c~ o _ c~/, = erf [21b]

The interfacial compositions are obtained from the condi- tions of continuous flux and activity as:

c ~° + c ~ ° ~ / D ~ c ~/t~ = [22a]

1 +

c ~/'~= (f"/ft3)c"/tJ [22b]

where f " and f~ denote the activity coefficients in the two parts of the weld. Disregarding the printing error in Stark's paper, pointed out by Ku~era and Str~nsk~, 7 we find that Eqs. [21] and [22] are identical to the ones derived by Stark.

This approach not only allows the use of different activity coefficients but also different diffusion coefficients in the two parts of the weld. Since only exact solutions to Fick's second law are applied, the mass balance is ensured.

Figure 1 shows the concentration profile in a compound material w i t h f ~ / f ~ = 1.65 calculated using Eqs. [21] and [22]. This diagram should be compared with Figures l(a) and (b) in Million's report. The dashed curve shows the relative activity, calculated from the concentration profile. It is interesting to notice that, although the activity curve is continuous, its slope is not. Of course, in view of Eq. [17], this is the only possibility for maintaining a continuous flux when Oac/OCc is discontinuous.

As a concluding remark, one may add that in approximate calculations one should use Eqs. [21] and [22]. If the dif- fusion of the alloy elements cannot be neglected, one should try some diagonalization procedure (see Reference 6), and for the more complicated case one has to apply nu- merical methods and solve Eqs. [10] or [11] (see Refer- ences 8 and 9).

REFERENCES

1. J. P. Stark: Metall. Trans. A, 1980, vol. l lA, p. 1793. 2. J. P. Stark: Metall. Trans. A, 1980, vol. l lA, p. 1797. 3. L. S. Darken: Trans. A1ME, 1949, vol. 180, p. 430. 4. B. Million: Z. Metallkde, 1983, vol. 74, p. 105. 5. Y. Adda and J. Philibert: La Diffusion dans les Solides, Paris P. U. F.,

1966, chapter 5. 6. P.K. Gupta and A. R. Cooper, Jr.: Physica, 1971, vol. 54, p. 39. 7. J. Kuc~ra and K. Strfinsk~,: Metall. Trans. A, 1982, vol. 13A, p. 1658. 8. J. /~gren: Scand. J. Metallurgy, 1981, vol. 10, p. 134• 9. J. /~gren: Scand. J. MetaUurgy, 1982, vol. 11, p. 3.

Creep Deformation of TiAI and TiAI + W Alloys

PATRICK L. MARTIN, MADAN G. MENDIRATTA, and HARRY A. LIPSITT

Intermetallic compounds based upon TiA1 and Ti3A1 ex- hibit attractive combinations of physical and mechanical properties for use at elevated temperatures..The main detracting feature is their lack of tensile ductility at low temperatures. TiAI (referred to as y) has the Llo structure with only a 2 pct deviation from cubic symmetry. Ti3A1 (referred to as a2) has the DOl9 superlattice of the hcp structure. Alloys containing less than 52 at. pct A1, such as those described in this paper, are two-phase mixtures of y and a2 usually coexisting in a lamellar morphology. 1.2

The slip of TiAI has been studied both in compression 1 and in tension. 3 Dislocations are observed to be the same as in fcc crystals except that two-thirds of the {111} (110) slip systems require the passage of superdislocations to main- tain long-range order. Twins of the type {112} (111) have also been observed. 1'2 These features combine to give (single-phase) TiA1 a sharp brittle-to-ductile transition at about 700 °C.3 In contrast, Ti3A1 shows a continuously • increasing tensile ductility with increasing test temperature. 4 The dislocations observed in this phase are of the basal (so-called "a") type with some "c + a" type being ob- served at 700 °C. 5'6 The study of creep behavior in these alloys has, thus far, been restricted to a2-base alloys; 7 this study revealed that a2 alloys obey the conventional power- law creep equation

/~min = (~r/E)" e x p ( - Q a / R T ) [1]

in the temperature range 550 to 825 °C.

The present study characterizes the creep-deformation modes in two al loys--binary TiA1 and ternary TiA1 + W. These alloys consist of at least two phases--predominantly y and a small volume fraction of a2. In addition, the aged TiA1 + W alloy contains a third phase whose morphology is characteristic of particle-strengthened alloys. Correlation of the minimum creep rates with the parameters in Eq. [1] allows determination of the activation energy and stress exponent. A simple explanation is proposed to account for the different creep rates observed in the two alloys.

Experimental Procedures. The alloys used in this study were produced by powder-metallurgy techniques. Pre- alloyed ingots were converted to powder via the Rotating Electrode Process (REP) by Nuclear Metals, Inc. The -325 mesh size fraction was consolidated by hot extrusion at 1413 °C with a 26:1 reduction in area. The chemical com-

positions of the consolidated alloys are shown in Table I. Mechanical test specimens were prepared by low-stress

PATRICK L. MARTIN, formerly with Carnegie-Mellon University, Pittsburgh, PA, is a Scientist with Los Alamos National Laboratories, Los Alamos, NM. MADAN G. MENDIRATTA is a Senior Scientist, Systems Research Labs., Inc., Dayton, OH. HARRY A. LIPSITT is a Group Leader, Air Force Wright Aeronautical Laboratories, Wright-Patterson Air Force Base, OH 45433.

Manuscript submitted November 15, 1982.

U. S. GOVERNMENT WORK 2170 VOLUME 14A, OCTOBER 1983 NOT PROTECTED BY U. S. COPYRIGHT METALLURGICAL TRANSACTIONS A

Table I. Compositions of Alloys Studied in Atomic Percent

Alloy Titanium Aluminum Tungsten

Binary 49.7 50.3 - - Ternary 49.1 48.7 2.2

{a)

(a)

(b) Fig. 1 --Light micrographs showing structure of (a) TiAl, (b) TiAI + W,

crush grinding to final form. Heat treatments were carried out in a vacuum of 5 × l0 -5 torr or better.

Testing was carried out on constant-load-type creep frames using LVDT extensometers to monitor and record the strain continuously. Temperature was controlled to -+2 °C in the range 700 to 950 °C. The initial and deformed microstrnctures were characterized by light and electron microscopy. The latter method made use of a JEOL 120 CX STEM.

Results and Discussion. Both alloys received a two-stage recrystallization heat treatment consisting of two hours at 1150 °C, followed by eight hours at 950 °C. Light micro- graphs of the resulting microstructures are shown in

(b)

Fig. 2 - - Transmission electron micrographs showing phase morphology in TiA1 + W. (a) Bright-field micrograph showing lamellar 3' + a2 micro- structure, third-phase particles, and lamella-free 3' grains, (b) dark-field micrograph showing a2-phase lamella.

Figure 1. The binary-alloy microstructure is not signifi- cantly different from that reported earlier, 2 in that a non- uniform distribution of a2 is observed. Some grains contain a2 laths, while in others a2 allotriomorphs decorate the 3' grain boundaries. The W-bearing ternary alloy contains both equiaxed-3' grains and a large proportion of lamellar areas. The lamellar regions consist of alternating-twin related-3' laths with an occasional (thin) a2 lath interdispersed. This is illustrated in the transmission electron micrographs of Figure 2. It can also be seen from Figure 2 that a third phase has nucleated on all internal surfaces and within some of the equiaxed-3' grains. The STEM X-ray spectrum comparing these precipitates to the 3' matrix indicates that these pre- cipitates are enriched in W with respect to the 3' phase (Figure 3). It is believed that these precipitates are the bcc /3 phase stabilized by the excess W.

METALLURGICAL TRANSACTIONS A VOLUME 14A, OCTOBER 1983--2171

Fig. 3--Superimposed STEM X-ray spectra of third-phase particles and matrix. Particle spectrum is shown by white line.

Table II. Creep Deformation Conditions and Resulting Minimum Creep Rates

10-2 -

t d 3

,,c

10"4

°'~LI

10 -5

Test Temp. Initial Stress Alloy (°C) (MPa) ~mi~ (hr -1)

TiA1 700 103.4 1.98 X 10 -6

206.8 2.4 x 10 5 241.3 5.2 × 10 -5

750 103.4 9 × 10 .6 206.8 1.2 )< 10 -4

241.3 3 x 10 4 iO-2 --

800 103.4 5 x 10 5 206.8 7.2 x 10 -4 241.3 0.84 × 10 -3

850 103.4 1.1 x 10 -4 206.8 3.2 × 10 .3 241.3 5.3 X 10 -3

TiAI-W 750 172.4 1.1 x 10 -5 1°-3

800 172.4 4.5 x 10 -5

850 137.9 4.8 x 10 -5 172.4 2.1 X 10 4

241.3 6.8 X 10 -4

310.3 4.1 x 10 .3 ,.~

900 172.4 9.2 × 10 .4 ~ 10-4

950 172.4 5.5 X 10 3 .~

The alloys were creep deformed at temperatures from 700 to 950 °C and at nominal stresses from 138 to 310 MPa. The minimum creep rates measured for both alloys are shown in Table II. A graphic comparison of these two compositions is made in the Arrhenius plot of Figures 4 and 5, where the logarithm of the minimum creep rates are plotted against reciprocal temperatures (at constant stress). Two points are immediately clear upon examination of these data. First, the minimum creep rates of the W-bearing ternary alloys are lower than those of the binary alloy over the temperature regime studied. Second, the slope of these curves, corre- sponding to the apparent activation energy for creep, Qa, is similar but not identical for the two alloys. The magnitude

IO "6 l I I 8 9 I 0

4 I / T X I 0 , *C

Fig. 4--Temperature dependence of minimum creep rate for TiAI.

TIAI

22 Oa = 3 X I 103MPa

(~72 kca l /mole )

I II

10-5

T I A I - W X

7 2 . 4 MPa

Qo = 3.7 X 105 J/mole (~89 kcol/mole)

N'

10-6 I I I I 8 9 I0 II

I /T X 10 4, *C

Fig. 5--Temperature dependence of minimum creep rate for TiAI + W.

2172--VOLUME 14A, OCTOBER 1983 METALLURGICAL TRANSACTIONS A

¢.-

it1 o l

° ~

10 -3

850"C

10-4

io-S

10 -6

800"C

,%

, w

-N

i 0.C / / / OOO

I I I I

I00 200 300 400

o-, MN/m 2

Fig. 6 - -S t ress dependence of minimum creep rate for TiAI.

of Qa for the binary is 3 × 105 J/mole, while that of the W-bearing ternary is 3.7 × 105 J/mole. These activation- energy values are greater than normally found in many conventional titanium alloys, 8 probably due to the relative difficulty of diffusive processes in ordered alloys.

Figures 6 and 7 compare the two alloys on a plot of minimum creep rate v s logarithm of the applied stress (at constant temperature). The slope of the curves is the stress exponent, n. The value for the binary alloys is - 4 , while that of the ternary is slightly higher at 5.5. In this range of values, recovery is the assumed creep mechanism?

The lower creep rate of the ternary alloy compared to the binary is consistent with the observed microstructural differ- ences. The optical micrographs in Figure 1 clearly show the effective grain size to be smaller for the ternary alloy. This normally leads to a higher minimum creep rate due to grain- boundary sliding. However, the nucleation of the W-rich precipitates on these boundaries may inhibit this mode of deformation. The smaller grains thus reduce the effective slip length. The fine dispersion of W-rich particles within the equiaxed-7 grains may also inhibit the passage of dis-

10 -4

i0 -3

5.5

,0 .5 % ' I00

10-2 - TI AI io-2 -

I I F

200 3OO 400

TiA I - W 850"C

0"~ M N / m 2

Fig. 7- -St ress dependence of minimum creep rate for TiAI + W.

Fig. 8--Transmission electron bright-field micrograph showing dis- locations interacting with third-phase particle.

locations. Figure 8 shows these particles to interact with the dislocations such that either Orowan bowing and/or climb processes are necessary to effect bypassing. Single-unit dis- locations as well as superdislocations are shown in Figure 8. The high antiphase energy of the y lattice prevents resolu- tion of the two-unit dislocations making up the super- dislocations. The relative increase in the value of n in the ternary alloy is much less dramatic than that normally noted in (nondeforming) dispersoid-strengthened alloys where values approaching 20 have been reported. 9 The small volume fraction, variable distribution, and potential deformability of the W-rich precipitates observed in this

METALLURGICAL TRANSACTIONS A VOLUME 14A, OCTOBER 1983--2173

work make the direct analogy to dispersion-strengthened alloys difficult.

(1) The engineering creep data for the two ordered alloys can be expressed in terms of the conventional power law. Qa valuesof3 x 105and3.7 x 105 J/moleandn values of 4 and 5.5 were observed for TiAI and TiAI + W, respec- tively. (2) The creep activation energy was found to be higher for these ordered alloys than for conventional Ti alloys, which may be due to the characteristically slower diffusion in ordered matrices. (3) The slightly higher Qa and n values for the W-containing ternary may be due to W-rich particles inhibiting dislocation motion and grain- boundary sliding.

REFERENCES

1. D. Shechtman, M.J. Blackburn, and H.A. Lipsitt: Metall. Trans., 1974, vol. 5, p. 1373.

2. S. M. L. Sastry and H.A. Lipsitt: Metall. Trans. A, 1977, vol. 8A, p. 299.

3. H.A. Lipsitt, D. Shechtman, and R.E. Schafrik: Metall. Trans. A, 1975, vol. 6A, p. 1991.

4. H.A. Lipsitt, D. Shechtman, and R.E. Schafrik: MetaU. Trans. A, 1980, vol. IlA, p. 1369.

5. S. M. L. Sastry and H.A. Lipsitt: Metall. Trans. A, 1977, vol. 8A, p. 1543.

6. S. M. L. Sastry and H.A. Lipsitt: Acta Met., 1977, vol. 25, p. 1279. 7. M.G. Mendiratta and H.A. Lipsitt: J. Mater. Sci., 1980, vol. t5,

p. 2985. 8. M. Kehoe and R.W. Broomfield: Titanium Science and Technology,

R. I. Jaffee and H. M. Butte, eds., Plenum Press, New York, NY, 1973, vol. IV, p. 2167.

9. S. Purushothaman and J.K. Tien: Acta Met., 1978, vol. 26, p. 519.

Effects of Second-Phase Dispersoids on Deformation Behavior of AI-Li Alloys

K.K. SANKARAN, J. E. O'NEAL, and S. M. L. SASTRY

A1-Li alloys, strengthened by coherent precipitation of the ordered 6' phase (A13Li, L12 structure), ~'2 possess good combinations of high specific modulus and high specific strength. 3'4 These alloys, however, deform by planar slip and hence possess low ductility, 5'6 and they are difficult to cast and process by conventional ingot-casting techniques. 7 The ductility of A1-Li alloys can be improved by the slip homogenization effected by grain refinement and the intro- duction of incoherent, nonshearable dispersoids. 8'9 While attempts to obtain a uniform distribution of fine, incoherent dispersoids in A1-Li alloys by ingot-casting methods have not been successful, ~° rapid solidification processing of these alloys can overcome the difficulties of ingot process- ing, refine the grain size, and increase the solid solubility

K.K. SANKARAN, Design Specialist, is with McDonnell Douglas Astronautics Company, P.O. Box 516, St. Louis, MO 63166. J.E. O'NEAL, Scientist-Electron Microscopy, and S. M. L. SASTRY, Senior Scientist-Physical Metallurgy, are both with McDonnell Douglas Research Laboratories, P.O. Box 516, St. Louis, MO 63166.

Manuscript submitted July 23, 1982. J

of potential dispersoid-forming elements, which can subsequently be precipitated as fine dispersoids. 4'8'9

The purpose of this investigation was to determine the effectiveness of microstructural refinement and the fine in- coherent dispersoids resulting from minor additions of the dispersoid-forming elements Mn, Cr, Fe, Co, Ti, and Zr for homogenizing slip and enhancing ductility in rapidly solidi- fied binary A1-Li alloys. Particular emphasis was placed on the deformation and strengthening mechanisms in solid solution-, precipitation-, and dispersion strengthened A1-Li alloys.

Ihgots of AI-Li alloys, with the nominal compositions listed in Table I, were prepared by melting 99.99 pct pure A1 with the appropriate master alloys and casting in an inert atmosphere. The dispersoid forming elements listed in Table I have low equilibrium solid solubilities in aluminum and form aluminum rich compounds. These elements, there- fore, are good candidates for producing supersaturated solid solutions by rapid solidification, and large volume fractions of controlled dispersions by suitable heat treatment of the rapidly solidified alloys. Zr was added to produce subgrain strengthening in A1-Li alloys. Additions of the dispersoid- forming elements were made in the form of commercially available master alloys. Li additions were made in the form of A1-10Li master alloys, which were prepared from 99.999 pct pure A1 (<2 ppm Na) and high-purity Li (<70 ppm Na).

The ingots were remelted and rapidly solidified at rates ~104 °C per second by roller-quenching into flakes. 8 The flakes were consolidated by compacting into 75 pct dense billets in 6061-AI cans, degassing the billets while pro- gressively heating to the extrusion temperature of 400 °C, and extruding the billets at a reduction-of-area ratio of 20:1 into 11.5 mm diameter cylindrical rods with a core diameter of 10 mm. The chemical analyses of the extrusions are listed in Table I.

All extrusions were heat treated by solution treatment at 540 °C for 0.5 hour, cold-water quenching, and aging at 200 °C for 24 hours.

Microstructures of the Alloys. Figure 1 presents the mi- crostructures of typical extruded alloys following solution treatment and shows the dispersoids that precipitate during elevated-temperature processing. Because of the relatively low volume-fraction of the dispersoids, selected-area dif- fraction patterns could not be obtained to determine their identity. Since dispersoid precipitation occurs essentially during the elevated-temperature solution treatment of the alloys, it can be assumed that these are the phases that occur in equilibrium with the primary Al-solid-solutions in the corresponding binary systems. The dispersoid parameters determined from transmission electron micrographs of the alloys are listed in Table II. The dispersoid parameters are accurate to ---5 pct of the indicated values. Subsequent aging of the alloys results in the precipitation of 3', and the microstructure consists of a bimodal distribution of 6' and dispersoids such as shown for the AI-Li-Co alloy in Figure 2. The size of 6' precipitates following aging at 200 °C for 24 hours is about 40 nm.

Room-Temperature Tensile Properties and Deformation Behavior. The room-temperature tensile properties of AI-Li alloys are presented in Table III. The high strength of A1-Li-Zr alloys is caused by the presence of fine 1 to 2/xm

2174-- VOLUME 14A, OCTOBER 1983 METALLURGICAL TRANSACTIONS A