Microstructure and mechanical properties of Laves-phase alloys based on Cr2Nb

Materials Science and Engineering, A 132 (1991 ) 61-66

61

Microstructure and mechanical properties of Laves-phase alloys based on Cr2Nb M. Takeyama* and C. T. Liu Oak Ridge National Laboratory, p.o. Box 2008, Oak Ridge, TN 37831-6115 (U.S.A.)

(ReceivedMay 22. 1990)

Abstract Mechanical properties of Laves-phase alloys based on Cr2Nb at temperatures up to 1000°C were examined and correlated with rnicrostructures and phase relationships. Single-phase Cr2Nb alloys are very hard and brittle at ambient temperatures, indicating the difficulty in generation and glide of dislocations due to the complicated crystal structure (C-15). Examination of the Cr-Cr2Nb two-phase region revealed the following: (a) the hardness decreases with increasing amounts of the soft chromiumrich phase; (b) the eutectic composition has a niobium concentration of 17 at.%, instead of 12 at.% as reported in the currently existing phase diagram; (c) heat treatments produce uniform dispersion of fine Laves-phase precipitates in primary chromium-rich patches for the hypoeutectic alloys, and these particles are very stable even at temperatures above 1000 °C; (d) the soft particles are very effective in preventing crack propagation originating in the brittle Laves-phase matrix, which results in a high yield strength with moderate ductility up to 1000 °C. These results demonstrate that the introduction of a soft chromium phase has promising effects in improving the mechanical properties of brittle Cr2Nb Lavesphase alloys.

1. Introduction The Cr2Nb Laves phase has a C-15 (cF24) cubic structure with a stacking sequence of XYZ type, where X, Y and Z represent closed-packed layers, similar to an f.c.c, structure; however, each layer is composed of four interpenetrating atomic layers. Its unit cell contains 24 atoms with a lattice parameter of 6.98 A [1-4]. The crystal structure is so complicated that dislocation movement appears to be difficult at ambient temperatures. Concerning the Von Mises criteria, Cr2Nb does not have a sufficient number of independent slip systems for uniform, extensive plastic deformation. This Laves-phase alloy, however, is attractive for high-temperature structural applications because of its high melting temperature (1770°C) [5, 6], relative low density (approximately 7.7 g c m - 3) [7], and potential resistance to oxidation [5]. In addition, this intermetallic alloy has a wide range of composition homogeneity,

indicating the possibility of improving its mechanical and metallurgical properties by alloying additions. Recently, some attention has been given to the deformation behavior of Laves-phase alloys [8-1 1]. For example, Sauthoff has revealed excellent creep resistance of Cr2Ti Laves phase at temperatures up to 1100 °C [11]. The principal objective of this study is to develop Laves-phase intermetallic alloys based on Cr2Nb for structural use at high temperatures. Since we found that the single-phase Laves alloys are extremely brittle at ambient temperatures, our efforts were then placed on Cr-Cr2Nb two-phase compositions with the soft chromium-rich phase dispersed in the hard Laves phase (or vice versa). The high-temperature properties were characterized by hardness measurements and compressive tests, and correlated with the microstructures and phase relationships. 2. Experimental details

*Present address: National Research Institute for Metals, 2-3-12 Nakameguro, Meguro-ku,Tokyo 153,Japan. 0921-5093/91/$3.50

The alloys used in this study have nominal compositions of chromium with 6-34 at.% Nb. © Elsevier Sequoia/Printed in The Netherlands

62

(Hereafter, all compositions in this paper are given in at.%.) These alloys were prepared by arcmelting and drop-casting into a chilled copper mold in an argon atmosphere. The ingot with a size of 12 m m × 25 minx 130 mm showed some surface cracks because of severe thermal shock during drop-casting. The oxygen level in the material varied from 480 to 970 ppm by weight. The alloys were encapsulated in an evacuated quartz tube and heat treated at 1100 °C for 100 h or 900 °C for 5 days. Compression specimens having a cross-section of 4.7 mm × 4.7 mm and a height of 7.2 mm were cut by electrodischarged machining from the annealed ingots. The specimens were tested under compression on an Instron testing machine at room temperature, 500 and 1000°C at a nominal strain rate of 3.3 x 10 -3 s -]. The elevated temperature tests were carried out in vacuum (less than 7 x 10 -4 Pa). The microhardness was measured on as-cast alloys with a 50 g load. Microstructures of these alloys were examined by optical and scanning electron microscopies. Metallographic specimens were electropolished in a solution of 10 HC1 and 90 C2HsOH by parts. An electron probe microscope equipped for wavelength-dispersive spectrometry (WDS) was also used for analyzing the chemistry of phases in the alloys. 3. Results and discussion

3.1. Hardness Change in room temperature hardness with niobium concentration of several C r - N b alloys in the as-cast condition is shown in Fig. 1. The single-phase Laves alloys (greater than 30% Nb) are very hard, and their hardness values are larger lO

(D

than 8 GPa. Also, alloy stoichiometry does not strongly influence the microhardness, as evident from the alloys of Cr-32.3%Nb and Cr-34.3%Nb. Microcracks were observed at the tips of the indents. The high hardness and crack tendency simply indicate that the single-phase alloys were very brittle at room temperature, possibly because of the difficulty in nucleation and glide of dislocations. The hardness decreases with decreasing niobium concentration, and its value drops to 3.3 GPa for a niobium concentration of 6%. This reduction is apparently due to the introduction of a soft chromium-rich phase (1.6 GPa) [5], as mentioned below. From this result, we focused our studies on the two-phase (chromiumrich plus Laves) alloys. 3.2. Microstructures A backscattered electron image of as-cast Cr-20%Nb alloy is shown in Fig. 2. Layered lamellar structures denoted as (A) are observed, together with blocky plate-like phases denoted as (B). The lamellar structure indicates the eutectic reaction that occurs between the chromium-rich phase and the Laves phase, which is consistent with the currently existing chromium-niobium binary phase diagram [6]. Point analyses using WDS performed at both the eutectic regions and blocky phases reveal average niobium concentrations of 16.9% and 30.2% respectively. The blocky phase is, no doubt, the Laves phase. The same analyses were also performed for the as-cast alloy with 12% Nb as well as the alloys annealed at l l 0 0 ° C for 100h. The results allow us to redraw partially the phase diagram of the chro-

t As cast alloys

6

,,/-

t9 4

"t" ~

[]

2

/

I1: presentresults L Io: da= fromGold~.~,n=o~t S ] /

_ _ J

,~ /

0

10

/

20

30

40

Nb concentration (at. %) Fig. ]. Change in room-temperature hardness with niobium

concentration in as-cast Cr-Nb alloys.

Fig. 2. Backscattered electron image of as-cast Cr-20%Nb alloy. The marks (A) and (B) refer to the Cr-Cr2Nb eutectic region and primary Cr2Nb Laves phase respectively.

63

mium-rich side, which is superimposed on the currently existing phase diagram in Fig. 3. The results reveal two major differences: (a) the eutectic composition contains about 17% Nb, instead of 12% Nb; (b) the phase boundary between the Cr2Nb single phase and the Cr-Cr2Nb two-phase region is located at about 30% Nb regardless of temperature. Optical microstructures of as-cast and annealed (at 900°C for 5 days) samples of Cr-12%Nb and Cr-6%Nb alloys are shown in Fig. 4. In the as-cast condition, the Cr-12%Nb alloy exhibits patches of the chromium-rich phase surrounded by the eutectic phase matrix (Fig. 4(a)), indicating that this is a hypoeutectic structure. Some dark dots can be seen within the chromium-rich patches, and they are probably the Cr2Nb Laves-phase particles formed during cooling, as mentioned below. In the as-cast Cr-6%Nb, chromium-rich patches cover most areas, and a small amount of the Laves phase is seen as a network along the chromium-rich

2000 - --e--

1863oc

Go~dschmidt [5,6] Present data

1800 ~ k ~ ,

ov

1600

== 0~

1770°C

I-" ~/6 I(Cl':/

12

;7 "

3o~ I ',/ C:r2Nb I

,H

1400

E I.-

r

'

i

i

i

1200 i 1.7

30i

/

1000 0

10

20

30

40

50

Atomic percent Niobium

Fig. 3. Modified chromium-rich side of the Cr-Nb binary phase diagram from the present study, indicating two major differences: (a) the eutectic composition was 17 at.% Nb, instead of 12 at.% Nb; (b) the phase boundary between C r + C r , N b and CrNb was located at about 30 at.% Nb regardless of temperatures above 1000 °C.

Fig. 4, Optical microstructures of as-cast ((a), (b)) and annealed ((c), (d)) samples of Cr-I 2%Nb ((a), (c)) and Cr-6%Nb ((b), (d)) alloys.

64

patches (Fig. 4(b)). Interestingly, the primary chromium-rich regions turn dark after annealing, as shown in Figs. 4(c) and 4(d) for Cr-12%Nb and Cr-6%Nb alloys respectively. Now, let us examine the microstructure of chromium-rich patches in detail. A scanning electron micrograph showing a primary chromium-rich region of Cr-12%Nb alloy annealed at l l 0 0 ° C for 100 h is given in Fig. 5. Numerous fine particles (less than 1 ~m) are clearly visible in the region. (Some black dots seen in the region are holes where fine Laves-phase particles possibly dropped out during sample preparation.) These precipitates were identified as a Cr2Nb Laves phase. The precipitation of Laves-phase particles after annealing is induced by a decrease in solubility of niobium in the chromium-rich phase from about 5% at the eutectic temperature (1620°C) to 1.7% at l l 0 0 ° C (refer to Fig. 3). It should be noted that these Laves-phase particles are extremely fine and stable even after extensive annealing at l l00°C. Another feature to be noticed in this microstructure is that the primary chromium-rich region is generally surrounded by a blocky Laves phase, indicating that the matrix phase in the eutectic region is the Laves phase rather than the chromium-rich phase. Based on microstructural examination, an annealing treatment at 900 °C for 5 days was chosen to study the mechanical properties of these alloys. This is simply because a lower annealing treatment temperature produces a finer precipitation of Laves-phase particles in primary chromium-rich regions.

3.3. Mechanical properties The 0.2% yield strength and compressive strain of Cr-6%Nb and Cr-12%Nb alloys annealed at 900 °C are plotted as a function of temperature in Fig. 6, where the data for a cast nickel-base superalloy IN 713C are also plotted for comparison [12]. At room temperature, the yield strengths of Cr-6%Nb and Cr-12%Nb alloys are higher than 800 and 1200 MPa respectively, and decrease gently with increasing temperature. Even at 1000°C, the alloy Cr-12%Nb exhibits a yield strength of about 800 MPa (Fig. 6(a)), which is substantially stronger than conventional nickel-base superalloys; for example, the yield strength of IN713C at 1000°C is about 300 MPa. The compressive strains, however, of Cr-6%Nb and Cr-12%Nb alloys are about 11% and 5% at room temperature respectively, and remain almost unchanged up to about 500 °C. At 1000 °C, both alloys can be deformed extensively with compressive ductilities larger than 25% (Fig. 6(b)), at which point the tests were interrupted at a strain of 26%-30% for these alloys. These results reveal that the twophase intermetallic alloys based on Cr2Nb have high strength with decent ductility under compression at all test temperatures. Further studies

1600

I Annealed at 900"C

(a)

0. 1200

~

(n

80O

12

Nb

0

~"

400 Alloy 713C

~

40

(b)

¢-

,~

3o

.~ u~ 0) c~

20

E

p~ / //..

_

O

o

i

t

it

10

~.'>~-~- __

/ Alloy 713C

121Nb 0

o

400

800

1200

T e s t temt~erature (°C)

Fig. 5. Scanning electron image of Cr-12%Nb alloy annealed at 1100 °C for 100 h, showing fine Cr:Nb Lavesphase precipitates in primary chromium-rich patches surrounded by the Laves-phase matrix.

Fig. 6. (a) The 0.2% offset yield strength and (b) compressive strain of annealed Cr-6%Nb and Cr- 12%Nb alloys as a function of temperature, shown with typical tensile properties of cast alloy IN 713C for comparison.

65

Fig. 7. Scanning electron micrograph showing the fracture surface of Cr 12%Nb alloy tested at room temperature.

are required to characterize their tensile properties at ambient and elevated temperatures. A fracture surface of the C r - 1 2 % N b alloy tested at room temperature is shown in Fig. 7. The dimple-like patches and relatively smooth areas correspond to the primary chromium regions with Cr2Nb particles and the Laves-phase matrix respectively. Some cleavage planes and smooth grain-boundary facets are seen on the fracture surface corresponding to the Lavesphase matrix. Cracks are also observed along interfaces between the dimpled regions and the adjacent Laves-phase matrix. Optical microstructures of a cross-section of the C r - 1 2 % N b specimens tested at room temperature and 1000°C are shown in Fig. 8. On room-temperature deformation (Fig. 8(a)), cracks occur almost exclusively in the eutectic Lavesphase matrix or at interfaces between the matrix and the chromium-rich patches with Cr2Nb precipitates. It should noted that no macrocracks were observed in the primary chromium-rich patches. In the specimen deformed to 26% at 1000 °C (Fig. 8(b)), a number of relatively small flaws bridge across the eutectic Laves-phase matrix, while no such flaws are observed within the patches. These indicate that the chromiumrich region with fine Laves-phase precipitates is apparently effective in preventing crack propagation. Similar features are also observed in the Cr-6%Nb alloy. The present study clearly shows that the introduction of the soft chromium-rich phase is responsible for the improved mechanical properties of the Laves-phase alloys based on Cr,Nb. The single-phase alloys exhibit microcracks at

l

I

I

Fig. 8. Optical micrographs of Cr- 12%Nb alloy fractured at (a) room temperaturc and (b) 1000 °C.

tips of the hardness indents, while the two-phase alloys show no cracks on hardness measurement and have some ductility (5%-11%) in compression tests at room temperature. Precipitation of ductile second-phase particles has also been reported to be effective in improving the tensile ductility of T i3AI (hexagonal, DO,,~) alloyed with niobium additions [13-17]. The increased plasticity in two-phase alloys could be associated with introduced interfaces. Noebe et al. [18] demonstrated that a substantial increase in ductility was achieved in some brittle B2-type intermetallic alloys by introducing interfaces via surface films or second phases, which generate a sufficient number of mobile dislocations to enhance plastic deformation. In order to understand the enhanced plasticity, further study on the dislocation microstructure at Cr-Cr~Nb interfaces is necessary. Although the chromium-rich patches with fine, stable Laves-phase precipitates play an important

66

role in improving the mechanical properties in the two-phase alloys, the major cause for brittleness still comes from the Laves-phase matrix which is prone to crack initiation and propagation even at elevated temperatures. Further studies are thus required to improve the toughness of the Cr2Nb Laves phase. Since the CrzNb phase has a wide range of composition homogeneity, alloying additions could prove to be useful in enhancing the matrix properties, leading to further improvement in the mechanical properties of these alloys with a two-phase structure.

4. Summarizing remarks Single-phase Cr2Nb is very hard and brittle at ambient temperatures. The high hardness indicates difficulty in nucleation and glide of dislocations. One way to improve substantially the ductility and toughness is to introduce ductile, soft particles. Our study reveals that the chromium-rich phase with fine, stable Cr2Nb precipitates is very effective in stopping crack propagation originating at the Laves phase. A moderate ductility of 5%-11% in compression has been achieved in the two-phase alloys tested at room temperature. These two-phase alloys are substantially stronger than conventional nickelbase superalloys at high temperatures.

Acknowledgments The authors wish to thank J. A. Horton and C. G. McKamey for helpful discussions; T. J. Henson, E. H. Lee, and H. D. Pierce for technical assistance; and J. D. Vought for alloy preparations. We also thank C. L. Dowker and W. E. Gilliam for manuscript preparation. Research funded by the Director's Discretionary Funds of Oak Ridge National Laboratory, operated by Martin Marietta Energy Systems, Inc., under contract DE-AC05-84OR21400 with the U.S. Department of Energy.

References 1 C. S. Barrett and T. B. Massalski, Structure of Metals, McGraw-Hill, New York, 1966, 3rd edn. 2 E Laves, in Theory of Alloy Phases, American Society for Metals, Metals Park, OH, 1956, p. 124.

3 R. P. Messmer, R. C. Tatar and C. L. Briant, in J. L. Walter, M. R. Jackson and C. T. Sims (eds.), Alloying, American Society for Metals, Metals Park, OH, 1988, p.29. 4 E Villars and L. D. Calvert (eds.), Pearson's Handbook of Crystallographic Data for lntermetallic Phases, American Society for Metals, Metals Park, OH, 1985. 5 H. J. Goldschmidt and J. A. Brand, J. Less-Common Met., 3 (1961) 44. 6 T.B. Massalski, J, L. Murray, L. H. Bennett and H. Baker (eds.), Binary Alloy Phase Diagram, American Society for Metals, Metals Park, OH, 1986. 7 A.I. Taub and R. L. Fleischer, Science, 243 (1989) 616. 8 J. D. Livingston, E. L. Hall and E. E Koch, in C. T. Liu, A. I. Taub, N. S. Stoloff and C. C. Koch (eds,), High Temperature Ordered lntermetallic Alloys, II1, Materials Research Society, Pittsburgh, PA, 1989, Vol. 133. 9 E. L. Hall and J. D. Livingston, in W. G. Baily (ed.), Proc. 41st Ann. Meet. of EMSA, San Francisco Press, San Francisco, CA, 1989. 10 D. L. Anton, paper presented at The Metallurgical

Society~American Society for Metals Symp. on HighTemperature Aluminides and lntermetallics, October 2-5, 1989, Indianapolis, IN. 11 G. Sauthoff, in S. H. Wang, C. T. Liu, D. P. Pope and J. O. Stiegler (eds.), Proc. The Metallurgical Society[American

Society for Metals Syrup. on High-Temperature Aluminides and lntermetallics, The Metallurgical Society of AIME, Warrendale, PA, 1990, pp. 329-352. 12 Engineering Properties of Alloy 713C, International Nickel, New York, 1968. 13 M.J. Blackburn, D. L. Ruckle and C. E. Bevau, Research to conduct an exploratory experimental and analytical investigation of alloys, Tech. Rep. AFML-TR-78-18, 1978 (Air Force Materials Laboratory, United Technologies, Pratt & Whitney, Hartford, CT ). 14 M.J. Blackburn and M. E Smith, Research to conduct an exploratory experimental and analytical investigation of alloys, Tech. Rep. AFML-TR-81-4046, 1981 (Air Force Wright Aeronautical Laboratories, United Technologies, Pratt & Whitney, Hartford, CT ). 15 M. J. Blackburn and M. P. Smith, Development of improved toughness alloys based on titanium aluminides, Interim Tech. Rep. FR-19139, 1988 (United Technologies, Pratt & Whitney, Hartford, CT). 16 R. Strychor, J. C. Williams and W. A. Sofia, Phase transformations and modulated microstructures in T i - A I - N b alloys, Metall. Trans. A, 19 (2) (1988) 225-234. 17 W. A. Baeslack III, M. J. Cieslak and T. J. Headley, Structure, properties and fracture of pulsed Nd: YAG laser welded Ti-14.8wt%AI-21.3wt%Nb titanium aluminide, Scr. Metall., 22 (1988) 1155-1160. 18 R. D. Noebe, J. T. Kim, M. Larsen and R. Gibala, in S. H. Whang, C. T. Liu, D. E Pope and J. O. Stiegler (eds.),

Proc. The Metallurgical Society/American Society for Metals Syrup. on High-Temperature Aluminides and lntermetallics, The Metallurgical Society of AIME, Warrendale, PA, 1990, pp. 271-300.