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Nb5Si3 in situ composites alloyed with Mo
Intermetallics 9 (2001) 521–527 www.elsevier.com/locate/intermet
Microstructure and room temperature deformation of Nbss/Nb5Si3 in situ composites alloyed with Mo Won-Yong Kima,*, Hisao Tanakaa, Akio Kasamaa, Ryohei Tanakaa, Shuji Hanadab a
Japan Ultra-High Temperature Materials Research Institute(JUTEMI), 573-3 Okiube, Ube, Yamaguchi 755-0001, Japan b Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan Received 14 March 2001; received in revised form 13 April 2001; accepted 16 April 2001
Abstract Room temperature deformation behavior of Nb5Si3 based intermetallics in the ternary Nb–Si–Mo alloy system is investigated by compression testing in relation to microstructure and compositional effect. The partial ternary phase diagram of the Nb-Si-Mo alloy system containing phase equilibrium information at 1973 K is determined by metallography, X-ray diffraction and scanning electron microscopy (SEM) equipped with wavelength-dispersive X-ray fluorescence spectroscopy (WDS). A pseudo-binary compound of Nb5Si3–Mo5Si3 is confirmed to form at Nb5Si3-rich compositions by Mo addition, without having a solubility range of Si. The phase transformation from a-Nb5Si3 to b-Nb5Si3 is found to occur at about 5 at.% of Mo content on the Nb5Si3–Mo5Si3 pseudo-binary line. The observed yield stresses are largely dependent on not only the volume fraction and morphology of constituent phases but also the strength of Nb5Si3 phase equilibrating with Nb solid solution. The in situ composites consisting of aNb5Si3 phase and bcc solid solution exhibit higher yield stress than those consisting of b-Nb5Si3 phase and bcc solid solution when yield stress is compared at the same volume fraction of bcc phase. It is suggested that room temperature deformability of b-Nb5Si3 is superior to that of a-Nb5Si3 phase. Room temperature deformation behavior of the present composites will be discussed in relation to microstructural evolution. # 2001 Elsevier Science Ltd. All rights reserved. Keywords: A. Silicides, various; B. Phase transformations; B. Phase diagram; B. Twinning
1. Introduction Niobium silicide base alloys have a strong potential as high temperature structural materials due to their high melting point over 2273 K and excellent strength even at extremely high temperature above 1473 K [1–4]. Nb5Si3 displays a complex tetragonal D8m at high temperature and a D8l at low temperature representing the space group I4/mcm (140) for both crystal structures [5]. The lattice parameters of both phases are a=1.0018 nm and c=0.5072 nm for the stoichiometric high temperature phase (b-Nb5Si3, W5Si3-type), and a=0.6570 nm and c=1.1884 nm for the stoichiometric low temperature phase (a-Nb5Si3, Cr5B3-type), respectively. Owing to their complex crystal structures, plastic deformation of monolithic Nb5Si3 will be limited even in the high temperature range. An incorporation of a ductile phase to hard-deformable Nb5Si3 phase resulting in two-phase * Corresponding author. Fax: +81-836-51-1765. E-mail address: [email protected] (W.-Y. Kim).
silicide alloy will be one way to improve room temperature deformability. Concerning the role of bcc phase in the Nbss/Nb5Si3 in situ composites (where Nbss means niobium solid solution), it is expected that microcracks introduced at a production stage, or flaws formed due to thermal stress during cooling may be reduced or removed by incorporating the ductile bcc phase. Additionally, the ductile phase in the Nbss/ Nb5Si3 in situ composites may assist deformation of the hard phase by a role of strain accommodation if the alloys are plastically deformed at room temperature. In some Laves phase alloy systems [6–8], the brittle to ductile transition temperature (BDTT) of monolithic Laves phase is found to be higher than 1273 K like other complex-basis intermetallic compounds such as Mo5Si3, Nb3Al, Cr3Si and V3Si [9–12]. Nevertheless, it has been observed that the complex-basis intermetallics such as HfV2+Nb or Ta with two-phase microstructures consisting of bcc solid solution and Laves intermetallic phase deform by mechanical twinning at room temperature [6–8]. These reports were very
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attractive for alloy design of intermetallics, and hence most of studies have been concentrated on the incorporation of ductile phase into brittle intermetallic phase. The brittleness of intermetallics is still one of the most critical issues to develop intermetallics or intermetallic base materials as next generation high temperature materials. In the present study, we aim at correlating the room temperature deformation behavior with phase transformation of Nb5Si3 and solid solution hardening in Nb–Si–Mo alloys consisting of Nb5Si3 and bcc solid solution. Mo is chosen as the ternary alloying element in the Nbss/Nb5Si3 in situ composites to increase melting point, to change phase stability and to strengthen the composites by solid solution hardening.
2. Experimental procedure Various ternary Nb–Si–Mo alloys were prepared by arc-melting under an argon atmosphere in a watercooled copper hearth. The purities of raw materials were 99.9 wt.% of Nb, 99.99 wt.% of Si, 99.9 wt.% of Mo, respectively. The alloy buttons were re-melted at least 5 times to ensure chemical homogeneity. Homogenization heat treatment was carried out at 1973 K for 100 h in a high purity nitrogen atmosphere, followed by rapid cooling to room temperature at a rate of 200 K/ min in the high temperature region (1973 to 773 K) and then at a rate of 10 K/min in the low temperature region 773 K to room temperature). Some of the arc-melted samples were heat treated at 2123 K for 24 h in order to investigate the effect of Mo alloying on phase stability and mechanical property in the Nb–Si–Mo ternary alloy system. Samples for chemical composition analysis, microstructural observation and compression testing were made by an electro-discharge machine (EDM) from the heat treated samples. X-ray diffraction experiment
was conducted on the samples heat treated at 1973 and 2123 K to examine the crystal structures of constituent phases. The nominal compositions of alloys studied in the present study are shown in Table 1. Chemical compositions of the constituent phases are analyzed using a scanning electron microscope equipped with a wavelength-dispersive X-ray fluorescence spectroscope. Compression specimens with 2.52.5 mm cross-section and 6 mm height were mechanically polished using SiC paper and Al2O3 particles with water. Compression tests were carried out using an Instron model 8500 mechanical testing machine under a nitrogen atmosphere at room temperature and at an initial strain rate of 310 4 s 1. To observe the deformation microstructure by transmission electron microscopy (TEM), some of the specimens were deformed to 3% plastic strain, and then unloaded. Slices for TEM specimens (1.5 mm) were prepared by EDM and mechanically polished to 100 mm thickness. Final thinning was accomplished using the twin jet electropolishing technique (Struers Tenupol) to make a perforation of samples in a chemical solution of 94 ml of methanol and 5 ml of H2SO4 and 1 ml of HF. TEM observation was made in a Jeol 4000EX microscope operated at 400 kV.
3. Results and discussion Partial ternary phase diagram of Nb–Si–Mo alloy system investigated at 1973 K is shown in Fig. 1. The dotted line in the two-phase region of the phase diagram indicates the phase boundary between a-Nb5Si3 phase and b-Nb5Si3 phase. An extension of Nb5Si3 phase toward Mo5Si3 on the pseudo-binary Nb5Si3–Mo5Si3 line is found without a deviation from the line, indicating that Mo as a ternary alloying element substitutes for Nb rather than Si. From the determined tie-lines in the
Table 1 The nominal compositions (at.%) of 14 kinds of alloys and corresponding microstructural characteristics in Nb–Si–Mo system Sample no.
Nb
Si
Mo
Constituent phases
Volume fraction of Nb5Si3 (%)
Heat treatment
a b c d e f g h i j k l
79 69 59 77 67 57 75 65 55 73 63 53
16 16 16 18 18 18 20 20 20 22 22 22
5 15 25 5 15 25 5 15 25 5 15 25
a-Nb5Si3 a-Nb5Si3 b-Nb5Si3 a-Nb5Si3 a-Nb5Si3 b-Nb5Si3 a-Nb5Si3 b-Nb5Si3 b-Nb5Si3 a-Nb5Si3 b-Nb5Si3 b-Nb5Si3
and Nbss and Nbss and Nbss and Nbss and Nbss and Nbss and Nbss and Nbss and Nbss and Nbss and Nbss and Nbss
38 38 38 45 45 45 52 52 52 57 57 57
1973 K for 1973 K for 1973 K for 1973 K for 1973 K for 1973 K for 1973 K for 1973 K for 1973 K for 1973 K for 1973 K for 1973 K for
e1 h1
67 65
18 20
15 15
b-Nb5Si3 and Nbss b-Nb5Si3 and Nbss.
45 52
2123 K for 48 h 2123 K for 48 h
100 h 100 h 100 h 100 h 100 h 100 h 100 h 100 h 100 h 100 h 100 h 100 h
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Fig. 1. Partial Nb–Si–Mo phase diagram at 1973 K in the present study. The open circles represent the alloy compositions for alloys in the duplex phase field consisting of Nb5Si3 phase and Nb solid solution; the solid circles correspond to maximum solubility in each phase.
phase diagram, it is suggested that Mo is more predominantly partitioned into Nb solid solution than Nb5Si3 phase in the two-phase Nbss/ Nb5Si3 region. Nb solid solution is extended along the Nb–Mo binary phase line with a small amount of Si solubility less than 1 at.%. The solubility of Si in Nb solid solution is not so sensitive to Mo content. These results are very consistent with the fact that the miscible solid solution is formed between Nb and Mo, and the solubilities of Si in Nb and Mo solid solution are very small in a similar manner to both Nb–Si and Mo–Si binary alloy systems. The Si content of a ternary eutectic consisting of Nbss and Nb5Si3 increases slightly compared to that of binary Nb–Si alloy system. Back scattered electron images (BEI) of Nb–18at.%Si–15at.%Mo and Nb–20at.%Si– 15at.%Mo are shown in Fig. 2. The lightly contrasted phase is Nb solid solution, while the darkly contrasted phase is Nb5Si3 phase, as indicated by arrows in Fig. 2. Clearly, it is observed that hypo-eutectic is changed to hyper-eutectic microstructure with increasing Si content, as seen in Fig. 2(a) and (c). The Si content of the eutectic is estimated from the microstructures to be between 18 and 20 at.% within the composition range investigated. With increasing Mo content at constant Si content, the volume fraction of primary bcc solid solution increased, the size of bcc particles increased and bcc volume fraction in a eutectic decreased. However, there was no change in volume fraction of the constituent phases at constant Si content, indicating that the volume fraction of the phases depends mostly upon Si content. The results of X-ray diffraction on several samples heat treated at both 1973 and 2123 K are shown in Fig. 3. The open squares in the figure indicate peaks due to a-Nb5Si3, and the closed circles indicate peaks from
Fig. 2. Back-scattered electron images of Nb–18at.%Si–15at.%Mo heat treated at 1973 K (a) and at 2123 K (b), and Nb–20at.%Si– 15at.%Mo alloy heat treated at 1973 K (c). All the microstructures consist of Nb5Si3 and bcc solid solution. Note that a microstructural transition from hypo-eutectic to hyper-eutectic occurs at a composition between (a) and (c).
the bcc solid solution. The b-Nb5Si3 phase is marked by the closed squares. While most diffraction peaks are common to the two Nb5Si3 phases, the 112 and 202 peaks of the a-Nb5Si3 phase and the 211 and 310 peaks of the bNb5Si3 phase allow their presence to be uniquely determined. We could not find an intermediate Nb3Si phase in the present alloys whatever heat treatment was performed at 1973 or 2123 K. Therefore, Mo addition to the
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Nb5Si3/Nbss in situ composite makes Nb3Si unstable. In the case of samples heat treated at 1973 K, it is found that the stability of the Nb5Si3 phases is changed depending on Si and Mo contents. b-Nb5Si3 is stable in the high Si and Mo region, while a-Nb5Si3 was stable in the low Si and Mo region. In the case of samples heat treated at 2123 K, b-Nb5Si3 is found to be stable over the whole composition range examined in the present study. In the binary Nb–Si phase diagram, Nb5Si3 phase undergoes a phase transformation from a to b depending on temperature and composition. Further the crystal structure of Mo5Si3 forming the pseudo-binary phase with Nb5Si3 is the same as b phase. Therefore, as seen in the alloys h and k of Fig. 1, the accelerated phase transformation from a to b phase at 1973 K by addition of Mo may be associated with a change of phase stability, since a-Nb5Si3 phase is stable at 1973 K in the binary Nb–Si phase diagram [13]. Assuming that the phase transformation from a to b in the present Nb–Si–Mo ternary alloy system involves a diffusion process, transformation kinetics should be considered to explain the alloying effect. Since Mo has a higher melting point than Nb or Si, it can not be explained by enhancement of the diffusion process due to Mo addition. In Fig. 4(a), the lattice parameters of the bcc phases of the several alloys are plotted as a function of Mo content in the bcc phase in the two-phase alloys. A monotonic decrease in lattice parameter is seen,
although it appears not to be linear. From the atomic sizes of Nb, Mo and Si being 0.147, 0.140 and 0.116 nm, respectively, it is clear that the largest Nb atom is replaced by Mo with increasing Mo content. In Fig. 4(b), the c/a ratio in the unit cell of a-Nb5Si3 phase is plotted as a function of Mo content in the a-Nb5Si3 phase. Since the atomic size of Mo is smaller than Nb and the pseudo-binary line extends without a deviation of Si content from the stoichiometric composition, Mo would substitute for Nb sites, thereby leading to smaller crystal size or reducing c/a ratio of tetragonal Nb5Si3 phase. However the calculated c/a ratio is somewhat different from our expectation. It is found that with increasing Mo content, a-axis is shortened, whereas caxis is lengthened, resulting in an increase of c/a ratio in the a-Nb5Si3 phase. At the present time, the propensity of c- and a-axial length change is puzzling, but probably the anisotropy of inter-atomic bonding force between either unlike or like atoms is responsible for the substitution behavior of ternary alloying elements in this alloy system [14–15]. Two kinds of chain structures, –Si– Mo–Si– and Mo–Mo or –Si–Si–, were reported in Mo5Si3 having the same crystal structure (body-centered tetragonal, tI32) as Nb5Si3. Chu et al. [15] have demonstrated that the strong anisotropy of thermal expansion and hardness arises from the bonding along [100] direction corresponding to –Si–Mo–Si– chain structure and [001] direction corresponding to –Mo–Mo– or –Si–Si– chain
Fig. 3. X-ray powder diffraction spectra for 6 alloys at a constant Mo content of approximately 15 at.%Mo. Alloys b, e, h and k are heat treated at 1973 K, and alloys e1 and h1 are heat treated at 2123 K. The chemical compositions of alloys el and h1 are the same as alloy e and alloy h, respectively.
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structure. Therefore, the selective substitution behavior of Mo for Nb sites may lead to an increase in c/a ratio, and reduce the stability of a-Nb5Si3 in the present Nb– Si–Mo alloy system. Further investigation is required to elucidate the atomic substitution behavior and associated physical and mechanical properties. Compression tests were carried out on samples heat treated at 1973 K for 100 h and 2123 K for 24 h for both alloys with two-phase microstructures consisting of a-Nb5Si3 or b-Nb5Si3 and bcc solid solution at room temperature and a strain rate of 310 4 s 1. A large number of alloys with compositions in the range from the hypo-eutectic and hyper-eutectic were tested. Within this set of alloys (alloy b, e, h and k), two systematic variations were employed to ascertain the effect of silicide phases
equilibrating with bcc solid solution on mechanical properties, and the effect of volume fraction of each phase. Offset yield stresses (0.2%) are plotted as a function of volume fraction of Nb5Si3 phase, as shown in Fig. 5(a). Concerning the samples heat treated at 2123 K, yield stress increases with increasing volume fraction of b-Nb5Si3 phase. Correspondingly compressive ductility decreases with increasing volume fraction of b-Nb5Si3. A sharp increase in yield stress is seen at 52% of b-Nb5Si3 phase. This result may be associated with microstructural transition from hypo-eutectic (primary bcc+eutectic) to hyper-eutectic (primary b-Nb5Si3+eutectic). In the alloy heat treated at 1973 K, it is found that yield stress decreases with increasing volume fraction of Nb5Si3 phase from 45% (alloy e) to 52% (alloy h).
Fig. 4. Lattice parameters of Nb solid solutions equilibrating with Nb5Si3 phase as a function of Mo content (a), and the c/a ratio of aNb5Si3 versus Mo content at a constant Si content (b). All samples were heat treated at 1973 K for 100 h.
Fig. 5. Yield stress (a) and compressive ductility (b) of alloys tested at room temperature as a function of total volume fraction of b-Nb5Si3 or b-Nb5Si3 phase in each alloy.
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Furthermore, it was confirmed for all alloys studied here that Mo content in the bcc solid solution equilibrating with Nb5Si3 phase increases with increasing volume fraction of Nb5Si3 phases whatever they exist as a or b. Therefore, the yield stress drop at 52% of Nb5Si3 is closely correlated with phase transformation of aNb5Si3 to b-Nb5Si3, suggesting that b-Nb5Si3 deforms at a lower stress than a-Nb5Si3 at room temperature in this Nb–Si–Mo alloy system. It is interesting to note that yield stresses display similar values for the two-phase Nbss/b-Nb5Si3 alloys when the volume fraction of bNb5Si3 phase is not changed significantly by the heat treatment at 1973 K (alloy h) or 2123 K (alloy h1), see yield stresses at 52% (alloy h and h1) volume fraction of Nb5Si3. Thus, the large increase in compressive ductility at 52% of b-Nb5Si3 for the heat treatment at 1973 K, as shown in Fig. 5(b), is considered to be a strong evidence of phase transformation of Nb5Si3 from a to b. In addition to the yield stress, strain rate change test was performed on the present in situ composite and Nb solid solution to make a better understanding of room temperature deformation behavior of the present alloys studied. If the strain rate sensitivity of the present alloys is similar to that of Nb solid solution, a room temperature deformation of the present alloy could be closely related to the deformation of bcc phase in the in situ composite. Fig. 6 shows the relationship between the yield stress and strain rate for the Nb–22at.%Si–5at.%Mo and Nb–22at.%Si–15at.%Mo in situ composites (alloy j and k in Fig. 1) and Nb–30at.%Mo alloy. Linear relationship is well defined for both alloys at room temperature.
Fig. 6. Strain rate dependence of yield stress at room temperature for Nb solid solution and Nb–22at.%Si–5at.%Mo and Nb–22at.%Si– 15at.%Mo in-situ composites. Note that the Mo content in the bcc solid solution of Nb–22at.%Si–15at.%Mo is about 27at.% which is very close to that of Nb–30at.%Mo tested.
Regarding the Nb–30at.%Mo alloy, yield stress increases with increasing strain rate. That is, a positive dependence of strain rate is observed. On the other hand, the yield stresses for the Nb-22at.%Si–5at.%Mo and Nb22at.%Si–15at.%Mo in situ composites are primarily insensitive to strain rate, suggesting a different deformation mechanism as compared to the Nb–30at.%Mo alloy. Fig. 7(a) and (b) shows the TEM microstructures on undeformed and deformed samples for the Nb–18at.% Si-15at.%Mo alloy heat treated at 2123 K for 48 h. The microstructures are found to consist of b-Nb5Si3 and Nb solid solution. Plastic strain (3%) was given at room temperature for the deformed sample. The microstructure of the undeformed sample is featureless in the Nb5Si3 phase but randomly distributed dislocations in the bcc phase are occasionally observed with a low density. These dislocations might be produced due to the difference of thermal expansion coefficient between co-existing phases. Therefore, the difference of thermal expansion coefficients of co-existing Nb5Si3 and bcc phase in Nb–18at.%Si–15at.%Mo in situ composite is estimated to be very small due to the low dislocation density in bcc phase. Concerning the microstructure of the deformed sample, mechanical twins with 20–30 nm
Fig. 7. TEM micrographs of Nb–18at.%Si–15at.%Mo heat treated at 2123 K for 24 h (a) and 3% plastically deformed at room temperature (b).
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in width rather than dislocations are frequently observed in the Nb5Si3 phase, as shown in Fig. 7(b). This microstructural observation is closely associated with the experimental result that yield stresses are insensitive to the changes of strain rate at room temperature. Thus, very weak strain rate dependence of yield stresses may be due to the twinning deformation at room temperature in the present alloy. Twinning deformation has been found in many complex-basis intermetallic alloys [6]. Much more works will be needed to understand the deformation mechanism in the two phase Nbss/Nb5Si3 in situ composites.
4. Summary and conclusions The phase stability, microstructure and room temperature mechanical properties were investigated for Nbss/Nb5Si3 in situ composites in the ternary Nb–Si–Mo alloy system. The composites consist of a-Nb5Si3 or bNb5Si3 phase, depending on heat treatment temperature and alloy composition, and Nb-rich bcc solid solution. The obtained results are summarized as follows. 1. Nb5Si3 phase is extended toward Mo5Si3 on the Nb5Si3–Mo5Si3 pseudo-binary line without having a solubility range of Si, suggesting the substitution of Mo for Nb site in Nb5Si3, while Nb phase forms the solid solution over Nb–40at.%Mo with Si solubility of 0.6 at.%. From the tie-line results in the partial ternary Nb–Si–Mo phase diagram, Mo is predominantly partitioned into the Nb solid solution rather than Nb5Si3 phase. 2. Lattice parameter of Nb solid solution decreases with increasing Mo content in the Nb solid solution and the c/a ratio of a-Nb5Si3 increases with increasing Mo content. 3. Phase transformation from a-Nb5Si3 to b-Nb5Si3 takes place at around 5 at.%Mo at 1973 K, indicating maximum solubility of Mo in the a-Nb5Si3. 4. Offset yield stress (0.2%) increases with increasing volume fraction of Nb5Si3 or increasing Mo content
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in the bcc solid solution. The composites consisting of b-Nb5Si3+bcc microstructures exhibit lower yield stresses than those consisting of aNb5Si3+bcc microstructures, suggesting that bNb5Si3 has deformability superior to a-Nb5Si3 at room temperature. 5. Nbss/Nb5Si3 in situ composites exhibit a very weak strain rate dependence of yield stress, while positive strain rate dependence of yield stress is observed in the Nb solid solution. 6. Mechanical twinning is observed in Nb5Si3 of Nbss/Nb5Si3 in situ composites deformed at room temperature. Acknowledgements This work is supported by a grant from the New Energy and Industrial Technology Development Organization (NEDO) of Japan. References [1] Shah DM, Anton DL, Pope DP. Mater Sci Eng A 1995;192193:658. [2] Davidson DL, Chan KS, Anton DL. Metal Trans A 1996;27:3292. [3] Bewlay BP, Lewandowski JJ, Jackson MR. JOM 1997;49(8):44. [4] Rigney JD, Lewandowsky JJ. Metal Trans A 1996;27:3007. [5] Pearson WB. A handbook of lattice spacings and structure of metals and alloys, vol. 2. Oxford: Pergamon, 1967. [6] Luzzi DE et al. Acta Materialia 1998;221:2913. [7] Chu F, Pope DP. Scripta Mertallurgica et Materialia 1993;28:331. [8] Kim W-Y, Luzzi DE, Pope DP. Acta Metall Mater, in press. [9] Marieb TN, Kaiser AD, Nutt SR, Anton DL, Shah DM. MRS Symp Proc 1991;213:329. [10] Chang CS, Pope DP. MRS Symp Proc 1993;288:477. [11] Smith LS, Aindow M, Loretto MH. MRS Symp Proc 1993;288:477. [12] Yoshimi K et al. Unpublished research. [13] Massalski TB et al. Metals Park, OH, ASM, 1986. [14] Westbrook JH. Intermetallic compounds. New York: John Wiley and Sons, 1967. [15] Chu F, Thoma DJ, McClellan K, Peralta P, He Y. Intermetallics 1999;7:611.
Microstructure and room temperature deformation of Nbss/Nb5Si3 in situ composites alloyed with Mo Won-Yong Kima,*, Hisao Tanakaa, Akio Kasamaa, Ryohei Tanakaa, Shuji Hanadab a
Japan Ultra-High Temperature Materials Research Institute(JUTEMI), 573-3 Okiube, Ube, Yamaguchi 755-0001, Japan b Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan Received 14 March 2001; received in revised form 13 April 2001; accepted 16 April 2001
Abstract Room temperature deformation behavior of Nb5Si3 based intermetallics in the ternary Nb–Si–Mo alloy system is investigated by compression testing in relation to microstructure and compositional effect. The partial ternary phase diagram of the Nb-Si-Mo alloy system containing phase equilibrium information at 1973 K is determined by metallography, X-ray diffraction and scanning electron microscopy (SEM) equipped with wavelength-dispersive X-ray fluorescence spectroscopy (WDS). A pseudo-binary compound of Nb5Si3–Mo5Si3 is confirmed to form at Nb5Si3-rich compositions by Mo addition, without having a solubility range of Si. The phase transformation from a-Nb5Si3 to b-Nb5Si3 is found to occur at about 5 at.% of Mo content on the Nb5Si3–Mo5Si3 pseudo-binary line. The observed yield stresses are largely dependent on not only the volume fraction and morphology of constituent phases but also the strength of Nb5Si3 phase equilibrating with Nb solid solution. The in situ composites consisting of aNb5Si3 phase and bcc solid solution exhibit higher yield stress than those consisting of b-Nb5Si3 phase and bcc solid solution when yield stress is compared at the same volume fraction of bcc phase. It is suggested that room temperature deformability of b-Nb5Si3 is superior to that of a-Nb5Si3 phase. Room temperature deformation behavior of the present composites will be discussed in relation to microstructural evolution. # 2001 Elsevier Science Ltd. All rights reserved. Keywords: A. Silicides, various; B. Phase transformations; B. Phase diagram; B. Twinning
1. Introduction Niobium silicide base alloys have a strong potential as high temperature structural materials due to their high melting point over 2273 K and excellent strength even at extremely high temperature above 1473 K [1–4]. Nb5Si3 displays a complex tetragonal D8m at high temperature and a D8l at low temperature representing the space group I4/mcm (140) for both crystal structures [5]. The lattice parameters of both phases are a=1.0018 nm and c=0.5072 nm for the stoichiometric high temperature phase (b-Nb5Si3, W5Si3-type), and a=0.6570 nm and c=1.1884 nm for the stoichiometric low temperature phase (a-Nb5Si3, Cr5B3-type), respectively. Owing to their complex crystal structures, plastic deformation of monolithic Nb5Si3 will be limited even in the high temperature range. An incorporation of a ductile phase to hard-deformable Nb5Si3 phase resulting in two-phase * Corresponding author. Fax: +81-836-51-1765. E-mail address: [email protected] (W.-Y. Kim).
silicide alloy will be one way to improve room temperature deformability. Concerning the role of bcc phase in the Nbss/Nb5Si3 in situ composites (where Nbss means niobium solid solution), it is expected that microcracks introduced at a production stage, or flaws formed due to thermal stress during cooling may be reduced or removed by incorporating the ductile bcc phase. Additionally, the ductile phase in the Nbss/ Nb5Si3 in situ composites may assist deformation of the hard phase by a role of strain accommodation if the alloys are plastically deformed at room temperature. In some Laves phase alloy systems [6–8], the brittle to ductile transition temperature (BDTT) of monolithic Laves phase is found to be higher than 1273 K like other complex-basis intermetallic compounds such as Mo5Si3, Nb3Al, Cr3Si and V3Si [9–12]. Nevertheless, it has been observed that the complex-basis intermetallics such as HfV2+Nb or Ta with two-phase microstructures consisting of bcc solid solution and Laves intermetallic phase deform by mechanical twinning at room temperature [6–8]. These reports were very
0966-9795/01/$ - see front matter # 2001 Elsevier Science Ltd. All rights reserved. PII: S0966-9795(01)00034-6
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attractive for alloy design of intermetallics, and hence most of studies have been concentrated on the incorporation of ductile phase into brittle intermetallic phase. The brittleness of intermetallics is still one of the most critical issues to develop intermetallics or intermetallic base materials as next generation high temperature materials. In the present study, we aim at correlating the room temperature deformation behavior with phase transformation of Nb5Si3 and solid solution hardening in Nb–Si–Mo alloys consisting of Nb5Si3 and bcc solid solution. Mo is chosen as the ternary alloying element in the Nbss/Nb5Si3 in situ composites to increase melting point, to change phase stability and to strengthen the composites by solid solution hardening.
2. Experimental procedure Various ternary Nb–Si–Mo alloys were prepared by arc-melting under an argon atmosphere in a watercooled copper hearth. The purities of raw materials were 99.9 wt.% of Nb, 99.99 wt.% of Si, 99.9 wt.% of Mo, respectively. The alloy buttons were re-melted at least 5 times to ensure chemical homogeneity. Homogenization heat treatment was carried out at 1973 K for 100 h in a high purity nitrogen atmosphere, followed by rapid cooling to room temperature at a rate of 200 K/ min in the high temperature region (1973 to 773 K) and then at a rate of 10 K/min in the low temperature region 773 K to room temperature). Some of the arc-melted samples were heat treated at 2123 K for 24 h in order to investigate the effect of Mo alloying on phase stability and mechanical property in the Nb–Si–Mo ternary alloy system. Samples for chemical composition analysis, microstructural observation and compression testing were made by an electro-discharge machine (EDM) from the heat treated samples. X-ray diffraction experiment
was conducted on the samples heat treated at 1973 and 2123 K to examine the crystal structures of constituent phases. The nominal compositions of alloys studied in the present study are shown in Table 1. Chemical compositions of the constituent phases are analyzed using a scanning electron microscope equipped with a wavelength-dispersive X-ray fluorescence spectroscope. Compression specimens with 2.52.5 mm cross-section and 6 mm height were mechanically polished using SiC paper and Al2O3 particles with water. Compression tests were carried out using an Instron model 8500 mechanical testing machine under a nitrogen atmosphere at room temperature and at an initial strain rate of 310 4 s 1. To observe the deformation microstructure by transmission electron microscopy (TEM), some of the specimens were deformed to 3% plastic strain, and then unloaded. Slices for TEM specimens (1.5 mm) were prepared by EDM and mechanically polished to 100 mm thickness. Final thinning was accomplished using the twin jet electropolishing technique (Struers Tenupol) to make a perforation of samples in a chemical solution of 94 ml of methanol and 5 ml of H2SO4 and 1 ml of HF. TEM observation was made in a Jeol 4000EX microscope operated at 400 kV.
3. Results and discussion Partial ternary phase diagram of Nb–Si–Mo alloy system investigated at 1973 K is shown in Fig. 1. The dotted line in the two-phase region of the phase diagram indicates the phase boundary between a-Nb5Si3 phase and b-Nb5Si3 phase. An extension of Nb5Si3 phase toward Mo5Si3 on the pseudo-binary Nb5Si3–Mo5Si3 line is found without a deviation from the line, indicating that Mo as a ternary alloying element substitutes for Nb rather than Si. From the determined tie-lines in the
Table 1 The nominal compositions (at.%) of 14 kinds of alloys and corresponding microstructural characteristics in Nb–Si–Mo system Sample no.
Nb
Si
Mo
Constituent phases
Volume fraction of Nb5Si3 (%)
Heat treatment
a b c d e f g h i j k l
79 69 59 77 67 57 75 65 55 73 63 53
16 16 16 18 18 18 20 20 20 22 22 22
5 15 25 5 15 25 5 15 25 5 15 25
a-Nb5Si3 a-Nb5Si3 b-Nb5Si3 a-Nb5Si3 a-Nb5Si3 b-Nb5Si3 a-Nb5Si3 b-Nb5Si3 b-Nb5Si3 a-Nb5Si3 b-Nb5Si3 b-Nb5Si3
and Nbss and Nbss and Nbss and Nbss and Nbss and Nbss and Nbss and Nbss and Nbss and Nbss and Nbss and Nbss
38 38 38 45 45 45 52 52 52 57 57 57
1973 K for 1973 K for 1973 K for 1973 K for 1973 K for 1973 K for 1973 K for 1973 K for 1973 K for 1973 K for 1973 K for 1973 K for
e1 h1
67 65
18 20
15 15
b-Nb5Si3 and Nbss b-Nb5Si3 and Nbss.
45 52
2123 K for 48 h 2123 K for 48 h
100 h 100 h 100 h 100 h 100 h 100 h 100 h 100 h 100 h 100 h 100 h 100 h
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Fig. 1. Partial Nb–Si–Mo phase diagram at 1973 K in the present study. The open circles represent the alloy compositions for alloys in the duplex phase field consisting of Nb5Si3 phase and Nb solid solution; the solid circles correspond to maximum solubility in each phase.
phase diagram, it is suggested that Mo is more predominantly partitioned into Nb solid solution than Nb5Si3 phase in the two-phase Nbss/ Nb5Si3 region. Nb solid solution is extended along the Nb–Mo binary phase line with a small amount of Si solubility less than 1 at.%. The solubility of Si in Nb solid solution is not so sensitive to Mo content. These results are very consistent with the fact that the miscible solid solution is formed between Nb and Mo, and the solubilities of Si in Nb and Mo solid solution are very small in a similar manner to both Nb–Si and Mo–Si binary alloy systems. The Si content of a ternary eutectic consisting of Nbss and Nb5Si3 increases slightly compared to that of binary Nb–Si alloy system. Back scattered electron images (BEI) of Nb–18at.%Si–15at.%Mo and Nb–20at.%Si– 15at.%Mo are shown in Fig. 2. The lightly contrasted phase is Nb solid solution, while the darkly contrasted phase is Nb5Si3 phase, as indicated by arrows in Fig. 2. Clearly, it is observed that hypo-eutectic is changed to hyper-eutectic microstructure with increasing Si content, as seen in Fig. 2(a) and (c). The Si content of the eutectic is estimated from the microstructures to be between 18 and 20 at.% within the composition range investigated. With increasing Mo content at constant Si content, the volume fraction of primary bcc solid solution increased, the size of bcc particles increased and bcc volume fraction in a eutectic decreased. However, there was no change in volume fraction of the constituent phases at constant Si content, indicating that the volume fraction of the phases depends mostly upon Si content. The results of X-ray diffraction on several samples heat treated at both 1973 and 2123 K are shown in Fig. 3. The open squares in the figure indicate peaks due to a-Nb5Si3, and the closed circles indicate peaks from
Fig. 2. Back-scattered electron images of Nb–18at.%Si–15at.%Mo heat treated at 1973 K (a) and at 2123 K (b), and Nb–20at.%Si– 15at.%Mo alloy heat treated at 1973 K (c). All the microstructures consist of Nb5Si3 and bcc solid solution. Note that a microstructural transition from hypo-eutectic to hyper-eutectic occurs at a composition between (a) and (c).
the bcc solid solution. The b-Nb5Si3 phase is marked by the closed squares. While most diffraction peaks are common to the two Nb5Si3 phases, the 112 and 202 peaks of the a-Nb5Si3 phase and the 211 and 310 peaks of the bNb5Si3 phase allow their presence to be uniquely determined. We could not find an intermediate Nb3Si phase in the present alloys whatever heat treatment was performed at 1973 or 2123 K. Therefore, Mo addition to the
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Nb5Si3/Nbss in situ composite makes Nb3Si unstable. In the case of samples heat treated at 1973 K, it is found that the stability of the Nb5Si3 phases is changed depending on Si and Mo contents. b-Nb5Si3 is stable in the high Si and Mo region, while a-Nb5Si3 was stable in the low Si and Mo region. In the case of samples heat treated at 2123 K, b-Nb5Si3 is found to be stable over the whole composition range examined in the present study. In the binary Nb–Si phase diagram, Nb5Si3 phase undergoes a phase transformation from a to b depending on temperature and composition. Further the crystal structure of Mo5Si3 forming the pseudo-binary phase with Nb5Si3 is the same as b phase. Therefore, as seen in the alloys h and k of Fig. 1, the accelerated phase transformation from a to b phase at 1973 K by addition of Mo may be associated with a change of phase stability, since a-Nb5Si3 phase is stable at 1973 K in the binary Nb–Si phase diagram [13]. Assuming that the phase transformation from a to b in the present Nb–Si–Mo ternary alloy system involves a diffusion process, transformation kinetics should be considered to explain the alloying effect. Since Mo has a higher melting point than Nb or Si, it can not be explained by enhancement of the diffusion process due to Mo addition. In Fig. 4(a), the lattice parameters of the bcc phases of the several alloys are plotted as a function of Mo content in the bcc phase in the two-phase alloys. A monotonic decrease in lattice parameter is seen,
although it appears not to be linear. From the atomic sizes of Nb, Mo and Si being 0.147, 0.140 and 0.116 nm, respectively, it is clear that the largest Nb atom is replaced by Mo with increasing Mo content. In Fig. 4(b), the c/a ratio in the unit cell of a-Nb5Si3 phase is plotted as a function of Mo content in the a-Nb5Si3 phase. Since the atomic size of Mo is smaller than Nb and the pseudo-binary line extends without a deviation of Si content from the stoichiometric composition, Mo would substitute for Nb sites, thereby leading to smaller crystal size or reducing c/a ratio of tetragonal Nb5Si3 phase. However the calculated c/a ratio is somewhat different from our expectation. It is found that with increasing Mo content, a-axis is shortened, whereas caxis is lengthened, resulting in an increase of c/a ratio in the a-Nb5Si3 phase. At the present time, the propensity of c- and a-axial length change is puzzling, but probably the anisotropy of inter-atomic bonding force between either unlike or like atoms is responsible for the substitution behavior of ternary alloying elements in this alloy system [14–15]. Two kinds of chain structures, –Si– Mo–Si– and Mo–Mo or –Si–Si–, were reported in Mo5Si3 having the same crystal structure (body-centered tetragonal, tI32) as Nb5Si3. Chu et al. [15] have demonstrated that the strong anisotropy of thermal expansion and hardness arises from the bonding along [100] direction corresponding to –Si–Mo–Si– chain structure and [001] direction corresponding to –Mo–Mo– or –Si–Si– chain
Fig. 3. X-ray powder diffraction spectra for 6 alloys at a constant Mo content of approximately 15 at.%Mo. Alloys b, e, h and k are heat treated at 1973 K, and alloys e1 and h1 are heat treated at 2123 K. The chemical compositions of alloys el and h1 are the same as alloy e and alloy h, respectively.
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structure. Therefore, the selective substitution behavior of Mo for Nb sites may lead to an increase in c/a ratio, and reduce the stability of a-Nb5Si3 in the present Nb– Si–Mo alloy system. Further investigation is required to elucidate the atomic substitution behavior and associated physical and mechanical properties. Compression tests were carried out on samples heat treated at 1973 K for 100 h and 2123 K for 24 h for both alloys with two-phase microstructures consisting of a-Nb5Si3 or b-Nb5Si3 and bcc solid solution at room temperature and a strain rate of 310 4 s 1. A large number of alloys with compositions in the range from the hypo-eutectic and hyper-eutectic were tested. Within this set of alloys (alloy b, e, h and k), two systematic variations were employed to ascertain the effect of silicide phases
equilibrating with bcc solid solution on mechanical properties, and the effect of volume fraction of each phase. Offset yield stresses (0.2%) are plotted as a function of volume fraction of Nb5Si3 phase, as shown in Fig. 5(a). Concerning the samples heat treated at 2123 K, yield stress increases with increasing volume fraction of b-Nb5Si3 phase. Correspondingly compressive ductility decreases with increasing volume fraction of b-Nb5Si3. A sharp increase in yield stress is seen at 52% of b-Nb5Si3 phase. This result may be associated with microstructural transition from hypo-eutectic (primary bcc+eutectic) to hyper-eutectic (primary b-Nb5Si3+eutectic). In the alloy heat treated at 1973 K, it is found that yield stress decreases with increasing volume fraction of Nb5Si3 phase from 45% (alloy e) to 52% (alloy h).
Fig. 4. Lattice parameters of Nb solid solutions equilibrating with Nb5Si3 phase as a function of Mo content (a), and the c/a ratio of aNb5Si3 versus Mo content at a constant Si content (b). All samples were heat treated at 1973 K for 100 h.
Fig. 5. Yield stress (a) and compressive ductility (b) of alloys tested at room temperature as a function of total volume fraction of b-Nb5Si3 or b-Nb5Si3 phase in each alloy.
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Furthermore, it was confirmed for all alloys studied here that Mo content in the bcc solid solution equilibrating with Nb5Si3 phase increases with increasing volume fraction of Nb5Si3 phases whatever they exist as a or b. Therefore, the yield stress drop at 52% of Nb5Si3 is closely correlated with phase transformation of aNb5Si3 to b-Nb5Si3, suggesting that b-Nb5Si3 deforms at a lower stress than a-Nb5Si3 at room temperature in this Nb–Si–Mo alloy system. It is interesting to note that yield stresses display similar values for the two-phase Nbss/b-Nb5Si3 alloys when the volume fraction of bNb5Si3 phase is not changed significantly by the heat treatment at 1973 K (alloy h) or 2123 K (alloy h1), see yield stresses at 52% (alloy h and h1) volume fraction of Nb5Si3. Thus, the large increase in compressive ductility at 52% of b-Nb5Si3 for the heat treatment at 1973 K, as shown in Fig. 5(b), is considered to be a strong evidence of phase transformation of Nb5Si3 from a to b. In addition to the yield stress, strain rate change test was performed on the present in situ composite and Nb solid solution to make a better understanding of room temperature deformation behavior of the present alloys studied. If the strain rate sensitivity of the present alloys is similar to that of Nb solid solution, a room temperature deformation of the present alloy could be closely related to the deformation of bcc phase in the in situ composite. Fig. 6 shows the relationship between the yield stress and strain rate for the Nb–22at.%Si–5at.%Mo and Nb–22at.%Si–15at.%Mo in situ composites (alloy j and k in Fig. 1) and Nb–30at.%Mo alloy. Linear relationship is well defined for both alloys at room temperature.
Fig. 6. Strain rate dependence of yield stress at room temperature for Nb solid solution and Nb–22at.%Si–5at.%Mo and Nb–22at.%Si– 15at.%Mo in-situ composites. Note that the Mo content in the bcc solid solution of Nb–22at.%Si–15at.%Mo is about 27at.% which is very close to that of Nb–30at.%Mo tested.
Regarding the Nb–30at.%Mo alloy, yield stress increases with increasing strain rate. That is, a positive dependence of strain rate is observed. On the other hand, the yield stresses for the Nb-22at.%Si–5at.%Mo and Nb22at.%Si–15at.%Mo in situ composites are primarily insensitive to strain rate, suggesting a different deformation mechanism as compared to the Nb–30at.%Mo alloy. Fig. 7(a) and (b) shows the TEM microstructures on undeformed and deformed samples for the Nb–18at.% Si-15at.%Mo alloy heat treated at 2123 K for 48 h. The microstructures are found to consist of b-Nb5Si3 and Nb solid solution. Plastic strain (3%) was given at room temperature for the deformed sample. The microstructure of the undeformed sample is featureless in the Nb5Si3 phase but randomly distributed dislocations in the bcc phase are occasionally observed with a low density. These dislocations might be produced due to the difference of thermal expansion coefficient between co-existing phases. Therefore, the difference of thermal expansion coefficients of co-existing Nb5Si3 and bcc phase in Nb–18at.%Si–15at.%Mo in situ composite is estimated to be very small due to the low dislocation density in bcc phase. Concerning the microstructure of the deformed sample, mechanical twins with 20–30 nm
Fig. 7. TEM micrographs of Nb–18at.%Si–15at.%Mo heat treated at 2123 K for 24 h (a) and 3% plastically deformed at room temperature (b).
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in width rather than dislocations are frequently observed in the Nb5Si3 phase, as shown in Fig. 7(b). This microstructural observation is closely associated with the experimental result that yield stresses are insensitive to the changes of strain rate at room temperature. Thus, very weak strain rate dependence of yield stresses may be due to the twinning deformation at room temperature in the present alloy. Twinning deformation has been found in many complex-basis intermetallic alloys [6]. Much more works will be needed to understand the deformation mechanism in the two phase Nbss/Nb5Si3 in situ composites.
4. Summary and conclusions The phase stability, microstructure and room temperature mechanical properties were investigated for Nbss/Nb5Si3 in situ composites in the ternary Nb–Si–Mo alloy system. The composites consist of a-Nb5Si3 or bNb5Si3 phase, depending on heat treatment temperature and alloy composition, and Nb-rich bcc solid solution. The obtained results are summarized as follows. 1. Nb5Si3 phase is extended toward Mo5Si3 on the Nb5Si3–Mo5Si3 pseudo-binary line without having a solubility range of Si, suggesting the substitution of Mo for Nb site in Nb5Si3, while Nb phase forms the solid solution over Nb–40at.%Mo with Si solubility of 0.6 at.%. From the tie-line results in the partial ternary Nb–Si–Mo phase diagram, Mo is predominantly partitioned into the Nb solid solution rather than Nb5Si3 phase. 2. Lattice parameter of Nb solid solution decreases with increasing Mo content in the Nb solid solution and the c/a ratio of a-Nb5Si3 increases with increasing Mo content. 3. Phase transformation from a-Nb5Si3 to b-Nb5Si3 takes place at around 5 at.%Mo at 1973 K, indicating maximum solubility of Mo in the a-Nb5Si3. 4. Offset yield stress (0.2%) increases with increasing volume fraction of Nb5Si3 or increasing Mo content
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in the bcc solid solution. The composites consisting of b-Nb5Si3+bcc microstructures exhibit lower yield stresses than those consisting of aNb5Si3+bcc microstructures, suggesting that bNb5Si3 has deformability superior to a-Nb5Si3 at room temperature. 5. Nbss/Nb5Si3 in situ composites exhibit a very weak strain rate dependence of yield stress, while positive strain rate dependence of yield stress is observed in the Nb solid solution. 6. Mechanical twinning is observed in Nb5Si3 of Nbss/Nb5Si3 in situ composites deformed at room temperature. Acknowledgements This work is supported by a grant from the New Energy and Industrial Technology Development Organization (NEDO) of Japan. References [1] Shah DM, Anton DL, Pope DP. Mater Sci Eng A 1995;192193:658. [2] Davidson DL, Chan KS, Anton DL. Metal Trans A 1996;27:3292. [3] Bewlay BP, Lewandowski JJ, Jackson MR. JOM 1997;49(8):44. [4] Rigney JD, Lewandowsky JJ. Metal Trans A 1996;27:3007. [5] Pearson WB. A handbook of lattice spacings and structure of metals and alloys, vol. 2. Oxford: Pergamon, 1967. [6] Luzzi DE et al. Acta Materialia 1998;221:2913. [7] Chu F, Pope DP. Scripta Mertallurgica et Materialia 1993;28:331. [8] Kim W-Y, Luzzi DE, Pope DP. Acta Metall Mater, in press. [9] Marieb TN, Kaiser AD, Nutt SR, Anton DL, Shah DM. MRS Symp Proc 1991;213:329. [10] Chang CS, Pope DP. MRS Symp Proc 1993;288:477. [11] Smith LS, Aindow M, Loretto MH. MRS Symp Proc 1993;288:477. [12] Yoshimi K et al. Unpublished research. [13] Massalski TB et al. Metals Park, OH, ASM, 1986. [14] Westbrook JH. Intermetallic compounds. New York: John Wiley and Sons, 1967. [15] Chu F, Thoma DJ, McClellan K, Peralta P, He Y. Intermetallics 1999;7:611.