Mechanical properties of Rh-based L12 intermetallic compoundsRh3Ti, Rh3Nb and Rh3Ta

Seiji Miura a,*, Keiichi Honma b, Yoshihiro Terada c, J.M. Sanchez d, Tetsuo Mohri a

aDivision ofMaterials Science andEngineering,Graduate School of Engineering,HokkaidoUniversity,Kita-13,Nishi-8, Kita-ku, Sapporo 060-8628, JapanbGraduate School of Engineering, Hokkaido University, (now Mitsubishi Material Corporation) Saitama, Japan

cDepartment of Metallurgy and Ceramics Science, Tokyo Institute of Technology, Meguro-ku, Tokyo 152-8552, JapandTexas Materials Institute, The University of Texas, Austin, TX 78712, USA

Received 24 November 1999; received in revised form 16 February 2000; accepted 17 February 2000

Abstract

A preliminary study of the mechanical properties of Rh-based L12 intermetallic compounds Rh3Ti, Rh3Nb and Rh3Ta is carriedout by micro-vickers hardness tests at room temperature and compression tests at various temperatures ranging from room tem-perature to 1673 K. A cold-rolling test is also conducted on a Rh3Ti plate. Although high temperature ductility-loss is observed inall compounds, Rh3Ti shows good ductility at all temperatures investigated. Both Rh3Nb and Rh3Ta show a weak positive tem-

perature dependence of strength (stress anomaly) at around 1273 K which is about half the melting point for both intermetalliccompounds. The stress anomaly is discussed in terms of a phase stability concept based on the Kear±Wilsdorf (K±W) mechanism.High work hardening rates of Rh3Nb and Rh3Ta, which cause high vickers hardness at room temperature, are also attributed to the

K±W mechanism. # 2000 Elsevier Science Ltd. All rights reserved.

Keywords: A. Intermetallics, miscellaneous; B. Brittleness and ductility; B. Plastic deformation mechanisms; E. Phase stability, prediction

1. Introduction

Application of intermetallic compounds for hightemperature materials has been hampered mainly bytheir poor ductility. One of the reasons is that most ofthe high melting-point intermetallic compounds havecomplex crystal structures with a limited number of slipsystems such as A15 and D88. Yamabe-Mitarai et al.have proposed a new class of alloys based on highmelting-point fcc metals, such as Ir and Rh, with L12intermetallic compounds as dispersions forming a similarmicrostructure to the commercial Ni-based superalloys[1±4]. It was shown that the decrease of strength is verygentle in some of these two-phase alloys until 1773 Kwithout a sudden drop of strength observed in Ni-basedsuperalloys at around 1200 K, and their compressiveductilities are also promising.Table 1 presents the melting points and densities of

pure Ir metal, pure Rh metal and their derivatives with

the L12 crystal structure. The density of each inter-metallic compound with stoichiometric composition isestimated from the lattice parameter and atomic weights[5,6]. Although pure Rh and its derivatives have ratherlower melting points compared to pure Ir and its deri-vatives [7], the Rh group possesses a large advantage ondensity. Moreover, it has been known that the oxidationresistance of Rh is superior to Ir [2]. However, funda-mental studies on the mechanical properties of both Ir-based and Rh-based L12 intermetallic compounds, suchas Ir3Nb or Rh3V, are still limited [1±4,8±10].Among the Ir- and Rh-based L12 compounds only

Ir3Nb is reported to show a strong positive temperaturedependence of strength (a stress anomaly) as Ni3Al andCo3Ti do [4]. A systematic investigation of the stressanomaly has been carried out on Ni and Pt based L12intermetallic compounds with various B-subgroup ele-ments as a minor component and it was revealed thatthe appearance of the stress anomaly can be explainedin terms of the phase stability of the L12 structure withrespect to other crystal structures [11±13]. On the otherhand, among L12 intermetallic compounds composed ofonly transition elements, a few of compounds, such as

0966-9795/00/$ - see front matter # 2000 Elsevier Science Ltd. All rights reserved.

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Intermetallics 8 (2000) 785±791

* Corresponding author. Tel.: +81-11-706-6347; fax: +81-11-706-

7812.

E-mail address: miura@eng.hokudai.ac.jp (S. Miura).

Co3Ti, (Co,Fe)3V and Ir3Nb, are known to show astrong stress anomaly [4,14±16]. As Rh3Ti, Rh3Nb andRh3Ta intermetallic compounds are expected to showhigh strength and high toughness from results of a seriesof screening tests on Rh- and Ir-based L12 intermetalliccompounds by Terada et al. [17], their mechanical prop-erties are investigated in the present study and attemptedto be explained by the phase stability concept.

2. Experimental procedures

The nominal composition for each alloy was 75.3 at%Rh- 24.7 at% X (X : Ti, Nb, Ta) in this study. All thespecimens were arc-melted several times in an Ar atmo-sphere on a water-cooled copper hearth from 99.9% Rhand 99.5% Ti, 99.5% Nb or 99.9% Ta purity. Alloyingots, being about 30 g, were heat-treated for homo-genization at 1773 K for 24 h in an Ar-¯ow atmosphere.Microstructural observation by scanning electronmicroscopy (SEM) and chemical analysis by wavelengthdispersive X-ray spectroscopy (WDS) with pure Rh, Ti,Nb and Ta as references were carried out on polishedsamples. Grain size of each ingot range from 200 to 300mm without any precipitates. The deviations from thenominal composition detected by WDS were less than0.1 at% for both Rh3Nb and Rh3Ta. Although thedeviation was about 1 at% for Rh3Ti, its compositionwas determined to be 74.3 at.% Rh-25.7 at% Ti, stillwithin the Rh-lean side of the L12 single phase ®eld.Compression test specimens with 3�3�6 mm3 in

dimension were cut from the alloy ingots by a wheelcutter and polished with emery papers. An instron-typetesting machine was used for the compression tests atvarious temperatures ranging from room temperature to1673 K with an initial strain rate of 2.8�10ÿ4 sÿ1. Athigh temperatures the compression tests were conductedunder an Ar-¯ow atmosphere. Strain rate was alter-nately changed several times to 2.8�10ÿ3 sÿ1, ten timeshigher than the initial strain rate, during plastic defor-mation.A Rh3Ti plate for a cold-rolling test with a size of 1.25

mm in thickness, 6.5 mm in width and 30 mm in lengthwas also prepared from the homogenized ingot andannealed at 773 K for 1 day in vacuum to reduce

expected hydrogen contamination. Micro-vickers testswere also carried out on all the alloys at the room tem-perature with a load of 300 g for 30 s.

3. Results

3.1. Temperature dependence of 0.2% ¯ow stress andwork hardening rate

Stress±strain curves of Rh3Ta tested at various tem-peratures are shown in Fig. 1 as a typical example. It isobvious that the work hardening rate is high and asudden drop of stress by cracking occurs after severalpercent of plastic deformation at a temperature rangelower than 1273 K. In the high temperature region asteady-state deformation similar to a creep deformationbehavior is observed. Rh3Nb shows similar deformationbehavior to that of Rh3Ta in the entire temperaturerange, and Rh3Ti has lower work hardening rates andlarger compressive ductility. Small strain rate depen-dence of strength at the temperature range lower than1273 K is common for all the alloys.The temperature dependence of 0.2% ¯ow stress of

each compound is presented in Fig. 2. Results of otherL12 intermetallic compounds are redrawn from the lit-erature [12]. Both Rh3Nb and Rh3Ta show a weakpositive temperature dependence of strength at around1273 K, whereas the strength of Rh3Ti decreasesmonotonically with increasing temperature. Wee andSuzuki reported the temperature dependence of micro-vickers hardness of Rh3Ta up to 1200 K from which atendency can be observed that the slope of its negativetemperature dependence of hardness becomes gentle ataround 1200 K [11]. This is consistent with the weakpositive temperature dependence of the 0.2% ¯ow stressobserved in the present study by compression tests.It has been known that the peak strength for Ni3Al is

located at around half the melting point (Tm) [18,19]. InFig. 3, the 0.2% ¯ow stresses of all compounds areplotted against the temperature normalized by Tm. Thepeak strengths for Rh3Nb and Rh3Ta are also located ataround half Tm. This suggests that the controllingmechanism of the stress anomaly observed in the Rh-based compounds is the same as in other L12 compounds,

Table 1

Properties of pure Ir, pure Rh and their stoichiometric derivatives

Ir Rh Ir3Nb Ir3Ta Rh3Ti Rh3V Rh3Nb Rh3Ta

Melting point (K) 2720 2236 2708 2727 �2023 2013 �2223 �2398Density (Mg/m3) 22.5 12.44 18.95 21.45 10.60 11.03 11.63 14.14

Electron±atom ratio, e/a ± ± 8.0 8.0 7.75 8.0 8.0 8.0

Atomic size factora [11] ± ± 0.09 0.09 0.09 0.01 0.10 0.10

a The atomic size factor is estimated by �rA ÿ rB�=rA, in which rA and rB are Goldschmidt's atomic radius of major element A and minor element

B, respectively.

786 S. Miura et al. / Intermetallics 8 (2000) 785±791

i.e. the Kear±Wilsdorf (K±W) mechanism. At tempera-tures higher than the peak temperatures, the stressdrops and another thermally assisted process(es) seemsto govern the deformation. The work hardening rate(WHR) of Ni3Al is also known to change with tem-perature, in a manner similar to the changes of strength[18,19]. Fig. 4 indicates temperature dependence ofWHR of the compounds at 1% plastic strain. TheWHR is identi®ed to be high for each compound and itstarts to decrease at the temperature corresponding tothe peak of strength for Rh3Nb and Rh3Ta. These

relationships between the temperature dependence ofstrength and WHR are similar to those of Ni3Al alloys[19], which also suggests the operation of the K±Wmechanism in Rh3Nb and Rh3Ta.Table 2 summarizes the results of micro-vickers

hardness tests at room temperature together with theresults reported by Wee and Suzuki [11]. In both resultshardness of Rh3Nb and Rh3Ta are much higher thanthose of Rh3Ti and Ni3Al, which is attributed to thehigh WHR of Rh3Nb and Rh3Ta.

Fig. 1. Stress±strain curves of Rh3Ta tested at various temperatures.

Fig. 2. Temperature dependence of 0.2% ¯ow stress of Rh-based

compounds investigated in the present study. Reported data of other

L12 compounds are also shown [12].

Fig. 3. Normalized temperature dependence of 0.2% ¯ow stress of

Rh-based compounds investigated in the present study. Reported data

of other L12 compounds are also shown [12].

S. Miura et al. / Intermetallics 8 (2000) 785±791 787

3.2. Temperature dependence of compressive ductility

Figs. 5 and 6 indicate the temperature dependence ofthe compressive ductility and the maximum strength ofthe compounds. The ductility of Rh3Ti is much higherthan those of Rh3Nb and Rh3Ta at any test tempera-tures except for 1073 K. It is noteworthy that both thecompressive ductility and the maximum strength ofRh3Ti show the minima at 1073 K, and no cracking wasobserved in the Rh3Ti specimens tested at the tempera-ture range higher than 1273 K. The compressive ducti-lities of both Rh3Nb and Rh3Ta are small, but bothcompounds show high maximum strength in accordancewith their own high WHRs.Shown in Fig. 7 are a series of photos of Rh3Ti spe-

cimens before and after the compression test at varioustemperatures. Specimens tested at the temperaturerange lower than 1273 K were deformed until the sud-den drop of stress caused by cracking. The deformationof each specimen is uniform, however as shown in Fig. 5the deformability at 1073 K is much lower than thattested at other temperatures.Compressive ductility of Rh3Ti at room temperature

was also examined by cold-rolling with about 2%

reduction in thickness by each pass. Fig. 8 shows a seriesof photographs of the Rh3Ti specimen before cold-rollingand after 9 and 12% reduction in thickness. Cracksinitiate at each side of the plate sample after severalpasses, and they seem to propagate through grainboundaries. The sample broke when its reduction inthickness reached 12%.

Table 2

Vickers hardness of various L12 compounds at room temperature

Rh3Ti Rh3Nb Rh3Ta Ni3Al

Present study 252 765 710 ±

From the literature [11] 220 500 500 170

Fig. 4. Temperature dependence of the work hardening rate (WHR)

of compounds investigated.

Fig. 6. Temperature dependence of compressive maximum strength.

Fig. 5. Temperature dependence of compressive ductility of com-

pounds.

788 S. Miura et al. / Intermetallics 8 (2000) 785±791

4. Discussion

4.1. Stress anomaly and the Kear±Wilsdorf mechanism

As shown in Figs. 2 and 3 it is revealed that bothRh3Nb and Rh3Ta show a weak positive temperaturedependence of strength. As the compounds investigatedin the present study have no order±disorder transition[7], the stress anomaly can not be explained by thechange in order parameter as is the case of CuZn andFeCo [20]. Details of the controlling mechanism of the

stress anomaly of the L12 intermetallic compounds atintermediate temperature range are still controversial[21±23]. However, the Takeuchi±Kuramoto modelbased on the Kear±Wilsdorf mechanism, i.e. a cross-slipof a part of screw dislocations from {111} to {001}, hasbeen accepted as a basis for the stress anomaly [24,25].It has been argued that the thermal activation of thecross-slip behavior causes the positive temperaturedependence of strength in various L12 structures. In themodel, two sets of Shockley partials are assumed toform a super-dislocation in the L12 structure. It can be

Fig. 7. Rh3Ti specimens before and after compression tests.

Fig. 8. A Rh3Ti specimen before and after cold-rolling.

S. Miura et al. / Intermetallics 8 (2000) 785±791 789

expected that the dislocation behavior is a�ected by theenergies of planar faults such as a complex stackingfault (CSF) between a pair of Shockley partials and ananti-phase boundary (APB) between the leading andtrailing dislocations; by increasing the energy of CSFbetween a pair of Shockley partial dislocations on {111}plane in the L12 crystal structure, they tend to form aperfect (screw) dislocation, resulting in the cross-sliponto the {001} plane. Similarly, with decreasing theenergy of APB on the {001} plane, the leading disloca-tion tends to cross-slip onto the {001} plane. Each of theplanar faults is related to a certain crystal structurederived from the L12 structure by a certain rearrange-ment of con®guration of atoms in close resemblance tothe fact that the stacking fault on {111} in fcc is relatedto hcp. Therefore, the energies of planar faults should bea�ected by the phase stability among competing phases.To explain the stress anomaly, Suzuki and his co-work-ers applied the concept of phase stability of the L12structure with respect to other crystal structures, whichis related to the planar faults accompanying the super-dislocations [11]. In this concept, the phase stability ofthe L12 structure against the DO19 structure is related tothe CSF energy in the L12 crystal structure, and theDO22 type stacking in the L12 structure is treated as theAPB on {001}. This argument claims that low stabilityof the L12 structure is necessary for the appearance ofthe stress anomaly. Hence the phase stability of geome-trically close-packed (GCP) phases including L12, DO19

and DO22 crystal structures should be evaluated.Suzuki and his co-workers showed that the size of

elements characterized by Goldschmidt's atomic radiusfor co-cordination number 12, which is signi®cant forthe stability of GCP phases [26], provides a goodexplanation for the tendency of the stress anomaly of Niand Pt-based L12 compounds with B-subgroup ele-ments [12,13]. The stress anomaly is apparentlyobserved only in the L12 alloys with the atomic sizefactor, �rA ÿ rB�=rA, in which rA and rB are Goldschmidt'satomic radius of major element A and minor element B,respectively, at a certain size range which is close to theboundary between the L12 and other structures. It showsthat the relative instability of the L12 structure againstother structure is inevitable for the appearance of thestress anomaly. However, there is no signi®cant di�erenceamong the atomic size factor of the present alloys.The electron±atom ratio (e/a) is another factor which

governs the stability of GCP phases [11,12,27±29]. Fig.9 illustrates the variation of e/a with the change ofmajor element A of A3B type intermetallic compoundsalong the second row of the periodic table. The elec-tron±atom ratio (e/a) varies from 7 to 9.25 in Fig. 9(a)with Ti as a minor element B and from 7.25 to 9.5 inFig. 9(b) and (c) with Nb and Ta, respectively. Cubicand hexagonal type stackings are represented by ``c'' and``h'', respectively [26], and rectangular and triangular

arrangements of minor element B of A3B compounds inclose-packed plane are denoted by ``R'' and ``T'',respectively [27,28]. In this notation the L12, the D019and the D022 structures are represented as `cT', `hT' and`cR', respectively. It is noticed that with increasing e/a,the crystal structure varies from `cT' (L12) to `chT'(D024) to `cR' (D022), indicating that the hexagonalstacking and/or the R-type arrangement of minor atomsbecome stable [11]. Also the Rh-based L12 crystalstructure composed of Ti is relatively stable comparedto those composed of Nb or Ta. This explains the morepronounced tendencies of the stress anomaly of Rh3Nband Rh3Ta than that of Rh3Ti.To pursue the extension of the phase stability concept

to this kind of compound, a series of alloys with a widerange of atomic size factor without changing e/a isrequired. An addition of V to Rh3Nb and Rh3Ta wouldmeet such a requirement, and this remains for futurework.

4.2. Strain rate sensitivity of deformation stress

As shown in Fig. 1, in the temperature range higherthan 1273 K, the strain-rate dependence of stressbecomes noticeable in the compounds investigated,whereas it is very small in the lower temperature range.However, a steady-state deformation was observed onlyin Rh3Ta alloy at 1673 K. We applied the followingequation to evaluate the creep property of the inter-metallic compound Rh3Ta;

": � A�n �1�

where A is a constant, ":the strain rate, � the stress and

n the stress exponent. The stress exponent n for Rh3Ta

Fig. 9. The variation of e/a by changing major element A of A3B type

intermetallic compounds along the second row of the periodic table.

The electron±atom ratio (e/a) varies from 7 to 9.25 in (a) with Ti as a

minor element and from 7.25 to 9.5 in (b) and (c) with Nb and Ta as a

minor element, respectively. Compounds which do not appear in the

equilibrium phase diagram are shown with `` ''.

790 S. Miura et al. / Intermetallics 8 (2000) 785±791

is evaluated to be 10.1 at 1673 K. The stress exponentfor L12 compound Ni3Al ranges from 3 to 5, which wasevaluated by steady-state creep tests and the steady-state deformation behavior of compression or tensiletests by instron-type testing machine [30]. The valueobtained by the present study is rather higher. Highwork hardening rate is observed even at 1473 K and itimplies that a low temperature deformation mechanismother than the creep mechanism still contributes to thedeformation even at 1673 K.

4.3. Temperature dependence of compressive ductilityand compressive maximum strength

As shown in Fig. 5 the maximum value of compres-sive ductility appears at around 1273 K for both Rh3Nband Rh3Ta, but their deformabilities are not yet su�-cient for most structural applications. On the otherhand, although the cold-rolling property of Rh3Ti atroom temperature is not excellent, its deformability isstill attractive. In all of A3B-type L12 compounds duc-tility is observed only with the composition in themajor-element ``A'' rich side. Interesting thing is thatRh3Ti shows small but not ignorable deformability evenwith a Rh-lean composition.As pointed out above, Rh3Ti shows a ductility loss at

1073 K. The ductile L12 compounds such as Co3Ti, (Co,Fe)3V and Ni3Al+B have been reported to show similarhigh temperature embrittlement, or a ductility-loss, ataround 1000 K, which almost agree with the peak tem-perature of strength [31]. In that regard, the ductility-lossobserved in both Rh3Nb and Rh3Ta at around 1500 Kare similar. The high temperature embrittlement isattributed to oxygen penetration in Co3Ti andNi3Al+B [16,32,33]. An observation of the fracturesurface of tensile or bending specimens tested at theductility-loss temperature range would identify the reasonfor this embrittlement. At the higher temperature rangethe deformation behavior of Rh3Ti seems to be governedby a creep-like mechanism. It is e�ective to alleviate thestress concentration, resulting in the high ductility.

5. Concluding remarks

Mechanical properties of Rh-based L12 intermetalliccompounds Rh3Ti, Rh3Nb and Rh3Ta are investigated.High micro-vickers hardnesses of Rh3Nb and Rh3Ta atroom temperature are due to their high work hardeningrates which are attributed to the Kear±Wilsdorfmechanism. Both Rh3Nb and Rh3Ta show a weakpositive temperature dependence of strength but Rh3Tidoes not; these tendencies could be explained by theconcept of phase stability. Rh3Ti shows good ductilityup to 30% in compression at almost all temperature

ranges investigated, however, rolling ductility at roomtemperature is not excellent. Embrittlement of Rh3Ti at1073 K and those of both Rh3Nb and Rh3Ta at around1500 K seem to be a common phenomenon in ductileL12 compounds.

Acknowledgements

The authors wish to thank Mr. Kenji Ohkubo fortechnical assistance.

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