded. While they believed that the overriding factor for ductility in the ternary L12modified Al,Ti alloys was the absence of porosity (from homogenization of second phases into the matrix), it is more likely that their bend ductility is a direct result of a low Ti content, possibly combined with a low interstitial content. A number of studies27-30have shown that the room temperature strength of Ll,-modified A 1,Ti increases with the Ti content. Since it is probable that the strength-Ti relationship is due in part to the formation of A12Ti3’,” and/or other second phases such as Ti,AlC’ and Ti2AlN,32 it is also likely that these precipitates can act as crack nucleation sites. Thus, the ideal Ll,-modified Al,Ti with low temperature ductility, oxidation resistance and creep strength equivalent to a2 + y alloys might be achieved in a low interstitial content material with 25 at% (or less) Ti (or equivalent elements such as Zr and Si). Unfortunately, with the possible exception of Nakayama and Mabuchi,2” such alloys have not been produced or tested to date. Furthermore, there has been a decided lack of attention to the alloy chemistries reported in the literature, where usually only nominal content or desired compositions have been given. If Ti level and/or amount of interstitials (C, N and 0) affect the properties, then these must be determined and reported in order to logically evaluate the observed behavior.

SUMMARY

Yang and Wert,

200

400

600

True compressive stress, MPa

Fig. 8. Comparison of the 1100 K creep strength properties of AltiTiz,Mn, &r,Fe,., to those for several Ti-AI materials?2

The deformation resistance of a quinary and four quaternary Ll,-modified Al,Ti alloys have been determined from 900 to 1200 K as a function of strain rate. While the quinary composition, A166Ti2,Mn,.,Cr3Fel.5, appeared to have the best *For this comparisoT2 the 980-l 130 K compression data of were fitted to eqn (2) by multiple reBartholomeuz et al. gression techniques and the appropriate parameters were then used to calculate the strength at 1100 K. The present estimates for the stress exponents and activation energies (Table 2 (b)) agree well with the values for n and Q in Ref. 22.

232

J. D. Whittenberger,

K. S. Kumar, M. S. DiPietro,

properties, an examination of all the known data suggests that grain size might have a greater influence on elevated temperature strength than composition. Furthermore, comparison of the slow strain rate behavior of Al-T1 intermetallics does not reveal any strength advantage for Ll,-modified Al ,Ti over single phase a2 or y or a lamellar ~2~+ y mixture.

REFERENCES 1. Zhang, S., Nit, J. P. & Mikkola, D. E., Scripta Metall. et Mater., 24 (1990) 57-62.

Scripta Metall. et Mater., 24 (1990) 200-06.

15. Sawamura, I. & Fine, M. E., Scripta Metall. et Mater., 26 (1992) 315-18.

16. Dipietro, M. S., Kumar, K. S. & Whittenberger,

J. D., J.

Mater. Res., 6 (1991) 530-38.

17. Kumar, K. S. 8~ Brown, S. A., Acta hjetall. et Mater., 40 (1992) 1923332.

18. Brown, S. A., Kumar,

K. S. & Whittenberger,

J. D.,

Mater. Sci. Eng., A161 (1993) 187-94.

19. Kumar,

K. S., Dipietro,

M. S. & Whittenberger,

J. D.,

Acta Metall. et Mater., 41 (1993) 1379-89.

20. Whittenberger, J. D., Kumar, K. S., Brown, S., Dipietro, M. S. & Farmer, S. C., in Light Weight Alloys for Aerospace Applications II., ed. E. W. Lee & N. J. Kim, TSM, Warrendale, PA, 1991 pp. 327-349. 21. Kumar, K. S. & Whittenberger, J. D., J. Mater. Res., 7 (1992) 104345.

M. F., Yang, Q. & Wert, J. A., Scripta Metall. et Mater., 29 (1993) 389-94. 23. Young-Won Kim, in High Temperature Ordered Intermetallic Alloys ZV, ed. L. A. Johnson, D. P. Pope & J. 0. 22. Bartholomeuz,

2. Kumar, K. S., Brown, S. A. & Whittenberger, J. D., in High Temperature Ordered Intermetallic Alloys IV, ed. L.

A. Johnson, D. P. Pope & J. 0. Stiegler Vol. 213, Materials Research Society, Pittsburgh, PA, 1991 pp. 481-86. 3. Xiaofu Chen, Xiaohua Wu, Shipu Chen Jr Gengxiang Hu, Scripta Metall. et Mater., 26 (1992) 1775-78. 4. Kumar, K. S. & Brown, S. A. Philos. Mag. A, 65 (1992) 91-109.

5. Kumar, K. S. & Brown, S. A. Acta Metall. et Mater., 40 (1992) 1923-32. 6. Schneibel, J. H., Horton, J. A. & Porter, W. D., Mater. Sci. Eng., Al52 (1992) 12631. 7. Kumar, K. S. & Brown, S. A., in High-Temperature Ordered Intermetallic Alloys V, ed. I. Baker, R. Darolia,

J. D. Whittenberger & M. H. Yoo Vol. 288, Materials Research Society, Pittsburgh, PA, 1993 pp. 781-86. 8. Wu, Z. L., Pope, D. P. & Vitek, V., Scripta Metall. et Mater., 24 (1990) 2187-90.

K. S., Brown, S. A., Montoya, K., Dipietro, M. S. & Whittenberger, J. D., in Proceedings of International Symposium on Intermetallic Compounds-Structure and Mechanical Properties- (JIMIS-6)) ed. 0. Izumi, Japan

Institute of Metals, Sendai, Japan, 1991 pp. 667-71. 10. Kumar, K. S. & Brown, S. A., Scripta Metall, et Mater., 26 (1992) 1977202.

11. Kumar, K. S., Dipietro, M. S., Brown S. & Whittenberger, J. D., NASA TM 105628, 1992. 12. Whittenberger, J. D., Mater. Sci. Eng., 57 77 (1983). 13. Whittenberger, J. D., Buzek, B. C. & Wirth, G., J. Mat. Sci., 21 (1986) 923330.

K. S. & Whittenberger,

Stiegler, Vol. 213, Materials Research Society, Pittsburgh, PA, 1991 pp. 777-94. 24. Parfitt, L. J., Smialek, J. M., Nit, J. P. & Mikkola, D. E., Scripta Metall. et Mater., 25 (1991) 727-31.

25. Meier, G. H., Birks, N., Pettit, F. S., Perkins, R. A. & Grabke, H. J., in International Symposium on Structural Intermetallics, ed. R. Darolia, J. J. Lewandowski, C. T. Liu, P. L. Martin, D. B. Miracle & M. V. Nathal, TMS, Warrendale, PA 1993, pp. 861-77. 26. Nakayama, Y. & Mabuchi, H. Intermetallics, 1 (1993) 4148.

27. Wu, Z. L., Pope, D. P. & Vitek, V., Scripta Metall. et Mater., 24 (1990) 2191-96.

28. Winnicka,

M. B. & Varin,

R. A. Scripta Metall. et

Mater, 25 (1991) 229772302.

9. Kumar,

14. Brown, S. A., Kumar,

S. A. Brown

J. .D.,

29. Wu, Z. L., as reported in George, E. P., Pope, D. P., Fu, Cl. & Schneibel, J. H., ISIJ International, 31 (1991) 1063-75. 30. Morris, D. G., Gunther, S. & Lerf, R., in High-Temperature Ordered Intermetallic Alloys V, ed. I. Baker, R. Darolia, J. D. Whittenberger and M. H. Yoo, Vol. 288, Materials Research Society, Pittsburgh, PA, 1993 pp. 177-82. 31. Morris, D. G. & Gunther, S., Acta Metall. et Mater., 40 (1992) 3065-73.

32. Wu, Z. L. & Pope, D. P., in International Symposium on Structural Intermetallics, ed. R. Darolia, J. J. Lewandowski, C. T. Liu, P. L. Martin, D. B. Miracle & M. V. Nathal, TMS, Warrendale, PA 1993, pp. 107-16.