Ductile L12Rh3Ti intermetallics for ultra-high temperature applications

Materials Science and Engineering A329– 331 (2002) 486– 491 www.elsevier.com/locate/msea

Ductile L12Rh3Ti intermetallics for ultra-high temperature applications S.M.L. Sastry a,*, R. Mahapatra b, A.W. Davis b a

Department of Mechanical Engineering, Washington Uni6ersity, Campus Box 1185, One Brookings Dri6e, St. Louis, MO 63130 -4899, USA b Materials Laboratory, Na6al Air Systems Command, Patuxent Ri6er, MD 20670, USA

Abstract The microstructures and deformation behavior at 25 – 1000 °C of Rh– xTi (x =15, 20, and 25 at.%) alloys were studied. The alloys were prepared by non-consumable electrode arc melting, The arc-melted buttons were isothermally forged at 1100 °C, and the mechanical properties at 25–1000 °C were evaluated by microhardness measurements, compression testing, and 3-point bend testing. The deformed specimens were examined by scanning electron microscopy for the determination of the extent and nature of slip and by transmission electron microscopy for the identification of dislocation activity. Unlike many other high temperature intermetallics, the Rh–Ti alloys do exhibit room temperature ductility. The high temperature strength and creep resistances are superior in Rh–Ti alloys to some of the other high temperature intermetallics. © 2002 Elsevier Science B.V. All rights reserved. Keywords: Mechanical properties; Intermetallics; Ductility; Rh– Ti alloys; Mar M200

1. Introduction In recent years there has been an increasing amount of interest in platinum group metals as base elements for ultra-high temperature materials. These elements possess extremely high melting temperatures while exhibiting superior oxidation resistance compared to the refractory metals [1]. Of the platinum group metals (Ir, Pt, Rh, and Pd), rhodium displays the most promising combination of properties for aerospace propulsion applications. When alloyed with Ti, ordered Rh3Ti based on L12 structure is formed over a range of composition [2]. L12 ordered phases such as Cu3Au consist of a symmetric fcc superlattice, exhibit several active slip systems at room temperature, and usually are ductile at mom temperature. Many L12 ordered phases exhibit an increase in flow stress with increasing temperature making them attractive for high temperature applications [3]. Furthermore, in Rh – Ti with 12–

* Corresponding author. Fax: + 1-314-935-4014. E-mail address: [email protected] (S.M.L. Sastry).

22 at.% Ti, a two-phase microstructure of Rh –Ti solid solution and Rh3Ti precipitates can be produced by suitable heat treatment. This microstructure is similar to the k/k% type microstructure of nickel-base superalloys which have good room temperature ductility and high temperature creep resistance. The principal benefits of Rh –Ti alloys are shown Table 1 which compares some physical properties of Rh and Rh3Ti to the nickel-base superalloy Mar M200. Rhodium melts at a temperature of 1963 °C, which is about 500 °C higher than the melting point of Mar M200. In addition, Rhodium has a lower thermal expansion coefficient and a higher thermal conductivity, which are beneficial properties for high temperature applications. Rh –Ti alloys have a slightly higher density than nickel-base superalloys; however, the density compensated properties of Rh –Ti alloys are expected to exceed those of nickel-base superalloys. In this paper we present preliminary results of microstructures and mechanical properties of Rh –15Ti, Rh –20Ti, and Rh –25Ti which indicate that these alloys have the potential to exceed the temperature capability of nickel-base superalloys.

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Table 1 Physical properties of Mar M200, Rh and Rh–25Ti alloys Alloy

Density (kg m−3)

Melting point (°C)

Thermal expansion (10−3 °C−1)

MarM200 Rh Rh–25 Ti

900 1244 1046

1550 1963 1800

1580 1080

Table 2 Chemical compositions of arc-melted Rh–Ti alloys Alloy

At.% Ti

Error (%)

Rh–15Ti Rh–20Ti Rh–25Ti

14.80 19.14 26.74

1.77 1.05 1.55

2. Experimental procedure Rh – Ti alloys of 15, 20, and 25 at.% Ti were prepared in the form of rods by non-consumable electrode arc melting of high purity Ti (99.998%) and Rh (99.9%) in

Thermal conductivity (10−2 W (mK)−1)

840 1470

a pure argon atmosphere. The as-cast compositions were confirmed by energy dispersive spectroscopy (EDS) performed on a scanning electron microscope (SEM). Arc-melted samples were heat treated at 1350 °C for 4 h in an argon atmosphere. Samples for scanning electron microscopy (SEM) and transmission electron microscopy (TEM) were cut perpendicular to the longitudinal axis by electro-discharge machining (EDM). The arc-melted and heat treated alloys were examined by X-ray diffraction, SEM and TEM. Specimens for TEM were prepared by ion milling. For the determination of mechanical properties, arcmelted samples with compositions Rh– 15Ti and Rh – 25Ti were isothermally forged at 1100 °C. Cube-shaped

Fig. 1. X-ray diffraction patterns for (a) arc-melted Rh – 15Ti and (b) Rh– 15Ti heated treated at 1350 °C for 4 h.

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Fig. 2. X-ray diffraction patterns for (a) arc-melted Rh – 20Ti and (b) Rh– 20Ti heat treated at 1350 °C for 4 h.

coupons with side lengths of 3 mm were cut with a thin diamond saw and deformed in compression at 25, 500, 800, 900, and 1000 °C. Three point bend tests were performed at room temperature on rectangular samples with dimensions 10×3 ×2 mm.

Fig. 3. X-ray diffraction pattern of arc melted Rh –25Ti.

3. Results and discussion The chemical compositions of the alloys determined by X-ray energy dispersive analysis in the SEM shown in Table 2 are close to the desired compositions. X-ray diffraction patterns obtained from surfaces perpendicular to the longitudinal axes of the arc-melted and heat treated rods are in Figs. 1–3. All of the arc-melted alloys show clear fcc fundamental peaks as expected. The Rh–20Ti and Rh – 25Ti (Fig. 3a and b and Fig. 4) alloys show additional peaks corresponding to the L12 superlattice. Upon heat treatment at 1350 °C, the superlattice peaks are more well defined in Rh –15Ti and Rh –20Ti alloys. The intensity of the (111), (220), and (311) peaks decreased indicating the possible development of texture. Transmission electron micrographs of the microstructures of arc-melted Rh–15Ti and Rh –20Ti alloys are shown in Fig. 4a and b. The arc-melted Rh–15Ti alloy contains evenly dispersed, equiaxed precipitates of :10 nm size as shown in the dark field micrograph in Fig. 4a. The selected area diffraction patterns of the Rh–

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Fig. 4. TEM micrographs (a, b) and selected area diffraction patterns (c, d) of arc melted Rh – 15Ti (a, c) and Rh – 20Ti (b, d). (a is DF micrograph, b is BF micrograph).

Fig. 5. Bright field transmission electron micrograph of Rh – 25Ti single phase alloy.

15Ti alloy (shown in Fig. 4c) contained rings corresponding to fundamental as well as superlattice reflections, which indicates that the precipitates solidified incoherently within the matrix with no significant lattice matching. Although this incoherent matching is not expected in these alloys, this behavior could be the result of rapid cooling which occurs during solidification after arc melting.

The arc melted Rh–20Ti alloy also consists of Rh3Ti precipitates in an Rh–Ti matrix as shown in the bright field micrograph in Fig. 4b. As would be expected from the lever rule, the volume fraction of the Rh3Ti precipitates is greater in the Rh–20Ti alloy than in Rh–15Ti. However, the precipitate sizes are comparable. These precipitates seem to solidify coherently with the matrix as indicated by the sharp spots in the selected area diffraction (SAD) pattern (Fig. 4d). Transmission electron microscopy of the Rh–25Ti alloy revealed a single phase structure seen in Fig. 5 and superlattice reflections in the diffraction pattern consistent with the L12 Rh3Ti phase. Annealing the Rh–15Ti and Rh–20Ti two-phase alloys at 1350 °C for 4 h changed the microstructure of the Rh –15Ti alloy significantly. Fig. 6a is a bright field micrograph of the heat treated Rh–15Ti alloy. The size, shape, and morphology of the precipitates have changed, considerably. The precipitates appear as discs about 50 nm in diameter and 20 nm thick. These discs are arranged orthogonally to each other in a

Fig. 6. TEM micrograph of Rh –15Ti alloy heat treated at 1350 °C for 4 h. (a) & (b) are bight field micrographs, (c) is dark field micrograph.

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Fig. 7. Dark field TEM micrograph of Rh –20Ti heat treated at 1350 °C for 4 h.

Fig. 8. Temperature dependence of yield strength of Rh – Ti alloys.

three-dimensional network structure. The bright field and dark field micrographs in Fig. 6b and c show the semicoherent precipitates containing what appear to be either fringes or interfacial misfit dislocations. Selected area diffraction (SAD) patterns of the discshaped precipitates in the annealed Rh– 15Ti alloy consist of sharp spots suggesting that there is good lattice matching between the precipitates and the matrix. These precipitates have grown considerably and have become semicoherent discs after annealing. The Rh –20Ti samples annealed at 1350 °C for 4 h did not show the rapid growth of precipitates observed in the Rh –15Ti alloy. As shown in Fig. 7, the precipitate size increased to about 20 nm but no change in precipitate morphology is seen. Furthermore, the selected area

diffraction (SAD) from these annealed alloys displayed ring patterns. This indicates the absence of coherency for the precipitates in the Rh(Ti) matrix. Table 3 shows the bend strength and hardness values of the two phase Rh –15Ti alloy and the single phase Rh –25Ti. The values for a gamma TiAl and a Ni-base superalloy are given for comparison. The single phase Rh –25Ti alloy has a comparable bend strength and hardness to the Ni-base superalloy. More importantly, the two phase Rh–15Ti alloy has significantly higher bend strength than the Ni-base superalloy. The specific strengths and hardnesses of the single and two-phase Rh alloys are compared to a gamma TiAl and Ni-base superalloys in Table 3. The strength of the single phase Rh –25Ti alloy is less than that of the Ni-base superalloy;

Table 3 Mechanical properties of Rh–Ti alloys Alloy

Bend strength (MPa)

Sp. Strength (MPa (g cm−3)−1)

VHN (kg mm−2)

Sp. Hardness (VHN (g cm−3)−1)

Gamma TiAl Ni-based alloy Rh–25 at.% Ti Rh–15 at.% Ti

420 840 881 1269

116 94 84 113

280 320 349 445

77 36 33 40

Fig. 9. Load-deflection curves obtained from three-point bend tests for (a) Ti – 44Al – 11Nb, (b) Rh– 15Ti, and (c) Rh – 25T.

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however, the Rh–15Ti two phase alloy does exhibit higher density compensated strength, approaching that of Gamma Titanium Aluminides. The advantage of these alloys over Gamma Titanium Aluminides is their ultrahigh temperature capability and also the potential for significant room temperature ductility. Fig. 8 shows the temperature dependence of yield strength in the Rh–25Ti alloy as compared to Mar M247 and a gamma TiAl. Rh – 25Ti has a superior strength while exhibiting the same anomalous behavior at intermediate temperatures as other L12 type compounds. The load-deflection curves obtained from three-point bend tests shown in Fig. 9 clearly demonstrate the increased ductility of Rh– 25Ti. While there is virtually no plastic deformation region for the Ti– 44Al –11Nb alloy, a significant plastic deformation region can be seen on the load-extension curves for Rh– 15Ti and Rh –25Ti. Further evidence of this ‘ductile’ behavior in the Ll2 Rh3Ti intermetallic phase was seen in the dislocation substructures and slip line morphology [4]. Although tensile tests are needed to quantify this ductility, it is encouraging to see plastic deformation in these alloys since ductility and fracture toughness are often the limiting factors for the application of intermetallic materials.

4. Summary and conclusions The potential of Rh– Ti alloys for ultra-high temper-

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ature applications has been investigated in this study. The microstructures of Rh–15Ti and Rh– 20Ti identified by TEM and X-ray diffraction consist of Rh3Ti (L12) precipitates in Rh–Ti solid solution matrix. The precipitates in the heat treated Rh–15Ti alloy are disc-shaped and they tend to align themselves orthogonally. The intermetallic precipitates in Rh–20Ti did not transform into discs but remained as spheres after heat treatment. Preliminary results of the density compensated strength and hardness properties of the two-phase Rh–15Ti alloy demonstrate that this alloy is equal to or superior to nickel-base superalloys, while offering the advantage of higher temperature capability. The alloys displayed significant ductility in three-point bend tests.

Acknowledgements This work is supported by the Office of Naval Research.

References [1] Y. Yamabe, Y. Koizumi, H. Murakami, Y. Ro, T. Maruko, H. Harada, Scripta Materialia 35 (1996) 211. [2] V. Paidar, D.P. Pope, V. Vitek, Acta Metall. 32 (1984) 435. [3] T. Sims Chester, C. Hagel William (Eds.), The Superalloys, John Wiley & Sons, 1972, p. 123. [4] R. Mahapatra, A.W. Davis, S.M.L. Sastry, Unpublished Results.