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The effect of second-phase dispersions on mechanical property of Ni3Si based multi-phase intermetallic alloys
Materials Science and Engineering A 476 (2008) 112–119
The effect of second-phase dispersions on mechanical property of Ni3Si based multi-phase intermetallic alloys M. Fujita a , Y. Kaneno a , T. Takasugi a,b,∗ a
Department of Materials Science, Graduate School of Engineering, Osaka Prefecture University, 1-1 Gakuen-cho, Naka-ku, Sakai, Osaka 599-8531, Japan b Osaka Center for Industrial Materials Research, Institute for Materials Research, Tohoku University, 1-1 Gakuen-cho, Naka-ku, Sakai, Osaka 599-8531, Japan Received 10 November 2006; received in revised form 17 April 2007; accepted 20 April 2007
Abstract Using alloys whose initial microstructures are composed of Ni3 Si(L12 ), Ni3 Si(L12 ) + Ni3 Ti(D024 ) and Ni3 Si(L12 ) + Ni3 Nb(D0a ), aging phenomenon and the associated high-temperature tensile property were investigated. The plate-like Ni3 Ti (D024 ) phase was precipitated from the Ni3 Si (L12 ) matrix at high temperature. It was shown by micro-hardness measurement that age hardening behavior due to the precipitation of the Ni3 Ti (D024 ) phase occurs in all alloys in the temperature range from 923 K to 1323 K. It was however shown by tensile test that the precipitated Ni3 Ti (D024 ) phase is not so much effective in improving the mechanical properties of alloys whose initial microstructures are composed of Ni3 Si(L12 ) + Ni3 Nb(D0a ) or Ni3 Si(L12 ) + Ni3 Ti(D024 ). In alloys whose initial microstructures are composed of Ni3 Si(L12 ) + Ni3 Nb(D0a ), a good combination of tensile strength and tensile elongation was found over a wide of test temperature regardless of presence or absence of the precipitated Ni3 Ti (D024 ) phase. It is suggested that the primary coarse Ni3 Nb (D0a ) phase precipitates are beneficial to improving the mechanical properties. © 2007 Elsevier B.V. All rights reserved. Keywords: Multi-phase intermetallics; Ni3 Si; Ni3 Ti; Ni3 Nb; GCP phase; Microstructure; Mechanical property
1. Introduction Geometrically close packed (GCP) Ni3 X phases combined with IIIb and IVb non-transition elements (e.g., Al and Si) [1–6], and also IVa and Va transition metals (e.g., Ti, V and Nb) [7–10] generally exhibit high thermal, chemical and microstructural stabilities, and also attractive mechanical properties such as strength anomaly at high temperatures [5–10], and therefore have been of great interest as a strengthener in Ni-based superalloys. However, the GCP Ni3 X intermetallic phases as well as many other intermetallic compounds show some drawbacks, such as poor ductility at room temperature and inferior creep strength at high temperature. It was recently shown that multi-phase intermetallic alloys, e.g., composed of GCP Ni3 X (X: Al, Ti and Nb) [11,12], (X: Si, Ti and Nb) [13–15], (X: Al, Ti and V) [16–18] and
∗ Corresponding author at: Department of Materials Science, Graduate School of Engineering, Osaka Prefecture University, 1-1 Gakuen-cho, Naka-ku, Sakai, Osaka 599-8531, Japan. Tel.: +81 72 254 9314; fax: +81 72 254 9912. E-mail address: [email protected] (T. Takasugi).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.04.085
(X: Al, Nb and V) [19] intermetallic phases are thermodynamically stable and exhibit stable microstructures at high temperatures. Also, it was expected that such GCP Ni3 Xbased multi-phase intermetallic alloys have coherent interfaces among their constituents because of their closely related crystal structures. With regard to the multi-phase Ni3 Si–Ni3 Ti–Ni3 Nb intermetallic alloys, an isothermal phase diagram constructed at 1323 K showed that there were a three-phase region consisting of Ni3 Si(L12 )–Ni3 Ti(D024 )–Ni3 Nb(D0a ), and twophase regions consisting of Ni3 Si(L12 )–Ni3 Ti(D024 ) and Ni3 Ti(D024 )–Ni3 Nb(D0a ) [13–15]. It was also found that the Ni3 Ti(D024 ) phase is precipitated from the Ni3 Si(L12 ) phase matrix by a solid-state reaction [15]. It was shown by hightemperature tensile test that alloys in which the Ni3 Si(L12 ) matrix was precipitated by the Ni3 Ti(D024 ) and Ni3 Nb(D0a ) phases had the most favorable properties as high-temperature mechanical and chemical materials although the majority of these multi-phase intermetallic alloys resulted in high tensile strength accompanied with certain levels of tensile ductility over a wide range of temperatures [14]. In this study, the effect of second-phase dispersions on mechanical properties of multi-phase intermetallic alloys
M. Fujita et al. / Materials Science and Engineering A 476 (2008) 112–119 Table 1 Alloy compositions used in this study Alloy
No. 1 No. 2 No. 3 No. 4
Alloy compositions Ni (at.%)
Si (at.%)
Ti (at.%)
Nb (at.%)
B (wt ppm)
79.5 79.5 79.5 79.5
8.2 6.2 8.2 6.15
9.2 11.3 8.2 8.2
3.1 3.0 4.1 6.15
50 50 50 50
whose initial microstructures are composed of Ni3 Si(L12 ), Ni3 Si(L12 ) + Ni3 Ti(D024 ) and Ni3 Si(L12 ) + Ni3 Nb(D0a ) is investigated. First, the changes of microstructure and hardness due to the Ni3 Ti(D024 ) phase precipitation by aging are investigated. Then, high-temperature tensile properties of these multi-phase intermetallic alloys are evaluated in the unaged and aged conditions. 2. Experimental procedures Alloys used in this study were prepared from starting raw materials of 99.9 wt.% Ni, 99.999 wt.% Si, 99.9 wt.% Ti and 99.9 wt.% Nb. All alloys were doped with 50 ppm boron. Alloys were prepared as button ingots with a 50 mm diameter by arc melting under an argon gas atmosphere on a copper hearth using a tungsten electrode. Alloy compositions of the button ingots are shown in Table 1 and also plotted in Fig. 1, i.e., in the form of the Ni3 Si–Ni3 Ti–Ni3 Nb pseudo-ternary phase diagram reported previously [15]. An identical 79.5 at.% Ni content was kept for the prepared alloys. All button ingots were homogenized at 1323 K for 48 h in a vacuum, followed by furnace cooling to room temperature. The button ingots were sectioned to a proper size mostly along a solidified direction of the button ingot by an electro-discharge machine (EDM). The sectioned specimens were evacuated into a silica tube and then annealed at 1373 K for 24 h in a vacuum by which the specimens were solution-treated,
Fig. 1. Isothermal Ni3 Si–Ni3 Ti–Ni3 Nb pseudo-ternary phase diagram at 1323 K. Alloy compositions used in this study are plotted in this figure.
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and then water-quenched. For aging, these specimens were reevacuated into a silica tube and then annealed at temperatures ranging from 823 K to 1323 K for various times, followed by water quenching. Metallographic and structural observations for these heattreated specimens were conducted by optical microscopy (OM), X-ray diffraction (XRD) and scanning electron microscopy (SEM). X-ray diffraction profiles were obtained using Cu K␣ radiation. The determination of the constituent intermetallic phases was conducted on the basis of SEM electron-probe analysis (EPMA) combined with XRD analysis. Mechanical properties were evaluated by means of microhardness test and tensile test. For Vickers hardness measurement, a load of 300 g was used. Tensile specimens with a gauge dimension 2 mm × 1 mm × 10 mm were cut from the button ingots homogenized at 1323 K for 48 h. Tensile test was conducted using unaged specimens as well as aged specimens: the former specimens were annealed at 1373 K for 24 h, while the latter specimens were annealed at 1373 K for 24 h and subsequently aged at 1273 K for 12 h, by which almost maximum hardening is obtained. The surface of the tensile specimens was abraded with a fine SiC paper. The tensile test in the temperature range between room temperature and 1173 K was conducted in a vacuum (∼1.5 × 10−3 Pa) within a metal tube surrounded by an electric furnace. A nominal strain rate used in the tensile test was 1.67 × 10−4 s−1 . Tensile yield strength was determined at 0.2% offset strain. Tensile elongation was calculated from the load–displacement curve and defined as the plastic strain to fracture. The fracture surface of the tensile specimens was examined by a scanning electron microscope. 3. Results 3.1. Microstructures Fig. 2 shows initial microstructures of alloys solution-treated at 1373 K for 48 h. Alloy 1 exhibits almost L12 single-phase microstructure although a very small volume fraction (∼1%) of the Ni3 Nb (D0a ) dispersions are contained. Alloy 2 exhibits a two-phase microstructure consisting of the Ni3 Si (L12 ) and Ni3 Ti (D024 ) phases. The Ni3 Ti (D024 ) phase with a volume fraction of ∼28% exhibits coarse plate-like morphology within the L12 phase matrix. On the other hand, alloys 3 and 4 exhibit microstructures composed of the Ni3 Si (L12 ) phase matrix and the Ni3 Nb (D0a ) phase particles. The volume fraction of the Ni3 Nb (D0a ) phase dispersions was larger in alloy 4 (∼14%) than in alloy 3 (∼3.5%). Thus, the feature of the microstructures observed at 1373 K (Fig. 2) was basically consistent with the phase diagram (at 1323 K) shown in Fig. 1. However, a minor inconsistency with the phase diagram may be due to the change of the solubility limits of each constituent phase (i.e., Ni3 Si, Ni3 Ti and Ni3 Nb) by temperature. For an example, the solubility limit of the Ni3 Si(L12 ) phase against the Ni3 Ti(D024 ) phase expands with increasing temperature. Also, it is deduced that the Ni3 Ti (D024 ) phases observed in alloy 2 (Fig. 2(b)) as well as the Ni3 Nb (D0a ) phases observed in alloys 3 (Fig. 2(c)) and
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M. Fujita et al. / Materials Science and Engineering A 476 (2008) 112–119
Fig. 2. Back scattering (BS)-SEM microstructures of alloys: (a) 1, (b) 2, (c) 3 and (d) 4 solution-treated at 1373 K for 48 h.
4 (Fig. 2(d)) were formed directly from liquid phase because both phase dispersions are large and composed of dendritic morphology. In all alloys, the precipitation of the Ni3 Ti (D024 ) phase occurred at temperatures above 823 K. Fig. 3 representatively shows microstructures of alloys 1–4 aged at a relatively high temperature 1223 K, indicating the microstructural changes during aging. The plate-like (or needle-like) Ni3 Ti (D024 ) phase precipitates appeared in the L12 phase matrix and the volume fraction and size of the precipitates tended to increase with increasing aging time for all alloys. Also, Fig. 3 indicates that the interfacial plane between the L12 and D024 phases are quite straight and coincides with the certain crystallographic planes. The previous study [15] suggested that the orientation relationship between both phases is expressed by; [0 1 1]L12 //[1 1 2¯ 0]D024
(1a)
(1¯ 1 1)L12 //(0 0 0 1)D024
(1b)
The closest packed atomic planes and directions in the L12 and D024 phases are parallel each other. For alloy 2, the aging resulted in a bimodal microstructure for the Ni3 Ti (D024 ) phase dispersions, i.e., (small precipitates introduced by aging) + (prior large dispersions formed during solidification). For alloys 3 and 4, the aging resulted in a three-phase microstructure in which the Ni3 Ti (D024 ) and Ni3 Nb (D0a ) phases are dispersed within the Ni3 Si (L12 ) phase matrix, being consequently consistent with the microstructure expected from the phase diagram (Fig. 1). Also, the volume fraction of the Ni3 Ti (D024 ) phase precipitates in alloy 4 was larger than that in alloy 3, again consistent with the phase diagram (Fig. 1). Thus, favorable microstructures in which fine Ni3 Ti (D024 ) phase precipitates
densely distribute in the Ni3 Si (L12 ) phase matrix were obtained by aging. 3.2. Age hardening behavior by hardness measurement Fig. 4 shows the variations of Vickers hardness with aging time for alloys 1, 2, 3 and 4. Here, it should be noted that the Vickers hardness was measured within the Ni3 Si (L12 ) phase matrix, avoiding coarse primary Ni3 Ti (D024 ) and Ni3 Nb phase (D0a ) particles existing in the initial microstructures. In unaged (i.e., as-solution-treated) condition, alloys showed hardness values of 370–390 Hv. For all alloys, age hardening occurred at temperatures above 823 K by precipitation of the Ni3 Ti (D024 ) phase, as understood from Fig. 3. As aging temperature increases, hardening occurred more quickly and then reached a maximum hardness at shorter aging time. For alloys 1, 3 and 4, the maximum hardness with approximately 460 Hv was observed by aging at 1273 K for 12 h while for alloy 2, the maximum hardness with approximately 445 Hv was observed by aging at 1173 K for 24 h. However, apparent over-age hardening, i.e., apparent softening did not occur even at the maximum temperature 1273 K and also even at the longest aging time more than 168 h (although their data were omitted from Fig. 4). This result indicates that the aged microstructures are fairly stable for prolonged annealing time at high temperature. In Fig. 4, the changes in Vickers hardness shows complicated behavior with aging temperature. That is, the hardness does not show monotonic increase with aging temperature (especially in alloy 3). At the moment, the explanation is not clearly presented, but may be associated with some changes in microstructure, according to an unusual increase of the over-saturation of solutes in L12 phase against to D024 phase with increasing temperature.
M. Fujita et al. / Materials Science and Engineering A 476 (2008) 112–119
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Fig. 3. Back scattered (BS)-SEM microstructures of alloys: (a) 1, (b) 2, (c) 3 and (d) 4 aged at 1223 K for 12 h and 72 h, respectively.
3.3. Tensile behavior Fig. 5 shows changes of stress versus strain curves with temperature for all alloys tested in this study. Depending on tem-
perature and alloy, largely different stress versus strain curves were observed. Alloy 1 deformed at low temperature showed a relatively large strain-hardening rate, and then ruptured. However, alloy 1 did not show an apparent plastic deformability at
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M. Fujita et al. / Materials Science and Engineering A 476 (2008) 112–119
Fig. 4. Changes of Vickers hardness by aging time for alloys: (a) 1, (b) 2, (c) 3 and (d) 4 aged at various temperatures.
temperature beyond 1100 K. For alloy 2, there was little plastic deformability over a whole test temperature. Like alloys 1, 3 and 4 deformed at low temperature showed a relatively large strain-hardening rate and then ruptured, leading to high plastic deformability. With increasing temperature, the strainhardening rate decreased. At a peak temperature (i.e., ∼900 K) where the yield stress shows the highest value, the strainhardening rate was still positive. Alloys 3 and 4 deformed at high temperature (beyond 900 K) exhibited a steady-state flow
after yielding, consequently resulting in high tensile elongation. From over-all evaluation for the tensile property, it is deduced that alloy 3 has a good combination of tensile stress and tensile elongation over a wide of test temperature while alloy 2 shows the worst result not only for strength property but also for tensile elongation property over a wide of test temperature. Figs. 6–9 plot yield stress, ultimate tensile stress (UTS) and tensile elongation as a function of temperature for alloys 1, 2, 3 and 4, respectively. For each alloy, the measurement was
Fig. 5. Stress vs. strain curves of unaged and aged alloys: (a) 1, (b) 2, (c) 3 and (d) 4 deformed at various temperatures, respectively.
M. Fujita et al. / Materials Science and Engineering A 476 (2008) 112–119
Fig. 6. Yield stress, tensile stress and tensile elongation as a function of test temperature for unaged and aged alloy 1, respectively. In aged alloy 1, the data for the yield stress at high temperatures are not plotted because the specimens fractured in the elastic region.
Fig. 7. Yield stress, tensile stress and tensile elongation as a function of test temperature for unaged and aged alloy 2, respectively. In alloy 2, the data for the yield stress at high temperatures are not plotted because the specimens fractured in the elastic region.
performed on unaged condition (as solution-treated at 1373 K for 24 h) and subsequently aged (at 1273 K for 12 h) condition, respectively. The yield stress increases with increasing temperature, and makes a broad maximum at around 873 K, followed by a decrease at high temperature. The reduced yield strength over a wide range of temperature was observed in alloy 2. Similar to the yield strength behavior, ultimate tensile stress increases with increasing temperature, and makes an apparent maximum
Fig. 8. Yield stress, tensile stress and tensile elongation as a function of test temperature for unaged and aged alloy 3, respectively.
117
Fig. 9. Yield stress, tensile stress and tensile elongation as a function of test temperature for unaged and aged alloy 4, respectively.
at around 673–873 K, followed by a decrease at high temperature regardless of alloys. The reduced ultimate tensile strength over a wide range of temperature was again observed in alloy 2. Regarding tensile elongation, all alloys more or less exhibit certain levels of tensile elongation over a wide rage of temperature. The tensile elongation shows a peak at an intermediate temperature (∼673 K), and then shows a minimum at ∼1073 K, followed by a steep increase at high temperature, except for alloys 1 and 2. Alloy 2 consistently suffers from low tensile elongation over an entire test temperature while alloy 1 shows high tensile elongation at low temperature but low tensile elongation at high temperature. For aging effect, the increased yield strength and the reduced tensile elongation were observed for alloy 1. However, for the other alloys, aging effect was not evident not only on the flow strength but also on the tensile elongation although there was an apparent difference in the Vickers hardness values between the unaged and aged conditions. Also, it is noted that steep increases of tensile elongation take place in alloys 3 and 4 deformed at the highest temperature (1173 K). The fracture patterns of aged alloys 1–4 are shown in Fig. 10. The fracture patterns of unaged alloys were primarily similar to those of aged alloys at a whole test temperature and therefore omitted from this figure. The observed fracture patterns are basically consistent with the tensile ductility shown in Figs. 5–9. In other words, alloys with high tensile elongation showed ductile transgranular fracture mode with dimple-like patterns while alloys with low tensile elongation showed brittle fracture mode with intergranular patterns or large facetted planes. For alloy 1 that showed high tensile elongation at low temperature and low tensile elongation at high temperature, dimple-like fracture patterns accompanied with large facetted planes were observed at low temperature and intergranular fracture patterns were observed at high temperature (Fig. 10(a)). The facetted planes may correspond to the interfacial planes between the L12 phase matrix and the D024 phase precipitate. For alloy 2 with the lowest tensile elongation among four alloys, largely facetted fracture patterns were observed at each test temperature. It is likely that largely facetted fracture planes correspond to the interfacial plane between the coarse primary D024 phase plate and the L12 phase matrix (see Fig. 2(b)). For alloys 3 and 4 that showed relatively large tensile elongation at a whole test temperature, quite ductile fracture mode, i.e., transgranular
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Fig. 10. SEM fractography of aged alloys: (a) 1, (b) 2, (c) 3 and (d) 4 deformed at RT, 673 K and 1173 K, respectively.
fracture mode with dimple-like patterns was dominated. Also, the evidence that dynamic recrystallization occurs was shown at the highest temperature (1173 K): the fractured surfaces were covered by a numerous number of (recrystallized) small grains. 4. Discussion For alloy 1 whose initial microstructure is composed of the Ni3 Si (L12 ) single-phase matrix, micro-hardness measurement certainly showed age hardening and tensile test showed slightly enhanced yield strength due to the precipitation of the Ni3 Ti (D024 ) phase. Ni3 Ti was precipitated in a plate or needle-like form in the matrix as shown in Figs. 2 and 3. Direct TEM observation and the misfit evaluated from lattice parameters of Ni3 Si (L12 ) and Ni3 Ti (D024 ) phases [15] reveal that Ni3 Ti precip-
itates are coherent or partially coherent to the matrix. For the strengthening mechanism by Ni3 Ti precipitates, (i) generation of interfacial dislocations during the dislocation propagation event, (ii) transformation of the dislocation core during intersection of the interface according to the change of the crystal structure, or (iii) intersection of the network of the misfit dislocations situated at the interface may be suggested. It is consequently difficult to expect a substantial strengthening effect from these mechanisms as far as the present issue is concerned with shortterm mechanical property. Conversely, if the present issue is concerned with the long-term mechanical property (e.g., creep), fine and coherent precipitates may effectively enhance the hightemperature mechanical properties (e.g., strength). In addition, the size and spacing of the precipitated Ni3 Ti phase appear to be too large to induce the precipitation strengthening like
M. Fujita et al. / Materials Science and Engineering A 476 (2008) 112–119
Orowan’s mechanism. When the other alloys 2–4 whose initial microstructures are composed of Ni3 Si(L12 ) + Ni3 Ti(D024 ) or Ni3 Si(L12 ) + Ni3 Nb(D0a ) are evaluated by tensile test, the age hardening due to the precipitation of the Ni3 Ti (D024 ) phase was little effective: strength and tensile ductility were little affected by the precipitation of the Ni3 Ti (D024 ) phase at a wide range of test temperature. In alloys 3 and 4, Ni3 Nb (D0a ) precipitates are incoherent with Ni3 Si (L12 ) matrix [14,15]. During deformation, it is likely that dislocations are emitted preferentially from the precipitates or contrariwise piled up to the precipitates, consequently inducing the strengthening and plastic deformability through the entire material [20,21]. Thus, the obtained results imply that Ni3 Nb(D0a ) dispersions are more effective than Ni3 Ti (D024 ) precipitates not only on strength property but also on tensile ductility of the present multi-phase intermetallic alloys. These results reveal that the strengthening due to the second-phase incoherent with the matrix is more effective than that due to the second-phase coherent or partially coherent with the matrix, and also that the second-phase incoherent with the matrix is beneficial, enhancing the tensile ductility of the matrix. Thus, it is deduced that over-all mechanical properties are better in alloys whose initial microstructure is composed of the L12 phase matrix dispersed by the Ni3 Nb (D0a ) phase precipitates rather than in alloys whose initial microstructure is composed of the L12 single phase matrix or the L12 phase matrix precipitated by the coarse primary Ni3 Ti (D024 ) or precipitated Ni3 Ti (D024 ). 5. Conclusions
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(3) It was found that the precipitated Ni3 Ti (D024 ) phase is not so much effective in improving the mechanical properties of alloys whose initial microstructures are composed of Ni3 Si(L12 ) + Ni3 Nb(D0a ) or Ni3 Si(L12 ) + Ni3 Ti(D024 ). (4) In alloys whose initial microstructures are composed of Ni3 Si(L12 ) + Ni3 Nb(D0a ), a good combination of tensile strength and tensile elongation was found over a wide of test temperature whether or not they contain the precipitated Ni3 Ti (D024 ) phase. This result suggests that the primary coarse Ni3 Nb (D0a ) phase precipitates are beneficial to improving the mechanical properties. Acknowledgement This work was supported in part by the Grant-in-aid for Scientific Research (B) from the Ministry of Education, Culture, Sports and Technology. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10]
The aging phenomenon and the associated hightemperature tensile property were evaluated, using alloys whose initial microstructures are composed of Ni3 Si(L12 ), Ni3 Si(L12 ) + Ni3 Ti(D024 ) and Ni3 Si(L12 ) + Ni3 Nb(D0a ). The following results were obtained from this study.
[11] [12] [13] [14]
(1) The plate-like Ni3 Ti (D024 ) phases were precipitated from the Ni3 Si (L12 ) phase matrix at high temperature, resulting in age hardening behavior in the temperature range from 923 K to 1323 K. (2) It was shown by tensile test that the enhanced flow strength and the reduced tensile elongation due to the precipitation of the Ni3 Ti (D024 ) phase were found for alloy 1 whose initial microstructure is composed of the Ni3 Si (L12 ) single-phase matrix.
[16]
[15]
[17] [18] [19] [20] [21]
D.P. Pope, S.S. Ezz, Int. Mater. Rev. 29 (1984) 136. N.S. Stoloff, Int. Mater. Rev. 34 (1989) 153. M. Yamaguchi, Y. Umakoshi, Prog. Mater. Sci. 34 (1990) 1. T. Suzuki, Y. Mishima, S. Miura, ISIJ Int. 29 (1989) 1. R. Lowire, Trans. Metall. Soc. AIME 194 (1952) 1093. P.H. Thornton, R.G. Davies, Metall. Trans. 1 (1970) 549. K. Hagihara, T. Nakano, Y. Umakoshi, Acta Mater. 48 (2000) 1469. K. Hagihara, T. Nakano, Y. Umakoshi, Scr. Mater. 48 (2003) 577. K. Hagihara, T. Nakano, Y. Umakoshi, Acta Mater. 51 (2003) 2623. K. Hagihara, T. Tanaka, T. Nakano, Y. Umakoshi, Acta Mater. 53 (2005) 5051. K. Tomihisa, Y. Kaneno, T. Takasugi, Intermetallics 10 (2002) 247. K. Ohira, Y. Kaneno, T. Takasugi, J. Mater. Sci. 39 (2004) 2295. K. Tomihisa, Y. Kaneno, T. Takasugi, Intermetallics 12 (2004) 317. K. Ohira, Y. Kaneno, T. Takasugi, Mater. Sci. Eng. A 399 (2005) 332. K. Ohira, Y. Kaneno, H. Tsuda, T. Takasugi, Intermetallics 14 (2006) 367. Y. Nunomura, Y. Kaneno, H. Tsuda, T. Takasugi, Intermetallics 12 (2004) 389. Y. Nunomura, Y. Kaneno, H. Tsuda, T. Takasugi, Acta Mater. 54 (2006) 851. S. Shibuya, Y. Kaneno, H. Tsuda, T. Takasugi, Intermetallics 15 (2007) 338. W. Soga, Y. Kaneno, T. Takasugi, Intermetallics 14 (2006) 170. M.F. Ashby, in: A. Kelly, R.B. Nicholoson (Eds.), Strengthening Methods in Crystal, New York, 1971, p. 137. F.J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, second ed., Elsevier, Amsterdam, 2004, p. 293.
The effect of second-phase dispersions on mechanical property of Ni3Si based multi-phase intermetallic alloys M. Fujita a , Y. Kaneno a , T. Takasugi a,b,∗ a
Department of Materials Science, Graduate School of Engineering, Osaka Prefecture University, 1-1 Gakuen-cho, Naka-ku, Sakai, Osaka 599-8531, Japan b Osaka Center for Industrial Materials Research, Institute for Materials Research, Tohoku University, 1-1 Gakuen-cho, Naka-ku, Sakai, Osaka 599-8531, Japan Received 10 November 2006; received in revised form 17 April 2007; accepted 20 April 2007
Abstract Using alloys whose initial microstructures are composed of Ni3 Si(L12 ), Ni3 Si(L12 ) + Ni3 Ti(D024 ) and Ni3 Si(L12 ) + Ni3 Nb(D0a ), aging phenomenon and the associated high-temperature tensile property were investigated. The plate-like Ni3 Ti (D024 ) phase was precipitated from the Ni3 Si (L12 ) matrix at high temperature. It was shown by micro-hardness measurement that age hardening behavior due to the precipitation of the Ni3 Ti (D024 ) phase occurs in all alloys in the temperature range from 923 K to 1323 K. It was however shown by tensile test that the precipitated Ni3 Ti (D024 ) phase is not so much effective in improving the mechanical properties of alloys whose initial microstructures are composed of Ni3 Si(L12 ) + Ni3 Nb(D0a ) or Ni3 Si(L12 ) + Ni3 Ti(D024 ). In alloys whose initial microstructures are composed of Ni3 Si(L12 ) + Ni3 Nb(D0a ), a good combination of tensile strength and tensile elongation was found over a wide of test temperature regardless of presence or absence of the precipitated Ni3 Ti (D024 ) phase. It is suggested that the primary coarse Ni3 Nb (D0a ) phase precipitates are beneficial to improving the mechanical properties. © 2007 Elsevier B.V. All rights reserved. Keywords: Multi-phase intermetallics; Ni3 Si; Ni3 Ti; Ni3 Nb; GCP phase; Microstructure; Mechanical property
1. Introduction Geometrically close packed (GCP) Ni3 X phases combined with IIIb and IVb non-transition elements (e.g., Al and Si) [1–6], and also IVa and Va transition metals (e.g., Ti, V and Nb) [7–10] generally exhibit high thermal, chemical and microstructural stabilities, and also attractive mechanical properties such as strength anomaly at high temperatures [5–10], and therefore have been of great interest as a strengthener in Ni-based superalloys. However, the GCP Ni3 X intermetallic phases as well as many other intermetallic compounds show some drawbacks, such as poor ductility at room temperature and inferior creep strength at high temperature. It was recently shown that multi-phase intermetallic alloys, e.g., composed of GCP Ni3 X (X: Al, Ti and Nb) [11,12], (X: Si, Ti and Nb) [13–15], (X: Al, Ti and V) [16–18] and
∗ Corresponding author at: Department of Materials Science, Graduate School of Engineering, Osaka Prefecture University, 1-1 Gakuen-cho, Naka-ku, Sakai, Osaka 599-8531, Japan. Tel.: +81 72 254 9314; fax: +81 72 254 9912. E-mail address: [email protected] (T. Takasugi).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.04.085
(X: Al, Nb and V) [19] intermetallic phases are thermodynamically stable and exhibit stable microstructures at high temperatures. Also, it was expected that such GCP Ni3 Xbased multi-phase intermetallic alloys have coherent interfaces among their constituents because of their closely related crystal structures. With regard to the multi-phase Ni3 Si–Ni3 Ti–Ni3 Nb intermetallic alloys, an isothermal phase diagram constructed at 1323 K showed that there were a three-phase region consisting of Ni3 Si(L12 )–Ni3 Ti(D024 )–Ni3 Nb(D0a ), and twophase regions consisting of Ni3 Si(L12 )–Ni3 Ti(D024 ) and Ni3 Ti(D024 )–Ni3 Nb(D0a ) [13–15]. It was also found that the Ni3 Ti(D024 ) phase is precipitated from the Ni3 Si(L12 ) phase matrix by a solid-state reaction [15]. It was shown by hightemperature tensile test that alloys in which the Ni3 Si(L12 ) matrix was precipitated by the Ni3 Ti(D024 ) and Ni3 Nb(D0a ) phases had the most favorable properties as high-temperature mechanical and chemical materials although the majority of these multi-phase intermetallic alloys resulted in high tensile strength accompanied with certain levels of tensile ductility over a wide range of temperatures [14]. In this study, the effect of second-phase dispersions on mechanical properties of multi-phase intermetallic alloys
M. Fujita et al. / Materials Science and Engineering A 476 (2008) 112–119 Table 1 Alloy compositions used in this study Alloy
No. 1 No. 2 No. 3 No. 4
Alloy compositions Ni (at.%)
Si (at.%)
Ti (at.%)
Nb (at.%)
B (wt ppm)
79.5 79.5 79.5 79.5
8.2 6.2 8.2 6.15
9.2 11.3 8.2 8.2
3.1 3.0 4.1 6.15
50 50 50 50
whose initial microstructures are composed of Ni3 Si(L12 ), Ni3 Si(L12 ) + Ni3 Ti(D024 ) and Ni3 Si(L12 ) + Ni3 Nb(D0a ) is investigated. First, the changes of microstructure and hardness due to the Ni3 Ti(D024 ) phase precipitation by aging are investigated. Then, high-temperature tensile properties of these multi-phase intermetallic alloys are evaluated in the unaged and aged conditions. 2. Experimental procedures Alloys used in this study were prepared from starting raw materials of 99.9 wt.% Ni, 99.999 wt.% Si, 99.9 wt.% Ti and 99.9 wt.% Nb. All alloys were doped with 50 ppm boron. Alloys were prepared as button ingots with a 50 mm diameter by arc melting under an argon gas atmosphere on a copper hearth using a tungsten electrode. Alloy compositions of the button ingots are shown in Table 1 and also plotted in Fig. 1, i.e., in the form of the Ni3 Si–Ni3 Ti–Ni3 Nb pseudo-ternary phase diagram reported previously [15]. An identical 79.5 at.% Ni content was kept for the prepared alloys. All button ingots were homogenized at 1323 K for 48 h in a vacuum, followed by furnace cooling to room temperature. The button ingots were sectioned to a proper size mostly along a solidified direction of the button ingot by an electro-discharge machine (EDM). The sectioned specimens were evacuated into a silica tube and then annealed at 1373 K for 24 h in a vacuum by which the specimens were solution-treated,
Fig. 1. Isothermal Ni3 Si–Ni3 Ti–Ni3 Nb pseudo-ternary phase diagram at 1323 K. Alloy compositions used in this study are plotted in this figure.
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and then water-quenched. For aging, these specimens were reevacuated into a silica tube and then annealed at temperatures ranging from 823 K to 1323 K for various times, followed by water quenching. Metallographic and structural observations for these heattreated specimens were conducted by optical microscopy (OM), X-ray diffraction (XRD) and scanning electron microscopy (SEM). X-ray diffraction profiles were obtained using Cu K␣ radiation. The determination of the constituent intermetallic phases was conducted on the basis of SEM electron-probe analysis (EPMA) combined with XRD analysis. Mechanical properties were evaluated by means of microhardness test and tensile test. For Vickers hardness measurement, a load of 300 g was used. Tensile specimens with a gauge dimension 2 mm × 1 mm × 10 mm were cut from the button ingots homogenized at 1323 K for 48 h. Tensile test was conducted using unaged specimens as well as aged specimens: the former specimens were annealed at 1373 K for 24 h, while the latter specimens were annealed at 1373 K for 24 h and subsequently aged at 1273 K for 12 h, by which almost maximum hardening is obtained. The surface of the tensile specimens was abraded with a fine SiC paper. The tensile test in the temperature range between room temperature and 1173 K was conducted in a vacuum (∼1.5 × 10−3 Pa) within a metal tube surrounded by an electric furnace. A nominal strain rate used in the tensile test was 1.67 × 10−4 s−1 . Tensile yield strength was determined at 0.2% offset strain. Tensile elongation was calculated from the load–displacement curve and defined as the plastic strain to fracture. The fracture surface of the tensile specimens was examined by a scanning electron microscope. 3. Results 3.1. Microstructures Fig. 2 shows initial microstructures of alloys solution-treated at 1373 K for 48 h. Alloy 1 exhibits almost L12 single-phase microstructure although a very small volume fraction (∼1%) of the Ni3 Nb (D0a ) dispersions are contained. Alloy 2 exhibits a two-phase microstructure consisting of the Ni3 Si (L12 ) and Ni3 Ti (D024 ) phases. The Ni3 Ti (D024 ) phase with a volume fraction of ∼28% exhibits coarse plate-like morphology within the L12 phase matrix. On the other hand, alloys 3 and 4 exhibit microstructures composed of the Ni3 Si (L12 ) phase matrix and the Ni3 Nb (D0a ) phase particles. The volume fraction of the Ni3 Nb (D0a ) phase dispersions was larger in alloy 4 (∼14%) than in alloy 3 (∼3.5%). Thus, the feature of the microstructures observed at 1373 K (Fig. 2) was basically consistent with the phase diagram (at 1323 K) shown in Fig. 1. However, a minor inconsistency with the phase diagram may be due to the change of the solubility limits of each constituent phase (i.e., Ni3 Si, Ni3 Ti and Ni3 Nb) by temperature. For an example, the solubility limit of the Ni3 Si(L12 ) phase against the Ni3 Ti(D024 ) phase expands with increasing temperature. Also, it is deduced that the Ni3 Ti (D024 ) phases observed in alloy 2 (Fig. 2(b)) as well as the Ni3 Nb (D0a ) phases observed in alloys 3 (Fig. 2(c)) and
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Fig. 2. Back scattering (BS)-SEM microstructures of alloys: (a) 1, (b) 2, (c) 3 and (d) 4 solution-treated at 1373 K for 48 h.
4 (Fig. 2(d)) were formed directly from liquid phase because both phase dispersions are large and composed of dendritic morphology. In all alloys, the precipitation of the Ni3 Ti (D024 ) phase occurred at temperatures above 823 K. Fig. 3 representatively shows microstructures of alloys 1–4 aged at a relatively high temperature 1223 K, indicating the microstructural changes during aging. The plate-like (or needle-like) Ni3 Ti (D024 ) phase precipitates appeared in the L12 phase matrix and the volume fraction and size of the precipitates tended to increase with increasing aging time for all alloys. Also, Fig. 3 indicates that the interfacial plane between the L12 and D024 phases are quite straight and coincides with the certain crystallographic planes. The previous study [15] suggested that the orientation relationship between both phases is expressed by; [0 1 1]L12 //[1 1 2¯ 0]D024
(1a)
(1¯ 1 1)L12 //(0 0 0 1)D024
(1b)
The closest packed atomic planes and directions in the L12 and D024 phases are parallel each other. For alloy 2, the aging resulted in a bimodal microstructure for the Ni3 Ti (D024 ) phase dispersions, i.e., (small precipitates introduced by aging) + (prior large dispersions formed during solidification). For alloys 3 and 4, the aging resulted in a three-phase microstructure in which the Ni3 Ti (D024 ) and Ni3 Nb (D0a ) phases are dispersed within the Ni3 Si (L12 ) phase matrix, being consequently consistent with the microstructure expected from the phase diagram (Fig. 1). Also, the volume fraction of the Ni3 Ti (D024 ) phase precipitates in alloy 4 was larger than that in alloy 3, again consistent with the phase diagram (Fig. 1). Thus, favorable microstructures in which fine Ni3 Ti (D024 ) phase precipitates
densely distribute in the Ni3 Si (L12 ) phase matrix were obtained by aging. 3.2. Age hardening behavior by hardness measurement Fig. 4 shows the variations of Vickers hardness with aging time for alloys 1, 2, 3 and 4. Here, it should be noted that the Vickers hardness was measured within the Ni3 Si (L12 ) phase matrix, avoiding coarse primary Ni3 Ti (D024 ) and Ni3 Nb phase (D0a ) particles existing in the initial microstructures. In unaged (i.e., as-solution-treated) condition, alloys showed hardness values of 370–390 Hv. For all alloys, age hardening occurred at temperatures above 823 K by precipitation of the Ni3 Ti (D024 ) phase, as understood from Fig. 3. As aging temperature increases, hardening occurred more quickly and then reached a maximum hardness at shorter aging time. For alloys 1, 3 and 4, the maximum hardness with approximately 460 Hv was observed by aging at 1273 K for 12 h while for alloy 2, the maximum hardness with approximately 445 Hv was observed by aging at 1173 K for 24 h. However, apparent over-age hardening, i.e., apparent softening did not occur even at the maximum temperature 1273 K and also even at the longest aging time more than 168 h (although their data were omitted from Fig. 4). This result indicates that the aged microstructures are fairly stable for prolonged annealing time at high temperature. In Fig. 4, the changes in Vickers hardness shows complicated behavior with aging temperature. That is, the hardness does not show monotonic increase with aging temperature (especially in alloy 3). At the moment, the explanation is not clearly presented, but may be associated with some changes in microstructure, according to an unusual increase of the over-saturation of solutes in L12 phase against to D024 phase with increasing temperature.
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Fig. 3. Back scattered (BS)-SEM microstructures of alloys: (a) 1, (b) 2, (c) 3 and (d) 4 aged at 1223 K for 12 h and 72 h, respectively.
3.3. Tensile behavior Fig. 5 shows changes of stress versus strain curves with temperature for all alloys tested in this study. Depending on tem-
perature and alloy, largely different stress versus strain curves were observed. Alloy 1 deformed at low temperature showed a relatively large strain-hardening rate, and then ruptured. However, alloy 1 did not show an apparent plastic deformability at
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Fig. 4. Changes of Vickers hardness by aging time for alloys: (a) 1, (b) 2, (c) 3 and (d) 4 aged at various temperatures.
temperature beyond 1100 K. For alloy 2, there was little plastic deformability over a whole test temperature. Like alloys 1, 3 and 4 deformed at low temperature showed a relatively large strain-hardening rate and then ruptured, leading to high plastic deformability. With increasing temperature, the strainhardening rate decreased. At a peak temperature (i.e., ∼900 K) where the yield stress shows the highest value, the strainhardening rate was still positive. Alloys 3 and 4 deformed at high temperature (beyond 900 K) exhibited a steady-state flow
after yielding, consequently resulting in high tensile elongation. From over-all evaluation for the tensile property, it is deduced that alloy 3 has a good combination of tensile stress and tensile elongation over a wide of test temperature while alloy 2 shows the worst result not only for strength property but also for tensile elongation property over a wide of test temperature. Figs. 6–9 plot yield stress, ultimate tensile stress (UTS) and tensile elongation as a function of temperature for alloys 1, 2, 3 and 4, respectively. For each alloy, the measurement was
Fig. 5. Stress vs. strain curves of unaged and aged alloys: (a) 1, (b) 2, (c) 3 and (d) 4 deformed at various temperatures, respectively.
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Fig. 6. Yield stress, tensile stress and tensile elongation as a function of test temperature for unaged and aged alloy 1, respectively. In aged alloy 1, the data for the yield stress at high temperatures are not plotted because the specimens fractured in the elastic region.
Fig. 7. Yield stress, tensile stress and tensile elongation as a function of test temperature for unaged and aged alloy 2, respectively. In alloy 2, the data for the yield stress at high temperatures are not plotted because the specimens fractured in the elastic region.
performed on unaged condition (as solution-treated at 1373 K for 24 h) and subsequently aged (at 1273 K for 12 h) condition, respectively. The yield stress increases with increasing temperature, and makes a broad maximum at around 873 K, followed by a decrease at high temperature. The reduced yield strength over a wide range of temperature was observed in alloy 2. Similar to the yield strength behavior, ultimate tensile stress increases with increasing temperature, and makes an apparent maximum
Fig. 8. Yield stress, tensile stress and tensile elongation as a function of test temperature for unaged and aged alloy 3, respectively.
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Fig. 9. Yield stress, tensile stress and tensile elongation as a function of test temperature for unaged and aged alloy 4, respectively.
at around 673–873 K, followed by a decrease at high temperature regardless of alloys. The reduced ultimate tensile strength over a wide range of temperature was again observed in alloy 2. Regarding tensile elongation, all alloys more or less exhibit certain levels of tensile elongation over a wide rage of temperature. The tensile elongation shows a peak at an intermediate temperature (∼673 K), and then shows a minimum at ∼1073 K, followed by a steep increase at high temperature, except for alloys 1 and 2. Alloy 2 consistently suffers from low tensile elongation over an entire test temperature while alloy 1 shows high tensile elongation at low temperature but low tensile elongation at high temperature. For aging effect, the increased yield strength and the reduced tensile elongation were observed for alloy 1. However, for the other alloys, aging effect was not evident not only on the flow strength but also on the tensile elongation although there was an apparent difference in the Vickers hardness values between the unaged and aged conditions. Also, it is noted that steep increases of tensile elongation take place in alloys 3 and 4 deformed at the highest temperature (1173 K). The fracture patterns of aged alloys 1–4 are shown in Fig. 10. The fracture patterns of unaged alloys were primarily similar to those of aged alloys at a whole test temperature and therefore omitted from this figure. The observed fracture patterns are basically consistent with the tensile ductility shown in Figs. 5–9. In other words, alloys with high tensile elongation showed ductile transgranular fracture mode with dimple-like patterns while alloys with low tensile elongation showed brittle fracture mode with intergranular patterns or large facetted planes. For alloy 1 that showed high tensile elongation at low temperature and low tensile elongation at high temperature, dimple-like fracture patterns accompanied with large facetted planes were observed at low temperature and intergranular fracture patterns were observed at high temperature (Fig. 10(a)). The facetted planes may correspond to the interfacial planes between the L12 phase matrix and the D024 phase precipitate. For alloy 2 with the lowest tensile elongation among four alloys, largely facetted fracture patterns were observed at each test temperature. It is likely that largely facetted fracture planes correspond to the interfacial plane between the coarse primary D024 phase plate and the L12 phase matrix (see Fig. 2(b)). For alloys 3 and 4 that showed relatively large tensile elongation at a whole test temperature, quite ductile fracture mode, i.e., transgranular
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Fig. 10. SEM fractography of aged alloys: (a) 1, (b) 2, (c) 3 and (d) 4 deformed at RT, 673 K and 1173 K, respectively.
fracture mode with dimple-like patterns was dominated. Also, the evidence that dynamic recrystallization occurs was shown at the highest temperature (1173 K): the fractured surfaces were covered by a numerous number of (recrystallized) small grains. 4. Discussion For alloy 1 whose initial microstructure is composed of the Ni3 Si (L12 ) single-phase matrix, micro-hardness measurement certainly showed age hardening and tensile test showed slightly enhanced yield strength due to the precipitation of the Ni3 Ti (D024 ) phase. Ni3 Ti was precipitated in a plate or needle-like form in the matrix as shown in Figs. 2 and 3. Direct TEM observation and the misfit evaluated from lattice parameters of Ni3 Si (L12 ) and Ni3 Ti (D024 ) phases [15] reveal that Ni3 Ti precip-
itates are coherent or partially coherent to the matrix. For the strengthening mechanism by Ni3 Ti precipitates, (i) generation of interfacial dislocations during the dislocation propagation event, (ii) transformation of the dislocation core during intersection of the interface according to the change of the crystal structure, or (iii) intersection of the network of the misfit dislocations situated at the interface may be suggested. It is consequently difficult to expect a substantial strengthening effect from these mechanisms as far as the present issue is concerned with shortterm mechanical property. Conversely, if the present issue is concerned with the long-term mechanical property (e.g., creep), fine and coherent precipitates may effectively enhance the hightemperature mechanical properties (e.g., strength). In addition, the size and spacing of the precipitated Ni3 Ti phase appear to be too large to induce the precipitation strengthening like
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Orowan’s mechanism. When the other alloys 2–4 whose initial microstructures are composed of Ni3 Si(L12 ) + Ni3 Ti(D024 ) or Ni3 Si(L12 ) + Ni3 Nb(D0a ) are evaluated by tensile test, the age hardening due to the precipitation of the Ni3 Ti (D024 ) phase was little effective: strength and tensile ductility were little affected by the precipitation of the Ni3 Ti (D024 ) phase at a wide range of test temperature. In alloys 3 and 4, Ni3 Nb (D0a ) precipitates are incoherent with Ni3 Si (L12 ) matrix [14,15]. During deformation, it is likely that dislocations are emitted preferentially from the precipitates or contrariwise piled up to the precipitates, consequently inducing the strengthening and plastic deformability through the entire material [20,21]. Thus, the obtained results imply that Ni3 Nb(D0a ) dispersions are more effective than Ni3 Ti (D024 ) precipitates not only on strength property but also on tensile ductility of the present multi-phase intermetallic alloys. These results reveal that the strengthening due to the second-phase incoherent with the matrix is more effective than that due to the second-phase coherent or partially coherent with the matrix, and also that the second-phase incoherent with the matrix is beneficial, enhancing the tensile ductility of the matrix. Thus, it is deduced that over-all mechanical properties are better in alloys whose initial microstructure is composed of the L12 phase matrix dispersed by the Ni3 Nb (D0a ) phase precipitates rather than in alloys whose initial microstructure is composed of the L12 single phase matrix or the L12 phase matrix precipitated by the coarse primary Ni3 Ti (D024 ) or precipitated Ni3 Ti (D024 ). 5. Conclusions
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(3) It was found that the precipitated Ni3 Ti (D024 ) phase is not so much effective in improving the mechanical properties of alloys whose initial microstructures are composed of Ni3 Si(L12 ) + Ni3 Nb(D0a ) or Ni3 Si(L12 ) + Ni3 Ti(D024 ). (4) In alloys whose initial microstructures are composed of Ni3 Si(L12 ) + Ni3 Nb(D0a ), a good combination of tensile strength and tensile elongation was found over a wide of test temperature whether or not they contain the precipitated Ni3 Ti (D024 ) phase. This result suggests that the primary coarse Ni3 Nb (D0a ) phase precipitates are beneficial to improving the mechanical properties. Acknowledgement This work was supported in part by the Grant-in-aid for Scientific Research (B) from the Ministry of Education, Culture, Sports and Technology. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10]
The aging phenomenon and the associated hightemperature tensile property were evaluated, using alloys whose initial microstructures are composed of Ni3 Si(L12 ), Ni3 Si(L12 ) + Ni3 Ti(D024 ) and Ni3 Si(L12 ) + Ni3 Nb(D0a ). The following results were obtained from this study.
[11] [12] [13] [14]
(1) The plate-like Ni3 Ti (D024 ) phases were precipitated from the Ni3 Si (L12 ) phase matrix at high temperature, resulting in age hardening behavior in the temperature range from 923 K to 1323 K. (2) It was shown by tensile test that the enhanced flow strength and the reduced tensile elongation due to the precipitation of the Ni3 Ti (D024 ) phase were found for alloy 1 whose initial microstructure is composed of the Ni3 Si (L12 ) single-phase matrix.
[16]
[15]
[17] [18] [19] [20] [21]
D.P. Pope, S.S. Ezz, Int. Mater. Rev. 29 (1984) 136. N.S. Stoloff, Int. Mater. Rev. 34 (1989) 153. M. Yamaguchi, Y. Umakoshi, Prog. Mater. Sci. 34 (1990) 1. T. Suzuki, Y. Mishima, S. Miura, ISIJ Int. 29 (1989) 1. R. Lowire, Trans. Metall. Soc. AIME 194 (1952) 1093. P.H. Thornton, R.G. Davies, Metall. Trans. 1 (1970) 549. K. Hagihara, T. Nakano, Y. Umakoshi, Acta Mater. 48 (2000) 1469. K. Hagihara, T. Nakano, Y. Umakoshi, Scr. Mater. 48 (2003) 577. K. Hagihara, T. Nakano, Y. Umakoshi, Acta Mater. 51 (2003) 2623. K. Hagihara, T. Tanaka, T. Nakano, Y. Umakoshi, Acta Mater. 53 (2005) 5051. K. Tomihisa, Y. Kaneno, T. Takasugi, Intermetallics 10 (2002) 247. K. Ohira, Y. Kaneno, T. Takasugi, J. Mater. Sci. 39 (2004) 2295. K. Tomihisa, Y. Kaneno, T. Takasugi, Intermetallics 12 (2004) 317. K. Ohira, Y. Kaneno, T. Takasugi, Mater. Sci. Eng. A 399 (2005) 332. K. Ohira, Y. Kaneno, H. Tsuda, T. Takasugi, Intermetallics 14 (2006) 367. Y. Nunomura, Y. Kaneno, H. Tsuda, T. Takasugi, Intermetallics 12 (2004) 389. Y. Nunomura, Y. Kaneno, H. Tsuda, T. Takasugi, Acta Mater. 54 (2006) 851. S. Shibuya, Y. Kaneno, H. Tsuda, T. Takasugi, Intermetallics 15 (2007) 338. W. Soga, Y. Kaneno, T. Takasugi, Intermetallics 14 (2006) 170. M.F. Ashby, in: A. Kelly, R.B. Nicholoson (Eds.), Strengthening Methods in Crystal, New York, 1971, p. 137. F.J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, second ed., Elsevier, Amsterdam, 2004, p. 293.