Intermetallic NiAlTa alloys with strengthening Laves phase for high-temperature applications. I. Basic properties

PII: SO966-9795(97)00031-9

Intermerallics 5 (1997) 563 -577 Cc; 1997 Elsevier Science Limited Printed in Great Britain. All rights reserved 0966.9795/97/$17.00 + 0.00

ELSEVIER

Intermetallic NiAl-Ta alloys with strengthening Laves phase for high-temperature applications. I. Basic properties B. Zeumert & G. SauthofF Max-Planck-Institut fir Eisenforschung GmbH, D-40074 Dtisseldorf, Germany

(Received 11 February

1997; accepted 20 March 1997)

Various Ta-containing NiAl-base alloys with the Laves phase TaNiAl, with Cl4 structure, as strengthening second phase were prepared by ingot metallurgy. They were studied with respect to the deformation behaviour at ambient and high temperatures - including elastic deformation and creep - as a function of alloy composition and microstructure. NiAl dissolves up to 0.2 at% Ta in solid solution. NiAl with up to 3 at% Ta forms precipitate particles of Laves phase primarily on grain boundaries, whereas the Laves phase covers the grain boundaries completely to form a continuous skeleton for higher Ta contents. The observed strengthening of NiAl by Ta is a result of both solid-solution strengthening and second-phase strengthening. The latter effect increases with increasing fraction of Laves phase and is accompanied by an increasing brittle-to-ductile transition temperature. The oxidation behaviour was checked. These alloys are promising for applications at temperatures above those of the Ni-base superalloys. 0 1997 Elsevier Science Limited Key words: A. nickel aluminides based on NiAl, Laves phases, B. mechanical properties at ambient temperatures, mechanical properties at high temperatures.

1. INTRODUCTION

The present investigation, based on the previous work on Laves phase alloys, was centred on intermetallic NiAl alloys with Ta, in which the ternary Laves phase TaNiAl with hexagonal Cl4 structure may form. It is noted that Ni and Al occupy the same kind of lattice sites in this Laves phase, and indeed the Ni/Al ratio can be varied within broad limits according to the ternary phase diagram.” This mutual substitution of Ni and Al in the ternary Laves phase corresponds to the formula Ta(Nio.5_,Alo.5+x)~ with -0.5 < x < + 0.5, which corresponds to the behaviour of other similar ternary Laves phases. lo The objectiveSof the present work was to study the effects of macroalloying, microalloying, and processing on the deformation behaviour of NiAl alloys with strengthening Laves phases, so as to establish a basis for the development and optimisation of structural alloys for high-temperature applications. Part I reports the behaviour of the basic NiAl-Ta alloys. The effects of further alloying are the subject of Part II, while Part III deals with the effects of processing. Some selected results have been reported mearlier.‘2-‘4A complete report with details of the work of Parts I and II is available in Ref. 15. This work has been

Various intermetallic phases have been selected as bases for materials developments aimed at structural applications at high temperatures. One of the most promising phases is the cubic ordered intermetallic phase NiAl, with B2 structure; however, this suffers from very low deformability at lower temperatures and comparatively low strength at higher temperatures, in particular above 1000”C.‘~2 As has been shown before (for example, in Refs 3 and 4) higher strength at high temperatures with still tolerable brittleness at low temperatures can be achieved by adding harder second phases, and the developments in progress make use of this.5 Laves phases - in particular ternary Laves phases with Al - that show outstanding high strengths at high temperatures with, however, brittle-to-ductile transition temperatures of the order of 0.6 T, (T, = temperature of melting), may be used for strengthening B2 phases, in particular NiAl.“‘O

*To whom correspondence should be addressed. tPresent address: McKinsey & Co, Inc., Dusseldorf, Germany. 563

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paralleled by work elsewhere on eutectic lamellar NiAl-Nb and NiAl-Ta alloys with the NiAl phase together with the Laves phase, prepared by directional solidification and studied with respect to creep and toughness.i”20

2. EXPERIMENTAL 2.1 Alloy preparation On the basis of the findings of previous work,iO study was centred on NiAl-Ta alloys that contain the ternary Laves phase TaNiAl with hexagonal Cl4 structure as well as the B2 phase NiAl. The constitution and phase equilibria in the alloy system Ni-Al-Ta were studied separately. ii The alloys studied are listed in Table 1. The alloys were prepared by ingot metallurgy. The purities of the metals used, Ni, Al, Ta, were 99.95, 99.97, 99.8 wt%, respectively. The alloys were melted in a vacuum induction furnace and cast in a cold Cu crucible. The resulting alloys contained pores, cavities, and microcracks. The alloys were then remelted in a Bridgman furnace and solidified directionally (with a temperature gradient of 40-45 K cm-’ in the temperature range 1500-1600°C). After this the alloys contained practically no pores nor microcracks. The composition deviations from the nominal alloy compositions were found to be less than ho.5 wt% for Ni, I.1 wt% for Al, and f O-3wt% for Ta. The impurity content was 980 at. ppm Fe, 140 at. ppm C, 25 at. ppm S, 8 at. ppm N, and 181 at. ppm 0, on average. The alloys with high contents of Laves phase with correspondingly high melting temperatures were melted in a magnetic levitation furnace and cast in a heated ceramic crucible, with a cooling time of 17 h from 1300°C to room temperature. These alloys were free of pores and cracks.

ll0.

Ni (at%)

Al (at%)

Ta (at%)

T-O T-0.2 T-0.5 T-l T-3 T-5 T-10 T-14 T-15 T-20 T-25 T-34(L)

50.0 49.9 49.75 49.5 48.5 47.5 45.0 43.0 42.5 40.0 37.5 33.0

50.0 49.9 49.75 49.5 48.5 47.5 45.0 43.0 42.5 40.0 37.5 33.0

0.2 0.5 I.0 3.0 5.0 10.0 14.0 15.0 20.0 25 34.0

The phase compositions were determined by energydispersive analysis (EDX) by scanning electron microscope (SEM) and wavelength-dispersive analysis (WDX) by electron microprobe. Phase compositions at fracture surfaces of specimens cracked in situ were determined by Auger electron spectroscopy (AES). Alloy crystal structures were studied by X-ray diffraction of annealed powders at room temperature. Textures were measured at the Institut fur Metallkunde und Metallphysik of the RWTH Aachen with the help of Dr J. Ball. The grain structure and phase distribution were analysed by light microscopy. Grain boundaries were etched by a solution of 10 vol.% concentrated HNOS and 90 vol. % ethyl alcohol, while phase contrast was achieved by etching in a solution of 68 vol.% glycerine, 16 vol.% HF and 16 vol.% concentrated HN03. Crack frequencies were determined by lineal analysis. Relative crack frequencies refer to the observed maximum frequency for each case. Microstructure was studied by transmission electron microscopy (TEM). Thin foils were prepared by electropolishing with a jet polisher (20 V, 60 mA) at 5°C in a solution of 68 vol.% concentrated acetic acid, 16 vol. % perchloric acid, and 16 vol. % ethylen glycol monobutyl ether and by ion-beam thinning. Thinning of Laves phases was extremely difficult because of premature fracture of the brittle foils. Densities were measured pycnometrically at room temperature. Thermogravimetric oxidation testing was done by Dr H. J. Schmutzler in the MaxPlanck-Institut fur Eisenforschung. Dynamic Young’s moduli and average thermal expansion coefficients in the temperature range 23-1000°C were determined by Dr V. Thien at Siemens KWU, Mtilheim. 2.3 Mechanical testing

Table 1. Composition of alloys studied Alloy

2.2 Alloy characterisation

Laves phase (vol.%)

0.5 2 8 14 25 49, eutectic 49 75 83 196

Testing was done at constant temperature with temporal and local temperature fluctuations in the specimen of less than &3’C below 800°C and f 1°C above 800°C. Stress-strain behaviour was studied using compression tests with a constant deformation rate between room temperature and 1400°C and using tension tests at 1100°C. The compressive creep behaviour was studied by creep tests with constant load or stepwise increased loading for determining the stress-creep rate relation with single specimens in the temperature range 800-1300°C. Specimens with dimensions

NiAl-Ta

5 x 5 x 10mm3 were prepared by electrodischarge machining with subsequent parallel polishing of the end surfaces. The measurement errors were about *O-05% for strain and +0.02% for stress. Additional tensile creep tests were done in the temperature range 900-1000°C. Ductility was studied by four-point bending tests with a constant strain rate between room temperature and 1400°C. Specimens with dimensions 3 x 6 x 18 mm3 were cut and polished carefully to avoid premature cracking. The measured flexural strains and stresses refer to the outer specimen fibre, assuming linear elasticity. The brittle-to-ductile transition temperature (BDTT) was defined as the temperature above which 1% plastic strain in the outer fibre could be attained. Specimens were notched 1.5mm for the determination of the

(4

565

alloys with Laves phase

flexural fracture toughness parameter I& in the temperature range 80&13Oo”C. Crack frequencies were measured by lineal analysis on specimen sections.

3. RESULTS 3.1 Constitution

NiAl-Ta alloys are single phase for Ta contents below 0.2at% Ta. With higher Ta coritents, up to about 3 at% Ta, the ternary Laves phase TaNiAl with hexagonal Cl4 structure is precfpitated primarily as particles on grain boundaries (see Fig. l(a)). The average diameter of thiese particles is about 2pm, and the particle number increases with increasing Ta content. With Ta contents

(b)

(4 Fig. 1. Micrographs of alloys T-l with 2 vol.% Laves phase (a), T-10 with 25 vol.% Laves phase (b), T-14 with’49 vol.% phase (c), and T-20 with 75 vol.% Laves phase (d) with the light Laves phase TaNiAl and the dark NiAl phase.15

Laves

B. Zeumer, G. Sauthof

566

above 3 at% the Laves phase precipitate covers the former NiAl grain boundaries completely, thus forming a continuous Laves phase skeleton with discontinuous NiAl grains surrounded by Laves phase (Fig. l(b)). In addition, small Laves phase particles with average diameter of the order of 300nm were observed in the NiAl grains. The eutectic composition is about 14at% Ta with about equal volume fractions of NiAl and Laves phase (Fig. l(c)). The eutectic temperature is about 1500°C. Hypereutectic alloys contain coarse primary Laves phase dendrites with residual eutectic phase mixtures in between (Fig. l(d)); these do not coagulate during annealing, even at 1450°C. The phase compositions, lattice constants, and phase equilibria have been studied separately in much detail, which allowed the construction of the isothermal sections of the Ni-Al-Ta phase diagram at 1000 and 1250°C.” Alloys with more than 34.2 at% Ta, which are expected to be single-phase Laves phase according to the phase diagram, still contain small amounts below 4 vol.% NiAl which do not dissolve during the applied conventional heat treatments. The melting temperature of the Laves phase TaNiAl is of the order of 1720°C. The remelted alloys (see section 2.1) show a texture with a preferred < 001 > grain orientation parallel to the pulling direction for single-phase NiAI. A similar texture was found for a two-phase NiAl-Ta alloy with 14 vol.% Laves phase, whereas an alloy with 25 vol.% Laves phase showed a preferred < 210 > grain orientation. Correspondingly, a distinct orientation relationship between NiAl and the Laves phase was observed with a (210) habit plane for the TaNiAl precipitate in NiAl, which agrees with earlier observations.21 3.2 Physical properties The densities of the alloys studied vary between 5.9 g cmV3for single-phase NiAl and 1O-4g cmd3 for single-phase TaNiAl (Table 2). A rule of mixtures gives only an order-of-magnitude agreement. Figure 2 shows Young’s modulus for NiAl alloys with various amounts of Laves phase. With the exception of the single-phase alloy NiAl*, all alloys Table 2. Densities of various NiAl alloys studied (see Table 1)

Alloy no. T-O T-10 T-34(L)

Density (g cmP3) 5.95 7.48 10.38

--__ -. _-

- *- - _ --_I-~.__._

---_

NiAMTa

-

--I-<-

_-- --. NiAl-1OTa _JJiAl NiAl-0.2Ta

00 0

200

400

600

800

1000

temperature in “C Fig. 2. Dynamic Young’s modulus as a function of tempera-

ture for NiAl with various alloying additions (all alloys shown are textured with the exception of the differently processed fine-grained NiAl*; the numbers give the amounts of the respective elements in at%).45

in Fig. 2 are remelted alloys with textures as described’ in the preceding section. Indeed the data for NiAl* agree with previous datalo as well as with other data in the literature,22 whereas the values for the other single-phase textured NiAl are much lower. This texture effect is plausible in view of the observed ~001 grain orientation and the large anisotropy parameter A = 2c~/(cir - ~12)= 3.7.23 Taking the data for textured NiAl as reference data, the stiffening effect of the addition of Laves phase to NiAl is clearly visible, whereas the addition of small amounts of Ta in solid solution (alloy NiAl-O.2Ta) has nearly no effect. The comparison of alloys NiAl-5Ta and NiAl-1OTa in Fig. 2 shows that the volume fraction of Laves phase is not important for Young’s modulus in this alloy range with the continuous Laves phase skeleton. The coefficient of thermal expansion for the NiAl alloy T-5 with 5 at% Ta and 14 vol.% Laves phase is 14.1 x lop6 K-l, while the value for textured NiAl without Laves phase is 15-1x 10e6 K-i. The latter value indicates an insensitivity to texture and grain orientation since it agrees with the generally accepted value for polycrystalline NiA1.22

>

3.3 Solid-solution hardening Microalloying of NiAl with Ta without formation of second phases leads to appreciable strengthening by solid-solution hardening in the whole studied temperature range up to 1100°C as can be seen in

NiAl-Ta

567

alloys with Laves phase

Fig. 3 (curve labelled ‘0%‘). The room-temperature data for NiAl-Ta in Fig. 3 are lower limits for the flow stress because of the observed microcracking. It is noted that high-temperature softening occurs at about 750°C for NiAl-Ta. The effect of microalloying on the brittle-to-ductile transition temperature (BDTT) is shown in Fig. 4. The BDTT of NiAl-Ta is distinctly lower at lower deformation rates than that of the binary NiAl alloy. A closer inspection of the fracture surface by scanning electron microscopy revealed intercrystalline fractures, with small Al203 particles (less than 1 pm diameter) on the fracture surfaces for the binary NiAl and transcrystalline fracture, without particles for NiAl-Ta. Obviously, Ta getters oxygen, thus preventing the formation of Al203 particles on grain boundaries. Four-point bending tests with notched specimens indicated a fracture toughness, i.e. a critical stress intensity factor Kr, for crack growth, in the range 46 MPa m % for the ternary alloys. 3.4 Precipitation hardening The 0.2% proof stress of the NiAl-Ta alloys studied increases at all temperatures with the volume fraction of the precipitated Laves phase TaNiAl, as can be seen in Fig. 3. At temperatures above 800°C the stress-strain curves showed maxima at about 1% strain and no fracture was observed for compressive strains up to 5%. Figure 5 shows the proof stress at 1100°C as a function of precipitate volume fraction. Obviously the observed proof stress

IO-7

10-6

IO-5

10-4 10-s lo-’

10-l

100

bending rate in l/s Fig. 4. Brittle-to-ductile transition temperature BDTT (for 1% plastic strain) as a function of deformation rate in fourpoint bending tests for binary NiAl and single-phase NiAl-Ta (alloys T-O, T-0.2 in Table 1) as well as for t+o-phase NiAl alloys (T-0.5-T-25) with various volume fractions of precipitated Laves phase TaNiAl as in Fig. 3.

increase with increasing volume fraction of hard phase does not agree with a simple rule of mixtures, as is supposed to result from the observed structure changes (see section 3.1). The steep increase of proof stress at very small volume fractions of a few per cent is related to the increasing covering of the NiAl grain boundaries with isolated

2500 1

$ I 2000 .E t 8 gQ1 1500

z-96 %

49 % MA 6000

z P P g d

1000

t

0 %--+ ‘..

NiAl

--I 500 temperature

1000

1500

in “C

Fig. 3. 0.2% proof stress (in compression with lop4 s-r deformation rate) as a function of temperature for the single-phase NiAl-Ta alloy T-0.2 without Laves phase (see Table 1) as well as for the two-phase NiAl-Ta alloys (T-0.2-T-34(L)) with various volume fractions of precipitated Laves phase TaNiAl, in comparison with binary NiAl (alloy T-O) and the ODS superalloy MA6000.46

0

20

40 vol.%

60

60

100

Laves phase

Fig. 5. Compressive 0.2% proof stress (with 10m4 s-t deformation rate) for NiAl alloys (T-0.2-T-34(L)) with the Laves phase TaNiAl as a function of volume fraction of Laves phase at 1100°C (full symbols, experimental; dashed curve, rule of mixtures according to a parallel model).

B. Zeumer, G. Sauthofl

568

particles. At volume fractions above 3%, with a less steep proof stress increase, the Laves phase is distributed continuously on the grain boundaries to form a skeleton. The comparison with the ODS superalloy MA 6000 in Fig. 3 shows that higher proof stresses at higher temperatures can be obtained for NiAl alloys by the Laves phase precipitation. 3.5 Creep

scales with normalisation of the applied compressive stress 0 by Young’s modulus E and of the deformation rate E by the Arrhenius factor exp[-Q/RT] for temperature compensation, i.e. the data are plotted as &/exp[-Q/Rq versus a/E. Indeed the data for the thermomechanically treated alloy are all near a common master curve, whereas the data for the as-cast alloy show appreciable scatter. Obviously the microstructure in the as-cast condition with the continuous Laves-phase skeleton changes with changing temperature, which

At higher temperatures the NiAl-Ta alloys deform by creep, which is decelerated by the Laves phase precipitate (see Fig. 6). The observed stress dependence of the secondary creep rate in Fig. 6 obviously follows the familiar power law, .@0: 0.“. This behaviour was observed in the temperature range 800-l 300°C. The observed temperature dependence of the creep rate could be described by an Arrhenius factor, exp[-Q/RT]. Apparent stress exponents n and activation energies Q were deduced from the obtained data and are listed in Table 3. In Fig. 7 the creep data for an NiAl-Ta alloy with 25 vol.% Laves phase are plotted on normalised 10”

,

_.

__

.__

800X 1ooo”c 11oo”c 1200°C

. : .. U. cl0 II 0

0

0

o-*

10”

10"

1

stress I Young’s modulus (4

1

cl

1loo”c

c

100 stress in MN/m2

1000

IO"

(b) Fig. 7. Temperature-compensated

secondary compressive creep rate &exp[Q/Rq (see text) as a function of normalised stress for the thermomechanically treated (a) and as-cast (b) NiAITa alloy T-10 with 25 vol.% Laves phase.

Table 3. Apparent activation energies Q and stress exponents n for creep in various NiAl alloys studied

Alloy T-5 T-10 T-25 T-34(L)

Precipitate TaNiAl TaNiAl TaNiAl TaNiAl

Precipitate volume fraction (%) 14 25 83 >96

lo"

stress I Young’s modulus

Fig. 6. Secondary creep rate at 1100°C as a function of com-

pressive stress (double-logarithmic plot) for NiAl alloys (T-3T-34(L)) with various volume fractions of precipitated Laves phase TaNiAl (data from single-specimen tests with stepwise loading).

10'

Q (kJ mol-‘)

n

463 578 643 288

> 10 (< 1000°C) 4-5 (~10000c) > 10 (< 1000°C) 4-5.5 (LlOOO”C) 4.7-8 (~looooc) 3-5 (2 1OOO°C)

NiAl-Ta

569

alloys with Laves phase

supposedly affects the factor of proportionality in the power law and/or the apparent activation energy Q. The creep resistance - i.e. the stress needed to produce a secondary creep rate of 10e7 s-i - at 1100°C is shown in Fig. 8 as a function of precipitated volume fraction. This creep resistance variation does not follow a simple rule of mixtures, as was already noted for the proof stress in Fig. 5. A subgrain structure develops during creep at 1100°C with inhomogeneously distributed dislocations of type < 100 > {1lo}. A similar behaviour was observed at 1000°C. Cracks form in the Laves phase during creep deformation at temperatures below 1000°C.

cracks are formed in the Laves phase below the BDTT of the Laves phase and are stopped at the phase boundary between Laves phase and NiAl as long as the deformation temperature is above the BDTT of the NiAl phase, which is about 4oo”c.25,26

3.6 BDTT The limited extension of the curves to lower temperatures in Fig. 3 indicates an increasing BDTT with increasing volume fraction of precipitate. The data for alloy T-34(L) with Laves phase matrix and less than 4 vol. % NiAl indicate plastic deformation at 1100°C and above, and no cracks were observed after compressive deformation at 1100°C. From this a BDTT of about two-thirds of the melting temperature is estimated for the Laves phase TaNiAl (which has a melting temperature of about 1720°C); this agrees with findings for other Laves phases.24 Figure 9 illustrates the microstructure of an NiAl-Ta alloy with 25 vol.% Laves phase after compressive deformation at intermediate and high temperatures. It can be seen that cracking occurs in the Laves phase at 600°C and 800°C in the as-cast and thermomechanically treated condition, whereas no cracks are observed at 1100°C. Obviously

0 t-

0

-,-

20

7

-~~

40

4

--T--

60

(b)

60

100

(cl

volume fraction in %

Fig. 8. Compressive creep resistance ondary creep rate shown in Fig. 7) 34(L)) with the Laves phase TaNiAl fraction of Laves phase

(stress for lop7 SK’ secfor NiAl alloys (T-3-Tas a function of volume at 1100°C.

Fig. 9. Optical micrographs of the as-cast NiAl-Ta alloy T-10 with 25 vol.% Laves phase after compression at 600°C with 1O-4 s-’ deformation rate (a), compressive thermomechanical treatment at 1100°C and compression at 800°C (b), and compression at 1100°C (c).

570

B. Zeumer, G. Sauthofl

The BDTT in bending is shown in Fig. 4 as a function of the deformation rate for various NiAlTa alloys with Laves phase. The increase of BDTT with increasing deformation rate indicates that plastic deformation is enabled primarily by thermal activation. A closer inspection by scanning electron microscopy revealed intercrystalline fracture in the Laves phase and transcrystalline fracture in the NiAl phase. The BDTT increase with increasing volume fraction of precipitate is illustrated in Fig. 10. Comparison of Figs 3 and 10 shows that the BDTT in compression is appreciably lower than in bending because of the more advantageous stress state. 3.7 Hot forming The hot-forming behaviour of an NiAl-Ta alloy with 25 vol.% Laves phase was tested by compressing at 1100°C a T-10 specimen fitted with a steel mantle. This test method has been used successfully for characterising the hot-forming behaviour of various materials including NiAl and an intermetallic TiAl alloy. 27~28Fig. 1l(a) shows the forming force as a function of strain, according to which a larger force is needed at the beginning to initiate the forming process. This agrees with the stress maximum for 1% strain, mentioned in section 3.4. The variation of hardness in the NiAl constituent of the two-phase NiAl-Ta alloy with strain in Fig. 11 (b) parallels the variation of forming force in Fig. 11(a). This indicates strengthening processes

in the NiAl constituent, which were not observed in single-phase NiAl. Figure 12 illustrates the microstructure evolution during a forming process. Initially inhomogeneous distributions of dislocations in the NiAl phase with densities below 3 x lo9 cmP2 and practically no free dislocations in the Laves phase were observed in the studied two-phase alloy with coarse phase distribution (Fig. 12(a)). 1% deformation at 1100°C increases the dislocation density in the NiAl phase with subgrain boundary formation (Fig. 12(b)) and produces free dislocations in the Laves phase, preferentially at phase boundaries (Fig. 12(c)). With further straining up to 5-6% the dislocation density increases in both phases, the subgrain formation in the NiAl phase progresses, and fine recrystallised grains of about 300 nm diameter with large-angle grain boundaries are visible in the Laves phase (Fig. 12(d)). This recrystallisation in the Laves phase was observed in both slowly cooled specimens and quenched specimens, which

z _Y

‘5

.-C

$ 10 8 5

04

-

0

I

I--.-~-~~

t

40 20 logarithmic strain in % (4

600 550 500 $j z

450 400 350

600 -I

-I-

0

20

40

60

60

100

vol.%Lavea phase

Fig. 10. Flexural BDTT (from Fig. 4 with lop4 s-l deforma-

tion rate) for NiAl alloys (T-O.ST-25) with the Laves phase TaNiAl as a function of volume fraction of Laves phase.

IO

20

30

I

40

50

60

logarithmic strain in %

-I

0

/

-1

(b) Fig. 11. Forming

force at 1100°C with 5x 10e4 s-l compression rate (a) and Vickers hardness in the NiAl constituent at room temperature after forming (b) as a function of strain for the two-phase NiAl-Ta alloy T-10 with 25 vol.% Laves phase.

571

NiAl-Ta alloys Mdth Laves phase

with annealing time is shown in Fig. 14. The final hardness increase at 1300°C is believed to be an effect of oxidation.

indicates dynamic or metadynamic recrystallisation in the Laves phase. With straining up to 50% the Laves phase disrupts, with NiAl flowing into the opening gaps (Fig. 12(e)). The observed microstructure evolution is sketched schematically in Fig. 13. Static recrystallisation was observed for annealing at 1250°C and 1300°C after deformation at 1100°C. The corresponding variation of hardness

3.8 Oxidation The NiAl-Ta alloy T-5 (Table 1) was selected for studying the oxidation behaviour. Figure 15 shows

(a)

(4

(e) Fig. 12. Bright field TEM micrographs of the two-phase NiAl-Ta alloy T-10 with 25 vol.% Laves phase after various deformations with lOA SK’ compression rate at 1100°C: (a) 0% (bright NiAl, dark Laves phase); (b) 1% (NiAl); (c) 1% (Laves phase); (d) 5.7% (bright NiAl, dark Laves phase with recrystallised grains); (e) 46% (bright NiAl, dark disrupted Laves phase).

572

B. Zeumer, G. Sauthofl

the observed mass change of the specimen with growing oxide scale at various temperatures in an He-02 atmosphere and the derived apparent parabolic rate constants. Obviously oxidation occurs according to the familiar parabolic rate law after some initial nonparabolic scale growth. The initial fast scale growth results from the formation of a

scale of @-Al203 with subsequent transformation to a dense o-Al20s scale. The parabolic rate constant increases with increasing temperature, as is expected. Metallographic observation revealed different scales on NiAl and the Laves phase. Fine-grained a-A1203 formed on the NiAl phase, whereas on the

F

Fig. 13. Microstructure

evolution during high-temperature

compression of a NiAl alloy with Laves phase as in Fig. 12 (schematic, see text).

573

NiAl-Ta alloys with Laves phase

v!

; I

0

400 1200” c .:.... ,:’ _..

L 350

1 ,y-y,300,c 1 0

20

40

60

80

100

time in min

25

Fig. 14. Vickers hardness in the NiAl phase of alloy T-10 at room temperature after 50% compression at 1100°C as a function of annealing time at I250 and 13OO”C,respectively.

Laves phase a bulging layer of a mixed oxide formed, which contained a Ta-Al oxide besides A1203 and which led to a Ta deficiency in the matrix with formation of the Heusler-type phase Ni2TaAI (L21 structure) as revealed by X-ray diffraction. Scale spallation was observed during cooling of specimens. It is concluded that the oxidation behaviour of NiAl-Ta alloys with the Laves phase TaNiAl is similar to that of the analogous NiAl-Nb alloys with the Cl4 Laves phase NbNiAl, which has been studied in detail earlier and has been found comparable to that of binary NiAl for volume fractions of Laves phase that were not too high.29,30 A closer comparison of the data shows that the parabolic rate constant of oxidation of the NiAl-Ta alloy T5 at 1200°C is even lower than that of a comparable NiAl-Nb alloy (with 7 at% Nb to produce 22 vol.% C 14 Laves phase), whereas it is higher at 1100°C and 1000°C.

4. DISCUSSION 4.1 Solid-solution hardening

A discussion of the effects of alloying on the strength of NiAl alloys has to consider contributions of solute atoms, point defects and precipitate particles.3’ The comparison of the data in Fig. 3 for the NiAl alloy with 0.2 at% Ta, which showed no visible precipitate particles, with the data for binary NiAl indicates strong solid-solution hardening. Solution hardening, which was studied extensively for the similar B2 phase CoAl, may be characterised by the hardening rate doo.2/d(Jc)

50 time

1o-'41l 0

25

in

in

100

75

100

h

50 time

75

h

Fig. 15. Mass gain (a) and instantaneous

parabolic rate constant k, (b) as functions of time during oxidation of the NiAITa alloy T-5 (Table 1) in an He-O2 atmosphere with an O2 partial pressure of 133 mbar at various temperatures.47

(c~~.~ = flow stress or 0.2% proof stress, c = amount

of solute).32 The data for the NiAl alloy with 0.2 at% Ta indicate doo.*/d(,/c) = 1 GPa, which is of the order of the effect of Ti in COA~.~~Such effects are usually discussed in terms of differences in atomic size. Indeed, the metallic radii of Ta and Ti on the one side and Ni and Co on the other side differ by only about 1%.33 The contribution of possible effects of constitutional point defects, which result from deviations from stoichiometry, cannot be estimated since the distribution of the Ta atoms on the Ni and Al sublattices, i.e. the effect of the Ta addition on the deviation from stoichiometry, is not yet clear. The only information available is the soltibility lobe in the isothermal section of the Ni-Al-Ta phase diagram, l’ which indicates a slight apparent preference

574

B. Zeumer, G. Sauthof

of Ta for Al sites in NiAl, and a first-principles calculation for the B2 phase FeA1,32 which indicates a preference of Ti for Al sites in FeAl. If Ta occupies Al sites in NiAl, this would mean an excess of atoms in the Al sublattice, in view of the Ni/Al ratio in the alloy T-0.2 (Table 1); this would be compensated for by vacancies in the Ni sublattice with an unknown contribution to strengthening. Another unknown contribution to the strength of the NiAl-Ta alloy under scrutiny may result from the presence of impurities. A recent detailed study of the strengthening of NiAl with microalloying additions of Ti, Zr, or Hf. revealed ageing effects, which were thought to result from the presence of interstitial impurities with formation of very small precipitate particles at or below the limit of microscopic resolution.35 Similar effects may be expected for NiAl-Ta alloys since the solubility of TaC, as a representative possible impurity precipitate in NiAl, is believed to be as low as that of TIC, ZrC, or HfC, given the similar cohesive energies of these carbides.36 The BDTT of the microalloyed NiAl-Ta is lower than that of the binary NiAl in Fig. 4. However, the BDTT of the microalloyed. NiAl-Ta corresponds to that for binary NiAl in a previous studyi and in other work.25y37 Obviously the binary NiAl of this study is embrittled by the presence of impurities, in particular oxygen which resulted in the formation of oxide particles which were observed on the fracture surfaces. 4.2 Precipitation hardening The further flow stress increases of the NiAl-Ta alloys with increasing Ta content above that of the microalloyed alloy with 0*2at% Ta in Fig. 3 should be attributed to hardening caused by the precipitation of the hard and brittle Laves phase TaNiAl. For Ta contents in the range O-2-3at% the Laves phase is precipitated discontinuously as particles, with preferential nucleation at sites of low nucleation energy, these being the grain boundaries in the present case. At higher Ta contents, the NiAl grain boundaries are covered continuously by the Laves phase. In the hypoeutectic composition range up to 14*5at% Ta the alloys solidify with primary NiAl dendrite growth primarily in the direction of solidification and eutectic solidification between the dendrites. It was noted that the microstructures showed some periodicity with Laves phase plates perpendicular to the pulling direction in the directionally solidified alloys, indicating oscillations of the solidification process.

The residual eutectic between the dendrites is dissolved by subsequent heat-treatments with formation of a continuous Laves phase skeleton and isolated NiAl grains in between. Hypereutectic alloys with compositions between 14.5 and about 33.5 at% Ta solidify with primary Laves phase dendrites and eutectic solidification between the dendrites. Again, the resulting microstructure is characterised by a continuously distributed Laves phase with embedded isolated NiAI. The variation of the flow stress at 1100°C with volume fraction of Laves phase precipitate (shown in Fig. 5) does not correspond to a simple rule of mixtures as was mentioned in section 3.4. A rule of mixtures for a composite with alternate lamellae or fibres of NiAl and Laves phase in the stress direction (parallel model) would give a linear increase of flow stress with volume fraction of Laves phase from pure NiAl to pure Laves phase, which indicates the additivity of the flow stresses of the constituent phases, whereas such a composite with the continuous phases perpendicular to the stress (series model) would result in a hyperbola-type variation, which would indicate additivity of strains.38 The flow stress increase of the order of 50 MPa for Laves phase fractions up to 2% in Fig. 5 is larger than expected with any rule of mixtures. The Laves phase is present in these alloys only as particles. The solution hardening effect is expected to be the same for the alloys with different volume fractions of Laves phase since the phases are saturated with constant solute content according to the chemical equilibrium. Then the observed flow stress increase for small amounts of Laves phase is expected to be a result of particle hardening. Particles can be passed by dislocations by the Orowan mechanism or by some climb process at higher temperatures. The latter process gives rise to a threshold stress which is a fraction of the Orowan stress.39 The Orowan stress is of the order of Gb/L with the shear modulus G, the Burgers vector length b and the particle distance L. A value of G of the order of 80MPa is deduced from Fig. 2 for the case shown in Fig. 5. Then the observed flow stress increase of the order of 50MPa would correspond to an Orowan stress if the particle distance were of the order of 500 nm (with b of the order of O-3nm). 22,37The volume fraction of Laves phase in the alloy T-l (Table I), which showed the observed flow stress increase, is of the order of 2%. The particle distance of the order of 500nm corresponds to a particle diameter of the order of 140nm. It is concluded that the steep flow stress

NiAl-Ta alloys with Laves phase

increase in Fig. 5 for alloys with small volume fractions of Laves phase indicates the presence of small precipitate particles, and indeed such small Laves phase particles were observed in NiAl-Ta alloys studied by electron microscopy. The further flow stress increase in Fig. 5 for NiAl-Ta alloys with larger volume fractions of Laves phase above 2 vol.% is practically linear up to about 70 vol.% Laves phase. This indicates additivity of the flow stresses according to a rule of mixtures for a parallel model, i.e. in this intermediate volume fraction range around the eutectic composition with 49 vol.% Laves phase the alloys behave as a composite with continuous phases along the loading direction. It is noted that the increase of the BDTT with increasing volume fraction of Laves phase parallels the corresponding variation of the flow stress, i.e. there is a steep increase for small volume fractions with increasing coverage of the grain boundaries and a less steep increase for intermediate volume fractions when a continuous Laves phase skeleton is present. Only above 70 vol.% Laves phase is there a transition to a hyperbola-like steep increase of flow stress. This is illustrated in Fig. 5 by the dashed interpolating curve which corresponds to a rule of mixtures for a series model as derived for an elastic constant.38 Obviously Laves phase alloys with only small fractions of the softer NiAl behave as composites with discontinuous phases in the loading direction. Such behaviour was observed and analysed in detail for a hard epoxy resin toughened by soft discrete rubber spheres.40 4.3 Creep The obtained creep resistances agree with the findings on the eutectic lamellar NiAl-Ta alloys, prepared by directional solidification.20 This indicates an insensitivity in the creep of such alloys at high temperatures to variations of the coarse Laves phase distribution. The observed creep behaviour corresponds to the familiar power law as has been observed and analysed for various other intermetallic phases and alloys2 In particular, the range 4-55 of the apparent stress exponents in Table 3 for the NiAlrich alloys with 14 or 25 vol.% Laves phase above lOOO”C, together with the observed subgrain formation, indicates dislocation creep with climb as the rate-controlling mechanism. This agrees with the behaviour of single-phase NiAl-base alloys.41 Likewise the stress range 3-5 in Table 3 for the practically single-phase Laves phase TaNiAl agrees

575

with the previous findings for the Laves phase NbNiAl.‘O Higher stress exponents, as shown in Table 3 for lower temperatures or higher volume fractions of Laves phase, usually indicate the priesence of a threshold stress as an effect of small pgrticles. This possibility could be excluded in the preisent case by detailed data analysis and microscopic observations. However, the temperatures for creep with exceptionally high stress exponents wefe below the respective BDTTs for the Laves phase, and indeed microcracks were observed in the Lave! phase after creep. Obviously, microcracking ocburs during creep, which is enhanced by higher stresses, and accelerates the creep. Thus, a str6nger stress dependence of apparent creep is believed to result. Such effects have been observed and analysed for brittle ceramics and graphite.42,43 The apparent thermal activation energy for creep in intermetallics is expected to appr4ximate that for diffusion as is the case for conventional metallic alloys.2 However, the apparent activation energies for creep of the NiAl-rich alloys in ‘Table 3 are distinctly higher than those for diffusion - about 310 kJ/mol for the diffusion of Ni in stioichiometric NiAl and less in non-stoichiometric NiA1.44 Such differences in activation energies have1 been reported frequently for intermetallic alloys and are supposed to be related to temperature-dependent structure changes. According to the data in Table 3, the apparent activation energy for cr$ep seems to increase with increasing fraction of Laves phase. However, the apparent activation energy for the practically single-phase TaNiAl is disdinctly lower, which precludes any application of a simple rule of mixtures. It is noted that the observed activation energy for TaNiAl is also distinctly lower than that for the previously studied NbNiAl.‘O 4.4 Hot forming The compressive BDTT of the Laves phase was found to be of the order of two-thirds of the melting temperature, i.e. about 1020°C for TaNiAl. The NiAl alloys with various amOUhtS of Laves phase could indeed be compressed plastically up to 50% at 1100°C without fracturing. The sequence of events during large-strain compression at 1100°C is visible in Fig. 12. The flow’ stress of the Laves phase is about 1700 MPa at this temperature whereas it is only 40 MPa for NiAl. Id spite of this enormous difference in flow stress thi: hard Laves phase, which is nearly free of dislocatCons initially, is affected by the compression, i.e. it deforms

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plastically with increase of dislocation density, as well as the soft NiAl phase. This initiates dynamic recrystallisation even in the Laves phase, which was indicated by the maximum in forming force in Fig. 1l(a). However, a new fine grain structure in the Laves phase was observed directly after compression and quenching, which could only be produced by dynamic recrystallisation during deformation or metadynamic recrystallisation after deformation during quenching. These observations agree with findings on dynamic recrystallisation in other Laves phases at about 0.7 of the melting temperature.24 Besides deformation of the Laves phase with recrystallisation, the compression produces disruption of the Laves phase with subsequent intrusion of the soft NiAl phase into the opening gaps between the Laves phase pieces, as can be seen in Fig. 12(e). This mechanism transforms the initial alloy structure with a skeleton of continuously distributed Laves phase into a duplex-like structure with a mixture of NiAl grains and Laves phase grains. This grain structure is comparatively stable since no spheroidisation of the Laves phase grains was observed during subsequent annealing. After compression at 1100°C with dynamic recrystallisation, the dislocation density in the Laves phase is still higher than after compression with subsequent creep. Anneals at higher temperatures after compression at 1100°C produce hardness variations as shown in Fig. 14, which is characteristic for static recrystallisation. Obviously, the residual dislocation density after compression was sufficient to initiate some further static recovery and/or recrystallisation processes.

5. CONCLUSIONS NiAl-base alloys with various amounts of Ta to produce the Laves phase TaNiAl with Cl4 structure as the strengthening second phase were studied with respect to their deformation behaviour in compression and bending at ambient and high temperatures - including elastic deformation and creep - as a function of alloy composition and microstructure. The following conclusions are drawn from the results: (1) Up to 0.2 at% Ta dissolves in NiAl and produces solid-solution strengthening. (2) Ta in solid solution getters impurities and avoids the increase of the brittle-to-ductile transition temperature BDTT due to impurities.

(3) Ta contents in the range 0.2-3 at% produce precipitated particles of the Laves phase TaNiAl with hexagonal Cl4 structure in NiAl, primarily on grain boundaries and secondarily in grains. (4) Ta contents above 3at% result in NiAl alloys with continuous distribution of the Laves phase TaNiAl surrounding the isolated NiAl grains. (5) The flow stress, creep resistance, and BDTT of the two-phase NiAl-Ta alloys increase with increasing volume fraction of Laves phase. (6) The BDTT of the Laves phase TaNiAl is about two-thirds of the melting temperature in agreement with data for other Laves phases. (7) The contribution of the Laves phase particles to the flow stress is higher than according to a rule of mixtures. (8) The Young’s modulus of the NiAl-Ta alloys depends sensitively on the alloy texture and the presence of the continuously distributed Laves phase, in contrast to the behaviour of the coefficient of thermal expansion. (9) Deformation at 1100°C leads to recrystallisation of the Laves phase in the two-phase NiAl-Ta alloys. (10) The two-phase NiAl-Ta alloys creep according to a power law with dislocation climb as the rate-controlling mechanism at 1100 and 1000°C and stresses above lOMPa, whereas creep at lower temperatures is accompanied by microcracking in the Laves phase, the cracks being stopped at NiAl-Laves phase boundaries. (11) The two-phase NiAl-Ta alloys show a good compares oxidation resistance which favourably with that of similar NiAl-Nb alloys as well as with binary NiAl for Laves phase contents that are not too high. The observed properties indicate alloys with the strengthening Laves are a promising base for materials aiming at structural applications at

that NiAl-Ta phase TaNiAl developments 1200°C.

ACKNOWLEDGEMENTS The authors are indebted to Professor Dr W. Pitsch for his continuous support of this work. The financial support of the Bundesministerium fiir Forschung und Technologie is gratefully acknowl-

AU-Ta

altoys with Laves phase

edged. It is a pleasure for the authors to thank Mr R. Staegemann, Mrs E. Bartsch, and Mr G. Bialkowski for alloy preparation, electron microscopy, and mechanical testing, respectively.

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