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Effects of Si and C additions on the thermal stability of directionally solidified TiAl–Nb alloys
Intermetallics 13 (2005) 1038–1047 www.elsevier.com/locate/intermet
Effects of Si and C additions on the thermal stability of directionally solidified TiAl–Nb alloys J.H. Kima, S.W. Kima, H.N. Leeb, M.H. Ohc, H. Inuid,*, D.M. Weea a
Department of Materials Science and Engineering, KAIST, Daejeon 305-701, South Korea b LG-Electronics Institute of Technology, Seoul 137-724, South Korea c Department of Materials Science and Engineering, KIT, Gumi 730-701, South Korea d Department of Materials Science and Engineering, Kyoto University, Sakyo-ku, Kyoto 606-8501, Japan Received 14 May 2004; received in revised form 14 September 2004; accepted 5 October 2004 Available online 3 May 2005
Abstract The thermal stability of the lamellar microstructure in TiAl–Nb alloys containing Si and C has been investigated by partial melting experiments. The proper compositions, where the lamellar structure is stable enough to be used as a seed material are found for both TiAl–Nb–Si and TiAl–Nb–Si–C alloy systems. The lamellar microstructure of the Ti–44.5Al–3Nb–0.6Si–0.2C and Ti–45Al–2Nb–0.6Si– 0.2C (at.%) alloys is indeed successfully aligned parallel to the growth direction by directional solidification (DS) with the seed material from the Ti–44.5Al–3Nb–0.6Si–0.2C alloy. The DS ingots of these alloys exhibit a good combination of room-temperature ductility (8.5%) and high-temperature yield strength (700 MPa at 800 8C). The proper composition range, where the lamellar structure is thermally stable enough to be used as the seed material for the TiAl–Nb–Si alloy is narrower than the corresponding range for previously investigated TiAl–Mo–Si and TiAl–Si alloys. The composition limits for such a region with the lamellar stability are discussed in terms of the critical volume fraction of the a (a2) phase for the Al-rich side limit and the a-peritectic composition for the Al-lean side limit. It is concluded that Nb is not an effective element to improve the lamellar stability because upon alloying with Nb, no significant change in the volume fraction of the a phase is expected to occur from the shift of the aCg/g phase boundary. q 2004 Published by Elsevier Ltd. Keywords: A. Titanium aluminides, based on TiAl; B. Thermal stability; C. Crystal growth; D. Microstructure; F. Mechanical testing
1. Introduction Alloys based on g-TiAl have long attracted considerable interest as a new class of material for high-temperature structural applications because of their low density, excellent high-temperature strength, and good oxidation resistance [1–3]. In particular, fully lamellar two-phase alloys consisting of TiAl (g) and Ti3Al (a2) have received considerable attention because of their high-temperature strength and room-temperature toughness, superior to those of other microstructure types such as near gamma, duplex, and nearly lamellar types [4]. However, the mechanical properties, especially the room-temperature ductility * Corresponding author. Tel.: C81-75-753-5467; fax: C81-75-753-5461 E-mail address: [email protected] (H. Inui). 0966-9795/$ - see front matter q 2004 Published by Elsevier Ltd. doi:10.1016/j.intermet.2004.10.010
and high-temperature strength of TiAl-based alloys of this microstructure type, have to be further improved for widespread industrial applications. Investigations on the mechanical properties of the g/a2 lamellar structure using so-called polysynthetically twinned (PST) crystals have revealed that the optimum combination of strength and ductility can be achieved when the lamellar boundaries are aligned parallel to the loading axis [5]. In order to take advantage of this plastic anisotropy of the lamellar structure, we have proposed a directional solidification (DS) processes by which the lamellar orientation is controlled to align parallel to the growth direction [6–8]. Because the lamellar microstructure is formed by precipitation of the g phase from the parent a phase upon cooling, following the Blackburn orientation relationship of (0001)a//(111)g and !1120O a ==! 110Og [9], one way to control the lamellar orientation is to control the orientation
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of the high-temperature a phase by finding alloy compositions, where the primary solidification phase is a. This was first achieved with an alloy with a particular composition; Ti–43Al–3Si (at.%) in the Ti–Al–Si ternary system [6]. With the use of this alloy for both the seed and feeder crystals, control of the lamellar orientation was successfully made [6]. This method has been further developed to produce a wide variety of TiAl-based alloys on top of the seed crystal of the Ti–43Al–3Si alloy [7,8]. TiAl-based alloys in which the lamellar orientation is successfully controlled with the seeding technique include those with a primary solidification phase of b [8,10–13]. However, the difference in composition between the seed (Ti–43Al–3Si) and feeder crystals is sometimes large enough to allow the formation of rather large mushy zones in the vicinity of the seed/feeder interface, which may become a problem in the industrial applications of the seeding technique. This problem can be largely avoided if the seed crystal has an alloy composition very close to that of the feeder crystal. In that case, for successful control of the lamellar orientation, the lamellar microstructure must be thermally stable upon heating to and cooling from just below the melting temperature without forming any new grains. We previously investigated the thermal stability of the lamellar microstructure of various alloys based on the Ti–46Al–1.5Mo alloy containing Si or C [14,15]. The lamellar microstructure of some of these alloys containing a small amount of Si or C, such as Ti–46Al–1.5Mo–1Si and Ti–46Al–1.5Mo–0.2C, was found to be stable, and the lamellar orientation was successfully controlled by using these seed crystals with stable compositions [14,15]. The Ti–46Al–1.5Mo–0.2C alloy, in particular, showed reasonable tensile properties at room temperature and good creep resistance at high temperatures [16]. In the present study, we investigate the thermal stability of the lamellar microstructure in TiAl–Nb alloys containing Si and C to find proper seed and feeder compositions, since Nb is reported to be one of the most important elements for providing TiAl-based alloys with good oxidation resistance and high-temperature mechanical properties [17–19]. We conduct directional solidification experiments to align the lamellar orientation parallel to the growth direction for TiAl–Nb–Si–C alloys in which we find high thermal stability of the lamellar structure. We then conduct mechanical tests on these directionally solidified ingots. We also discuss how proper alloy compositions, at which the g/a2 lamellar microstructure exhibits thermal stability, are determined depending on alloying element.
2. Experimental procedures Small buttons (25 g) of TiAl–Nb alloys containing Si and/or C were prepared by arc-melting in an argon atmosphere. The buttons were remelted at least eight times to promote homogeneity of the buttons. The nominal
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Table 1 The results from the partial melting experiments made for TiAl–3Nb–Si and TiAl–3Nb alloys Nominal composition
Calculated composition
Lamellar stability
Ti–44Al–3Nb–2Si Ti–45Al–3Nb–2Si Ti–46Al–3Nb–2Si Ti–44Al–3Nb–1Si Ti–44.5Al–3Nb–1Si Ti–45Al–3Nb–1Si Ti–45.5Al–3Nb–1Si Ti–46Al–3Nb–1Si Ti–45Al–3Nb–0.8Si Ti–47Al–3Nb–0.8Si Ti–46Al–3Nb Ti–47Al–3Nb Ti–48Al–3Nb
Ti–45.4Al–2Si Ti–46.4Al–2Si Ti–47.4Al–2Si Ti–45.4Al–1Si Ti–45.9Al–1Si Ti–46.4Al–1Si Ti–46.9Al–1Si Ti–47.4Al–1Si Ti–46.4Al–0.8Si Ti–48.4Al–0.8Si Ti–47.4Al Ti–48.4Al Ti–49.4Al
B ! ! ! ! ! ! ! ! ! ! ! !
Open circles and cross marks indicate that the lamellar microstructure is stable and unstable, respectively. The calculated compositions in the middle column are those calculated based on the Eq. (1).
compositions of alloys we investigated are tabulated in Tables 1 and 2 (alloy compositions are given in atomic% unless otherwise stated). Feeder ingots used for directional solidification processing were made by re-melting the arcmelted buttons into cylinder-type bars of 14 mm in diameter and 100 mm in length. The thermal stability of the lamellar microstructure of these alloys was evaluated by observing any changes in microstructures of directionally-solidified (DS) ingots after heating to and cooling from just below the melting temperature. This was accomplished by partially melting specimen cut from the DS ingot in an optical floating-zone furnace and comparing the heat-treated microstructure to that of the as-directionally solidified state. After finding out appropriate alloy compositions of high thermal stability for the lamellar structure, directional solidification experiments were made in the same optical Table 2 The results from the partial melting experiments made for TiAl–Nb–Si–C and TiAl–3Nb–C alloys Nominal composition
Calculated composition
Lamellar stability
Ti–44.5Al–3Nb–1Si–0.2C Ti–45Al–3Nb–0.8Si–0.2C Ti–43.5Al–3Nb–0.6Si–0.2C Ti–44.5Al–3Nb–0.6Si–0.2C Ti–45Al–3Nb–0.6Si–0.2C Ti–45Al–2Nb–0.6Si–0.2C Ti–45.5Al–3Nb–0.6Si–0.2C Ti–46Al–3Nb–0.6Si–0.2C Ti–45Al–3Nb–0.4Si–0.2C Ti–46Al–3Nb–0.4Si–0.2C Ti–47Al–3Nb–0.4Si–0.2C Ti–46Al–3Nb–0.2C Ti–47Al–3Nb–0.2C Ti–48Al–3Nb–0.2C
Ti–46Al–1Si Ti–46.5Al–0.8Si Ti–45Al–0.6Si Ti–46Al–0.6Si Ti–46.5Al–0.6Si Ti–46Al–0.6Si Ti–47Al–0.6Si Ti–47.5Al–0.6Si Ti–46.5Al–0.4Si Ti–47.5Al–0.4Si Ti–48.5Al–0.4Si Ti–47.5Al Ti–48.5Al Ti–49.5Al
! ! ! B ! B ! ! ! ! ! ! ! !
Open circles and cross marks indicate that the lamellar microstructure is stable and unstable, respectively. The calculated compositions in the middle column are those calculated based on the Eq. (1).
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floating-zone furnace at crystal growth rates from 5 to 10 mm/h with seed and feeder crystals of the same composition. A scanning electron microscope (SEM) equipped with an energy-dispersive spectrometer (EDS) and an optical microscope (OM) were used for microstructure characterization. Chemical etching with a solution of HF, HNO3 and H2O (1:1:18 by volume) was sometimes employed to facilitate the observation. Finer details of the lamellar microstructure of DS ingots were examined by a transmission electron microscope (TEM). Thin foils for TEM observations were prepared by twin-jet electro-polishing using a solution of methanol and perchloric acid. Tensile tests were performed at room temperature at a strain rate of 2!10K4 sK1 using a screw-driven universal testing machine. Flat tensile specimens were made from DS ingots with a spark-cutting machine so that the tensile axis was parallel to the growth direction. Each tensile specimen had dimensions of 5 mm in gage length and 2 mm!1 mm in cross section. Electro-polishing was made for all tensile specimens prior to testing. In addition, compression tests were also performed at room temperature, 800 and 1000 8C at a strain rate of 2!10K4 sK1 in vacuum for specimens with dimensions of 2.5 mm!2.5 mm!5 mm. The compression axis was parallel to the growth direction.
3. Results 3.1. Determination of suitable seed compositions The lamellar microstructure of the seed material must be stable upon heating to and cooling from just below the melting temperature without forming any new grains. In order to find out suitable alloy compositions for the seed materials, directional solidified ingots of the TiAl–3Nb–Si/C alloy systems grown at a growth rate of 20 mm/h were cut into two; one of the halves was partially melted while the other was kept in the as-directionally solidified state, as described previously [14,15]. Comparison of the two halves was then made by optical microscopy. The results from the partial melting experiments made on TiAl–3Nb–Si and TiAl–3Nb alloys are summarized in Table 1, where open circles and cross marks indicate that the lamellar microstructure is stable and unstable, respectively. As seen in Table 1, the lamellar microstructure is unstable for all alloys containing 1Si. While the lamellar microstructure is stable for the Ti–44Al–3Nb–2Si alloys, it is unstable in the case of the other two alloys containing 2Si. This clearly indicates that increasing the Si content does not simply improve the thermal stability of lamellar structure for all the TiAl–3Nb alloys, but improves it only for alloys of the particular composition range, i.e. the Al-lean side compositions. In other words, appropriate amounts of Al and Si contents have to be chosen in other for the material to be suitable seed material with high thermal stability for the lamellar structure.
Although the lamellar microstructure of the Ti–44Al– 3Nb–2Si alloy is thermally stable, the Si content of this alloy is high enough to allow the formation of numerous eutectic silicide particles, which are detrimental to the mechanical properties of TiAl-based alloys [6]. The Si content to avoid the formation of eutectic silicide particles is reported to be less than about 0.6Si [7]. Since C is reported to promote a-solidification as Si [15], we added a small amount of C to reduce the Si content to 0.6Si so that the formation of eutectic silicides is largely avoided. Fig. 1 shows some typical results of the partial melting experiments conducted on such TiAl–Nb–Si–C alloys. The microstructure after the partial melting experiment is very similar to that of the as-directionally solidified state for Ti–44.5Al–3Nb–0.6Si–0.2C and Ti–45Al–2Nb–0.6Si– 0.2C alloys (Fig. 1(b) and (d)), indicating the high thermal stability of the lamellar structure of these alloys. Indeed, the lamellar orientation of the partially melted piece of the Ti–44.5Al–3Nb–0.6Si–0.2C alloy is exactly the same as that of the as-directionally solidified state, as shown in Fig. 2. On the other hand, the formation of new lamellar grains is observed after the partial melting experiment for the other two alloys (Fig. 1(a) and (c)). The results from the partial melting experiments made on TiAl–Nb–Si–C alloys as well as on TiAl–3Nb–C alloys are summarized in Table 2. By adding a small amount of C (0.2C), some alloy compositions with high thermal stability of the lamellar microstructure can be found even with Si content as small as 0.6Si. However, the thermal stability of the lamellar microstructure of these alloys (Ti–44.5Al–3Nb–0.6Si– 0.2C and Ti–45Al–2Nb–0.6Si–0.2C) is easily lost with
Fig. 1. Microstructures of directionally-solidified ingots before (left) and after (right) the partial melting experiment; (a) Ti–43.5Al–3Nb–0.6Si– 0.2C, (b) Ti–44.5Al–3Nb–0.6Si–0.2C, (c) Ti–45Al–3Nb–0.6Si–0.2C and (d) Ti–45Al–2Nb–0.6Si–0.2C. The regions, where new lamellar grains are formed are marked with a circle.
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Fig. 2. Microstructures showing the counterparts of the lamellar microstructure of the Ti–44.5Al–3Nb–0.6Si–0.2C alloy before (a) and after (b) the partial melting experiment.
a small change in the Al, Si, and Nb contents. The lamellar microstructure is always unstable for alloys containing Si at less than 0.6 at.%. 3.2. Directional solidification using a stable seed alloy Of the two alloys with high thermal stability for the lamellar microstructure tabulated in Table 2, we have chosen the Ti–44.5Al–3Nb–0.6Si–0.2C alloy as the seed material for directional solidification experiments. A part of the directionally-solidified ingot of the Ti–44.5Al–3Nb– 0.6Si–0.2C alloy grown without a seed crystal was cut to be used as a seed crystal and was attached to a polycrystalline bar of the same composition by welding the two pieces
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within a floating-zone furnace. Care was taken to ensure that the lamellar boundaries of the seed crystal were oriented parallel to the growth direction. Using the seed crystal, directional solidification was carried out at a rate of 5 mm/h for the first 10 mm growth, and then at a rate of 10 mm/h for the remaining length. Fig. 3 shows a directionally-solidified Ti–44.5Al–3Nb–0.6Si–0.2C ingot grown from the seed material of the same composition. The lamellar microstructure of the grown ingot is continuous with that of the seed material and keeps the same lamellar orientation as the seed material (Fig. 3(b)). Thus, seeding to align the lamellar microstructure parallel to the growth direction is easily accomplished with directional solidification. The lamellar orientation of the Ti–45Al–2Nb–0.6Si–0.2C alloy is also found to be easily aligned parallel to the growth direction by using the Ti–44.5Al–3Nb–0.6Si–0.2C seed crystal. We have thus produced a number of DS ingots of these two compositions for mechanical testing. Fig. 4 shows microstructures of the Ti–44.5Al–3Nb– 0.6Si–0.2C alloy after arc-melting (a) and directional solidification at a growth rate of 5 mm/h (b). The primary solidification phase of this alloy can easily be determined by examining the side arm orientation of the dendrite spines of the arc-melted microstructure, which is cooled rapidly when compared to the other ingots. The primary phase is identified as b by noting that the dendrite arms are orthogonal to one another, as shown in Fig. 4(a) [6,7,15]. The core (bright) regions within the prior b dendrites consist of the B2 phase, which is found by EDS examination to be enriched with Nb, and silicide particles are found in the interdendritic regions. On the other hand, while some silicide particles are observed in the directionally solidified ingots, the B2 phase was never observed (Fig. 4(b)). This was further confirmed by TEM observations as shown in Fig. 5. Neither the B2 nor the carbide phase is observed in the as-solidified microstructures.
Fig. 3. Microstructures of the seed (left) and directionally solidified Ti–44.5Al–3Nb–0.6Si–0.2C ingot (right) showing the aligned g/a2 lamellar microstructure at (a) low and (b) high magnifications.
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Fig. 4. Microstructures (SEM backscattered electron images) of the Ti– 44.5Al–3Nb–0.6Si–0.2C alloy after arc-melting (a) and directional solidification at growth rates of 5 mm/h (b).
3.3. Mechanical properties of DS ingots The room-temperature yield stress and tensile elongation of the directionally solidified Ti–44.5Al–3Nb–0.6Si–0.2C and Ti–45Al–2Nb–0.6Si–0.2C (at.%) alloys with the lamellar microstructure aligned parallel to the loading axis are shown in Fig. 6. When these data are compared with those of the directionally solidified Ti–43Al–3Si and Ti– 46Al–1.5Mo–0.2C alloys in the previous studies [7,15], the yield strength of the present two alloys is not much different from that of the Mo-containing alloy, while the value of tensile elongation of the present two alloys is significantly
Fig. 5. TEM microstructure of the directionally solidified Ti–44.5Al–3Nb– 0.6Si–0.2C alloy (at a growth rate of 10 mm/h) showing the absence of small particles of the B2, silicide and carbide phases.
Fig. 6. Room-temperature yield stress (a) and tensile elongation (b) of directionally solidified ingots of the Ti–44.5Al–3Nb–0.6Si–0.2C and Ti– 45Al–2Nb–0.6Si–0.2C alloys with the lamellar boundaries aligned parallel to the loading axis. The corresponding data from directionally solidified ingots of the Ti–46Al–1.5Mo–0.2C [15] and Ti–43Al–3Si [7] alloys are also shown.
higher than those of the Ti–43Al–3Si and Ti–46Al–1.5Mo– 0.2C alloys. Of the four alloys, the Ti–45Al–2Nb–0.6Si– 0.2C alloy exhibits the best balance of yield stress (660 MPa) and tensile elongation (8.5%). The directionally solidified Ti–43Al–3Si alloy contains very large silicide particles in the interdendritic regions [7]. These brittle silicide particles are known to significantly reduce the tensile ductility of TiAl alloys. On the other hand, the B2 phase contained in the Ti–46Al–1.5Mo–0.2C alloy is known to be detrimental to the tensile ductility [15]. In contrast, both the Ti–44.5Al–3Nb–0.6Si–0.2C and Ti–45Al–2Nb– 0.6Si–0.2C alloys presently investigated contain no particles of the B2 and carbide phase, although they contain a small amount of the silicide phase. We thus believe that the high value of tensile elongation for these two alloys is attributed to the absence of second-phase particles. The yield stress obtained from compression tests conducted for directionally solidified ingots of the Ti– 44.5Al–3Nb–0.6Si–0.2C and Ti–45Al–2Nb–0.6Si–0.2C
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Fig. 7. Yield stress obtained in compression for directionally solidified ingots of the Ti–44.5Al–3Nb–0.6Si–0.2C and Ti–45Al–2Nb–0.6Si–0.2C alloys with the lamellar boundaries aligned parallel to the loading axis. The corresponding data from directionally solidified ingots of the Ti–46Al– 1.5Mo–0.2C [15] alloy are also shown.
alloys are shown in Fig. 7 as a function of temperature. The lamellar orientation for these ingots is aligned parallel to the compression axis. The corresponding data from directionally solidified ingots of the Ti–46Al–1.5Mo–0.2C alloy [15] are also shown in Fig. 7. For all the three alloys, the yield stress at 800 8C is just a little smaller than that at room temperature but it decreases considerably at 1000 8C. While there is no significant difference in yield stress between the Ti–45Al–2Nb–0.6Si–0.2C and Ti–46Al–1.5Mo–0.2C alloys, the yield stress of the Ti–44.5Al–3Nb–0.6Si–0.2C alloy is considerably higher than that for the other two alloys in the whole temperature range. In view of the fact that the yield stress of this alloy in tension is almost the same as those of the other two alloys (Fig. 6), the high yield stress of the Ti–44.5Al–3Nb–0.6Si–0.2C alloy in compression may be due to the variation of microstructure and/ or alloy compositions.
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of new a grains [7,20,21]. To meet the requirement (1), the alloy composition has to be located higher in the Al concentration than the peritectic composition of the LC a/b reaction (hereafter simply called the a-peritectic composition). If the Al content is lower than the a-peritectic composition, the a-phase is formed from the b-phase during cooling, resulting in the loss of the original lamellar microstructure through the Burgers orientation relationship between the two phases. Furthermore, for the requirement (2), it is preferable to have an enough a2-phase volume fraction at room temperature or an enough a-phase volume fraction at eutectoid temperature to avoid a sudden change in the a-phase volume fraction during rapid heating and cooling in the (aCg) two-phase region [21]. The alloy composition has to be well below the (aCg)/g phase boundary composition at the eutectoid temperature for this to occur. However, these requirements cannot be met for any alloy compositions in the Ti–Al binary system and proper alloying with elements of b-stabilizers (Nb, Mo) and a-stabilizers (Si, C) is required to achieve this [14,15]. Alloying b-stabilizer elements such as Nb and Mo promotes b-solidification and shifts the binary Ti–Al phase diagram to the Al-rich side [8,22,23]. Thus, alloying only with bstabilizer elements simply leads to the passage of b singlephase or bCa two-phase region during cooling, which results in the loss of the initial lamellar orientation. However, if the (aCg)/g boundary composition shits to the Al-rich side by alloying with b-stabilizers and at the same time, the a-peritectic composition shifts to the Al-lean side by alloying with a-stabilizers such as Si and C [6,15] upon simultaneous alloying, the appropriate situation for the thermal stability of the lamellar microstructure can be realized, as schematically shown in Fig. 8. The results from the partial melting experiments obtained previously for TiAl–Si ternary [7] and TiAl– 1.5Mo–Si quaternary [14] alloys are summarized in Table 3. It is clear from Table 3 that a wider composition range can be found for the thermal stability of the lamellar
4. Discussion 4.1. Effects of alloying elements on the thermal stability of TiAl alloys For seeding to be successful in directional solidification, the original lamellar orientation of the seed material should be maintained upon heating to and cooling from just below the melting temperature without forming any new grains. The seed material must thus have alloy compositions to meet the following requirements. (1) The b single-phase or bCa two-phase region must not be passed during cooling. (2) The volume fraction of the a-phase must not change significantly during heating and cooling so that the change in the volume fraction occurs by the thickening/thinning of a lamellae in the lamellar structure and not by the formation
Fig. 8. Ti–Al binary phase diagram modified to have a higher thermal stability for the lamellar structure as a result of simultaneous alloying with a- and b-stabilizer elements.
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Table 3 The results from the partial melting experiments obtained previously for TiAl–Si ternary [7] and TiAl–1.5Mo–Si quaternary [14] alloys Nominal composition
Calculated composition
Lamellar stability
Ti–47Al Ti–48Al Ti–43Al–1Si Ti–43Al–3Si Ti–45Al–2Si Ti–46Al–1Si Ti–47Al–1Si Ti–46Al–1.5Mo–3Si Ti–44.5Al–1.5Mo–2Si Ti–46Al–1.5Mo–2Si Ti–46A–1.5Mo–1.5Si Ti–44.5A–1.5Mo–1Si Ti–46Al–1.5Mo–1Si Ti–47Al–1.5Mo–1Si Ti–48Al–1.5Mo–1Si Ti–47Al–1.5Mo–0.5Si Ti–46Al–1.5Mo–0.4Si Ti–46Al–1.5Mo Ti–48Al–1.5Mo
– – – – – – – Ti–46.7Al–3Si Ti–45.2Al–2Si Ti–46.7Al–2Si Ti–46.7Al–1.5Si Ti–45.2Al–1Si Ti–46.7Al–1Si Ti–47.7Al–1Si Ti–48.7Al–1Si Ti–47.7Al–0.5Si Ti–46.7Al–0.4Si Ti–46.7Al Ti–48.7Al
! ! ! B B ! ! ! B B B ! B ! ! ! ! ! !
Open circles and cross marks indicate that the lamellar microstructure is stable and unstable, respectively. The calculated compositions in the middle column are those calculated based on the Eq. (1).
microstructure in the case of simultaneous alloying with Mo and Si than in the case of alloying with Si alone. It can be considered that the Mo additions effectively shift the phase diagram to the Al-rich side around the a-eutectoid temperature while the Si addition additions effectively shift the phase diagram to the Al-lean side around the a-peritectic temperature, modifying the diagram so as for the lamellar microstructure to have a high thermal stability, as schematically shown in Fig. 8. However, the addition of Si above a certain content level is detrimental to the thermal stability of the lamellar microstructure, because such a Si addition shifts not only the primary a-phase region but also the g-phase region to the Al-lean side, and that results in a decreased a2-phase volume fraction at low temperatures. Unlike the addition of Mo (Table 3), the addition of Nb to the TiAl–Si alloys cannot improve the thermal stability of the TiAl–Si ternary alloys (Table 1). Previous studies on liquidus surfaces of the TiAl–Nb and TiAl–Mo systems have indicated that the addition of 1Nb and 1Mo shifts the primary b-phase region to the Al-rich side by the amount corresponding to 0.3Al and 0.6Al, respectively [22,23]. This means that the addition of 3Nb is equivalent to the addition of 1.5Mo in terms of the shift the primary b-phase region. However, the effect of these elements on shifts of the phase boundaries around the a-eutectoid temperature is reported to be very different, as shown in Fig. 9 [24,25]. The addition of a small amount of Mo shifts the aCg/g phase boundary to the Alrich side considerably (Fig. 9(a)), resulting in the increased volume fraction of the a-phase around the a-3eutectoid temperature. Therefore, the change in the a-phase volume fraction during heating or cooling in the (aCg) two-phase
Fig. 9. Isothermal sections of (a) Ti–Al–Mo system at 1100–1200 8C [24] and (b) Ti–Al–Nb system at 1150 8C [25].
region can be reduced, resulting in the increased thermal stability of the lamellar structure of the TiAl–Si alloys upon alloying with Mo. On the other hand, the aCg/g phase boundary remains almost at a constant Al content upon alloying with Nb (Fig. 9(b)), and the volume fraction of the aphase may not change around the a-eutectoid temperature. Thus, Nb addition cannot contribute to reducing the change in a-phase volume fraction during heating and cooling and cannot improve the thermal stability of the TiAl–Si alloys. 4.2. Thermal stability of the lamellar microstructure in TiAl–Nb alloys containing Si and C In this section, we first discuss factors controlling the thermal stability of the lamellar microstructure of TiAl–Si ternary alloys and then extend the discussion to TiAl–Nb–Si and TiAl–Mo–Si quaternary alloys by taking into account of ‘Al-equivalent’ for Nb and Mo [22,23]. Previous studies on liquidus surfaces for the Ti–Al–Si ternary system have indicated that the addition of 1Si shifts the primary a-phase region to the Al-lean side by an amount corresponding to 2.8Al [6,8]. Assuming that the shift of the a-peritectic composition occurs in the same way and by the same amount as the primary a-phase region, the change in the a-peritectic composition (Al content) with Si additions can be indicated as shown in Fig. 10. The results of the partial melting experiments conducted for the Ti–Al–Si ternary alloys (Table 3) are plotted in Fig. 10 together with the line expressing the change in the a-peritectic composition. If the Al content is lower than the corresponding a-peritectic composition, the a-phase is formed from the b-phase during cooling, resulting in the loss of the original lamellar microstructure through the Burgers orientation relationship between the two phases. Therefore, the a-peritectic composition can be considered as the composition limit of
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Fig. 10. Schematic illustration of the composition limits of both the lower and higher Al sides for the lamellar stability and results of the partial melting experiments conducted for the TiAl–Si alloy system [7].
the lower Al side for lamellar stability. However, Ti–46Al– 1Si and Ti–47Al–1Si alloys were unstable even though they have a higher Al content than that in the corresponding a-peritectic composition. The reason for this instability is considered due to the small volume fraction of the a2 phase of these two alloys. In other words, the composition limit of the higher Al side for the lamellar stability is considered to be determined by the critical volume fraction of the a (a2) phase, since a significant change in the a/g phase volume fraction, which leads to the instability of the lamellar structure, is expected to occur during heating and cooling [21], if the volume fraction of the a2 phase is small, i.e. at higher Al contents. The supposed composition limit of the higher Al side for the lamellar stability, which was determined by partial melting experiments, is indicated as a dotted line in Fig. 10. The region delineated by these two limit lines corresponds to the alloy composition range, where the lamellar structure is thermally stable. In order to discuss the lamellar stability of the TiAl– 1.5Mo–Si and TiAl–3Nb–Si quaternary alloys in the same way, the line of the a-peritectic composition in the Ti–Al–Si system (Fig. 10) has to be shifted to the Al-rich side by an amount corresponding to the Al-equivalent for quaternary elements Nb and Mo, which was defined previously [11] as the shift of the a/b primary composition per 1 at.% addition of each of these elements; DCNbZC0.3, DCMoZC0.6 [22, 23]. Moreover, we will convert quaternary compositions (expressed as C) into ternary ones (expressed as C 0 ) corresponding to the TiAl–Si systems with the following equation. Ci0 =100ZCi =ð100KðCX CCY ÞÞ ðiZTi;Al;XZMo;Nb;Y ZCÞ 0 CSi yCSi
ðforsmallSiadditionÞ (1)
Thus, for example, using Eq. (1), the alloy Ti– 46Al–1.5Mo–1Si can be represented in the ternary system as Ti–46.7Al–1Si; C 0 Ti/100ZC Ti/(100KC Mo)Z51.5/ (100K1.5)Z52.3/100, C 0 AlZ46.7 and C 0 Siy1. These converted compositions for alloys of the TiAl–1.5Mo–Si and TiAl–3Nb–Si quaternary systems are shown in Tables 1 and 3, and these data are plotted in Fig. 11(a) and (b) together with the line expressing the change in the a-peritectic composition of those systems. The TiAl–1.5Mo–Si system (Fig. 11(a)) is seen to have a wider composition range for the
Fig. 11. Results of the partial melting experiments conducted for (a) TiAl– 1.5Mo–Si system [14], (b) TiAl–3Nb–Si system, and (c) TiAl–Nb–Si–C system. Quaternary compositions were transformed to ternary composition. Compositions, where the lamellar microstructure is thermally stable and unstable are expressed with closed circles and crosses, respectively. The dashed line in (c) indicates the change in the a-peritectic composition of TiAl–2Nb–Si–C system.
thermal stability of the lamellar microstructure than any other alloys investigated. The composition range for the lamellar stability of the TiAl–3Nb–Si system seems, however, to be narrower than that of the TiAl–Si ternary system, because the addition of Nb hardly affects the high Al limit for the lamellar stability of the TiAl–Si ternary system but only shifts the a-peritectic composition to the Al-rich side. This suggests that the size of the stable composition range depends on how much the b-stabilizers shift the aCg/g phase boundary to the Al-rich side; the range may be wider for the quaternary system with a stronger b-stabilizer. In view of this, the quaternary system containing W, which is a stronger b-stabilizer than Mo [23,26], may have a much wider composition range for lamellar stability. We reduced the Si content to 0.6% to avoid the formation of large silicide particles, which are detrimental to tensile ductility. It is reported that the formation of silicide particles results in a slight reduction in the a2-phase volume fraction [27,28]. Thus, the Si content needs to be reduced to improve both the tensile properties and the thermal stability of TiAl–
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3Nb–Si alloys. We put a small amount of C to solve above problems. Carbon promotes a-solidification more effectively than Si does [15]; the addition of 1C shifts the primary a-phase region to the Al-lean side by approximately 4Al. Thus, we can expect that the addition of 0.2C to shift the a-peritectic composition of the TiAl–3Nb–Si system (Fig. 11(b)) to the Al-lean side by 0.8Al. Alloy compositions in the TiAl–Nb–Si–C system can also be converted into those in the TiAl–Si ternary system by using Eq. (1). These converted compositions are tabulated in Table 2 and they are also plotted in Fig. 11(c). In Fig. 11(c), the dashed line indicates the change in the a-peritectic composition of the TiAl–2Nb–Si–C system. What is noteworthy here is that the stable composition range, which do not exist at a low Si content in the case of the TiAl–3Nb–Si system, is created at 0.6Si by adding 0.2C, although the stable composition range is very narrow. The thermal stability of the unstable Ti–44.5Al–3Nb–1Si alloy cannot be improved by adding 0.2C alone, but a stable composition can be found when the Si content is decreased to 0.6Si. This suggests that Ti–44.5Al–3Nb–1Si(–0.2C) alloys may be unstable due to the small a2-phase volume fraction, and that the reduction of Si contents and silicide particles may result in a slight increase of the a2-phase volume fraction and may contribute to improvement of the thermal stability. Therefore, the excellent thermal stability of Ti–44.5Al–3Nb–0.6Si–0.2C and Ti–45Al–2Nb–0.6Si–0.2C alloys may result from the increased a2-phase volume fraction by reducing the Si content without significantly changing the a-peritectic composition. Consequently, the ideal phase diagram with thermal stability, as shown in Fig. 8, can be created by alloying a small amount of Si and C to the TiAl–Nb system.
5. Conclusions The thermal stability of the lamellar microstructure in TiAl–Nb alloys containing Si and C was investigated by the partial melting experiments. The proper compositions, where the lamellar structure is stable enough to be used as a seed material were found for both TiAl–3Nb–Si and TiAl–Nb–Si–C alloy systems. The lamellar microstructure of the Ti–44.5Al–3Nb–0.6Si–0.2C and Ti–45Al–2Nb– 0.6Si–0.2C alloys was indeed successfully aligned parallel to the growth direction by directional solidification with the seed material from the Ti–44.5Al–3Nb–0.6Si–0.2C alloy. The DS ingots exhibited a promising combination of roomtemperature ductility and high-temperature yield strength. The composition range, where the lamellar structure is thermally stable enough to be used as the seed material for the TiAl–3Nb–Si alloy is narrower than the corresponding range for TiAl–1.5Mo–Si and TiAl–Si alloys. The composition limit for such a region with the lamellar stability was discussed in terms of the critical volume fraction of the a (a2) phase for the Al-rich side limit and the a-peritectic composition for the Al-lean side limit. It was concluded that
Nb was not an effective element to improve the lamellar stability because upon alloying with Nb, no significant change in the volume fraction of the a phase was expected to occur from the shift of the aCg/g phase boundary.
Acknowledgements This work was supported by Korea Research Foundation Grant (KRF-2004-042-D00111). One of the authors (J.H.Kim) is grateful to the BK21 (Brain Korea 21) project of the Korea Ministry of Education for financial support to conduct this research work at Kyoto University for three months. The authors would like to thank Dr David R. Johnson of Purdue University for useful advice and helpful discussion.
References [1] Yamaguchi M, Inui H. TiAl compounds for structural applications. In: Daloria R, Lewwandowski JJ, Liu CT, Martin PL, Miracle DB, Nathal MV, editors. Structural intermetallics. Warrendale, PA: TMS; 1993. p. 127. [2] Yamaguchi M, Inui H, Ito K. Acta Mater 2000;48:307. [3] Kim YW. J Metals 1994;46:30. [4] Lu YH, Zhang YG, Qiao LJ, Wang YB, Chen CQ, Chu WY. Mater Sci Eng 2000;A289:91. [5] Inui H, Nakamura A, Oh MH, Yamaguchi M. Acta Mater 1992;40: 3059. [6] Johnson DR, Inui H, Yamaguchi M. Acta Metall Mater 1996;44:2523. [7] Johnson DR, Masuda Y, Inui H, Yamaguchi M. Acta Metall Mater 1997;45:2523. [8] Muto S, Yamanaka T, Johnson DR, Inui H, Yamaguchi M. Mater Sci Eng 2000;A329–331:424. [9] Blackburn MJ. In: Jaffee RI, Promisel NE, editors. The science, technology, and application of titanium. Oxford: Pergamon Press; 1970. p. 663. [10] Yamanaka T, Johnson DR, Inui H, Yamaguchi M. Intermetallics 1999;7:779. [11] Muto S, Yamanaka T, Lee HN, Johnson DR, Inui H, Yamaguchi M. Adv Eng Mater 2001;3:391. [12] Johnson DR, Masuda Y, Inui H, Yamaguchi M. Mater Sci Eng 1997; A239–240:577. [13] Johnson DR, Lee HN, Muto S, Yamanaka T, Inui H, Yamaguchi M. Intermetallics 2001;9:923. [14] Lee HN, Johnson DR, Inui H, Oh MH, Wee DM, Yamaguchi M. Mater Sci Eng 2002;A329–331:19. [15] Lee HN, Johnson DR, Inui H, Oh MH, Wee DM, Yamaguchi M. Acta Mater 2000;48:3221. [16] Lee HN, Johnson DR, Inui H, Oh MH, Wee DM, Yamaguchi M. Intermetallics 2002;10:841. [17] Kim YW. In: Kim YW, Wagner R, Yamaguchi M, editors. Gamma titanium aluminides. Warrendale, PA: TMS; 1995. p. 637. [18] Kim YW. Mater Sci Eng 1995;A192–193:519. [19] Yoshihara M, Miura K. Intermetallics 1995;3:357. [20] Yamaguchi M, Johnson DR, Lee HN, Inui H. Intermetallics 2000;8: 511. [21] Kim SW, Lee HN, Oh MH, Yamaguchi M, Wee DM. In: George EP, Inui H, Mills MJ, Eggeler G, editors. Defect properties and related phenomena in intermetallic alloys. Warendale (PA): MRS; 2002. p. 249.
J.H. Kim et al. / Intermetallics 13 (2005) 1038–1047 [22] Omiya Y, Muto S, Yamanaka T, Johnson DR, Inui H, Yamaguchi M. In: George EP, Inui H, Mills MJ, Eggeler G, editors. Defect properties and related phenomena in intermetallic alloys. Warendale (PA): MRS; 2002. p. 237. [23] Jung IS, Jang HS, Oh MH, Lee JH, Wee DM. Mater Sci Eng 2002; A329–331:13.
[24] [25] [26] [27]
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Singh AK, Banerjee D. Metall Mater Trans 1997;28A:1735. Dign JJ, Hao SM. Intermetallics 1998;6:329. Hashimoto K, Kimura M, Mizuhara Y. Intermetallics 1998;6:667. Hsu FY, Klaar HJ, Wang GX, Dahms M. Mater Charact 1996; 36:371. [28] Tsuyama S, Mitao S, Minakawa K. Mater Sci Eng 1992;A153:451.
Effects of Si and C additions on the thermal stability of directionally solidified TiAl–Nb alloys J.H. Kima, S.W. Kima, H.N. Leeb, M.H. Ohc, H. Inuid,*, D.M. Weea a
Department of Materials Science and Engineering, KAIST, Daejeon 305-701, South Korea b LG-Electronics Institute of Technology, Seoul 137-724, South Korea c Department of Materials Science and Engineering, KIT, Gumi 730-701, South Korea d Department of Materials Science and Engineering, Kyoto University, Sakyo-ku, Kyoto 606-8501, Japan Received 14 May 2004; received in revised form 14 September 2004; accepted 5 October 2004 Available online 3 May 2005
Abstract The thermal stability of the lamellar microstructure in TiAl–Nb alloys containing Si and C has been investigated by partial melting experiments. The proper compositions, where the lamellar structure is stable enough to be used as a seed material are found for both TiAl–Nb–Si and TiAl–Nb–Si–C alloy systems. The lamellar microstructure of the Ti–44.5Al–3Nb–0.6Si–0.2C and Ti–45Al–2Nb–0.6Si– 0.2C (at.%) alloys is indeed successfully aligned parallel to the growth direction by directional solidification (DS) with the seed material from the Ti–44.5Al–3Nb–0.6Si–0.2C alloy. The DS ingots of these alloys exhibit a good combination of room-temperature ductility (8.5%) and high-temperature yield strength (700 MPa at 800 8C). The proper composition range, where the lamellar structure is thermally stable enough to be used as the seed material for the TiAl–Nb–Si alloy is narrower than the corresponding range for previously investigated TiAl–Mo–Si and TiAl–Si alloys. The composition limits for such a region with the lamellar stability are discussed in terms of the critical volume fraction of the a (a2) phase for the Al-rich side limit and the a-peritectic composition for the Al-lean side limit. It is concluded that Nb is not an effective element to improve the lamellar stability because upon alloying with Nb, no significant change in the volume fraction of the a phase is expected to occur from the shift of the aCg/g phase boundary. q 2004 Published by Elsevier Ltd. Keywords: A. Titanium aluminides, based on TiAl; B. Thermal stability; C. Crystal growth; D. Microstructure; F. Mechanical testing
1. Introduction Alloys based on g-TiAl have long attracted considerable interest as a new class of material for high-temperature structural applications because of their low density, excellent high-temperature strength, and good oxidation resistance [1–3]. In particular, fully lamellar two-phase alloys consisting of TiAl (g) and Ti3Al (a2) have received considerable attention because of their high-temperature strength and room-temperature toughness, superior to those of other microstructure types such as near gamma, duplex, and nearly lamellar types [4]. However, the mechanical properties, especially the room-temperature ductility * Corresponding author. Tel.: C81-75-753-5467; fax: C81-75-753-5461 E-mail address: [email protected] (H. Inui). 0966-9795/$ - see front matter q 2004 Published by Elsevier Ltd. doi:10.1016/j.intermet.2004.10.010
and high-temperature strength of TiAl-based alloys of this microstructure type, have to be further improved for widespread industrial applications. Investigations on the mechanical properties of the g/a2 lamellar structure using so-called polysynthetically twinned (PST) crystals have revealed that the optimum combination of strength and ductility can be achieved when the lamellar boundaries are aligned parallel to the loading axis [5]. In order to take advantage of this plastic anisotropy of the lamellar structure, we have proposed a directional solidification (DS) processes by which the lamellar orientation is controlled to align parallel to the growth direction [6–8]. Because the lamellar microstructure is formed by precipitation of the g phase from the parent a phase upon cooling, following the Blackburn orientation relationship of (0001)a//(111)g and !1120O a ==! 110Og [9], one way to control the lamellar orientation is to control the orientation
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of the high-temperature a phase by finding alloy compositions, where the primary solidification phase is a. This was first achieved with an alloy with a particular composition; Ti–43Al–3Si (at.%) in the Ti–Al–Si ternary system [6]. With the use of this alloy for both the seed and feeder crystals, control of the lamellar orientation was successfully made [6]. This method has been further developed to produce a wide variety of TiAl-based alloys on top of the seed crystal of the Ti–43Al–3Si alloy [7,8]. TiAl-based alloys in which the lamellar orientation is successfully controlled with the seeding technique include those with a primary solidification phase of b [8,10–13]. However, the difference in composition between the seed (Ti–43Al–3Si) and feeder crystals is sometimes large enough to allow the formation of rather large mushy zones in the vicinity of the seed/feeder interface, which may become a problem in the industrial applications of the seeding technique. This problem can be largely avoided if the seed crystal has an alloy composition very close to that of the feeder crystal. In that case, for successful control of the lamellar orientation, the lamellar microstructure must be thermally stable upon heating to and cooling from just below the melting temperature without forming any new grains. We previously investigated the thermal stability of the lamellar microstructure of various alloys based on the Ti–46Al–1.5Mo alloy containing Si or C [14,15]. The lamellar microstructure of some of these alloys containing a small amount of Si or C, such as Ti–46Al–1.5Mo–1Si and Ti–46Al–1.5Mo–0.2C, was found to be stable, and the lamellar orientation was successfully controlled by using these seed crystals with stable compositions [14,15]. The Ti–46Al–1.5Mo–0.2C alloy, in particular, showed reasonable tensile properties at room temperature and good creep resistance at high temperatures [16]. In the present study, we investigate the thermal stability of the lamellar microstructure in TiAl–Nb alloys containing Si and C to find proper seed and feeder compositions, since Nb is reported to be one of the most important elements for providing TiAl-based alloys with good oxidation resistance and high-temperature mechanical properties [17–19]. We conduct directional solidification experiments to align the lamellar orientation parallel to the growth direction for TiAl–Nb–Si–C alloys in which we find high thermal stability of the lamellar structure. We then conduct mechanical tests on these directionally solidified ingots. We also discuss how proper alloy compositions, at which the g/a2 lamellar microstructure exhibits thermal stability, are determined depending on alloying element.
2. Experimental procedures Small buttons (25 g) of TiAl–Nb alloys containing Si and/or C were prepared by arc-melting in an argon atmosphere. The buttons were remelted at least eight times to promote homogeneity of the buttons. The nominal
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Table 1 The results from the partial melting experiments made for TiAl–3Nb–Si and TiAl–3Nb alloys Nominal composition
Calculated composition
Lamellar stability
Ti–44Al–3Nb–2Si Ti–45Al–3Nb–2Si Ti–46Al–3Nb–2Si Ti–44Al–3Nb–1Si Ti–44.5Al–3Nb–1Si Ti–45Al–3Nb–1Si Ti–45.5Al–3Nb–1Si Ti–46Al–3Nb–1Si Ti–45Al–3Nb–0.8Si Ti–47Al–3Nb–0.8Si Ti–46Al–3Nb Ti–47Al–3Nb Ti–48Al–3Nb
Ti–45.4Al–2Si Ti–46.4Al–2Si Ti–47.4Al–2Si Ti–45.4Al–1Si Ti–45.9Al–1Si Ti–46.4Al–1Si Ti–46.9Al–1Si Ti–47.4Al–1Si Ti–46.4Al–0.8Si Ti–48.4Al–0.8Si Ti–47.4Al Ti–48.4Al Ti–49.4Al
B ! ! ! ! ! ! ! ! ! ! ! !
Open circles and cross marks indicate that the lamellar microstructure is stable and unstable, respectively. The calculated compositions in the middle column are those calculated based on the Eq. (1).
compositions of alloys we investigated are tabulated in Tables 1 and 2 (alloy compositions are given in atomic% unless otherwise stated). Feeder ingots used for directional solidification processing were made by re-melting the arcmelted buttons into cylinder-type bars of 14 mm in diameter and 100 mm in length. The thermal stability of the lamellar microstructure of these alloys was evaluated by observing any changes in microstructures of directionally-solidified (DS) ingots after heating to and cooling from just below the melting temperature. This was accomplished by partially melting specimen cut from the DS ingot in an optical floating-zone furnace and comparing the heat-treated microstructure to that of the as-directionally solidified state. After finding out appropriate alloy compositions of high thermal stability for the lamellar structure, directional solidification experiments were made in the same optical Table 2 The results from the partial melting experiments made for TiAl–Nb–Si–C and TiAl–3Nb–C alloys Nominal composition
Calculated composition
Lamellar stability
Ti–44.5Al–3Nb–1Si–0.2C Ti–45Al–3Nb–0.8Si–0.2C Ti–43.5Al–3Nb–0.6Si–0.2C Ti–44.5Al–3Nb–0.6Si–0.2C Ti–45Al–3Nb–0.6Si–0.2C Ti–45Al–2Nb–0.6Si–0.2C Ti–45.5Al–3Nb–0.6Si–0.2C Ti–46Al–3Nb–0.6Si–0.2C Ti–45Al–3Nb–0.4Si–0.2C Ti–46Al–3Nb–0.4Si–0.2C Ti–47Al–3Nb–0.4Si–0.2C Ti–46Al–3Nb–0.2C Ti–47Al–3Nb–0.2C Ti–48Al–3Nb–0.2C
Ti–46Al–1Si Ti–46.5Al–0.8Si Ti–45Al–0.6Si Ti–46Al–0.6Si Ti–46.5Al–0.6Si Ti–46Al–0.6Si Ti–47Al–0.6Si Ti–47.5Al–0.6Si Ti–46.5Al–0.4Si Ti–47.5Al–0.4Si Ti–48.5Al–0.4Si Ti–47.5Al Ti–48.5Al Ti–49.5Al
! ! ! B ! B ! ! ! ! ! ! ! !
Open circles and cross marks indicate that the lamellar microstructure is stable and unstable, respectively. The calculated compositions in the middle column are those calculated based on the Eq. (1).
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floating-zone furnace at crystal growth rates from 5 to 10 mm/h with seed and feeder crystals of the same composition. A scanning electron microscope (SEM) equipped with an energy-dispersive spectrometer (EDS) and an optical microscope (OM) were used for microstructure characterization. Chemical etching with a solution of HF, HNO3 and H2O (1:1:18 by volume) was sometimes employed to facilitate the observation. Finer details of the lamellar microstructure of DS ingots were examined by a transmission electron microscope (TEM). Thin foils for TEM observations were prepared by twin-jet electro-polishing using a solution of methanol and perchloric acid. Tensile tests were performed at room temperature at a strain rate of 2!10K4 sK1 using a screw-driven universal testing machine. Flat tensile specimens were made from DS ingots with a spark-cutting machine so that the tensile axis was parallel to the growth direction. Each tensile specimen had dimensions of 5 mm in gage length and 2 mm!1 mm in cross section. Electro-polishing was made for all tensile specimens prior to testing. In addition, compression tests were also performed at room temperature, 800 and 1000 8C at a strain rate of 2!10K4 sK1 in vacuum for specimens with dimensions of 2.5 mm!2.5 mm!5 mm. The compression axis was parallel to the growth direction.
3. Results 3.1. Determination of suitable seed compositions The lamellar microstructure of the seed material must be stable upon heating to and cooling from just below the melting temperature without forming any new grains. In order to find out suitable alloy compositions for the seed materials, directional solidified ingots of the TiAl–3Nb–Si/C alloy systems grown at a growth rate of 20 mm/h were cut into two; one of the halves was partially melted while the other was kept in the as-directionally solidified state, as described previously [14,15]. Comparison of the two halves was then made by optical microscopy. The results from the partial melting experiments made on TiAl–3Nb–Si and TiAl–3Nb alloys are summarized in Table 1, where open circles and cross marks indicate that the lamellar microstructure is stable and unstable, respectively. As seen in Table 1, the lamellar microstructure is unstable for all alloys containing 1Si. While the lamellar microstructure is stable for the Ti–44Al–3Nb–2Si alloys, it is unstable in the case of the other two alloys containing 2Si. This clearly indicates that increasing the Si content does not simply improve the thermal stability of lamellar structure for all the TiAl–3Nb alloys, but improves it only for alloys of the particular composition range, i.e. the Al-lean side compositions. In other words, appropriate amounts of Al and Si contents have to be chosen in other for the material to be suitable seed material with high thermal stability for the lamellar structure.
Although the lamellar microstructure of the Ti–44Al– 3Nb–2Si alloy is thermally stable, the Si content of this alloy is high enough to allow the formation of numerous eutectic silicide particles, which are detrimental to the mechanical properties of TiAl-based alloys [6]. The Si content to avoid the formation of eutectic silicide particles is reported to be less than about 0.6Si [7]. Since C is reported to promote a-solidification as Si [15], we added a small amount of C to reduce the Si content to 0.6Si so that the formation of eutectic silicides is largely avoided. Fig. 1 shows some typical results of the partial melting experiments conducted on such TiAl–Nb–Si–C alloys. The microstructure after the partial melting experiment is very similar to that of the as-directionally solidified state for Ti–44.5Al–3Nb–0.6Si–0.2C and Ti–45Al–2Nb–0.6Si– 0.2C alloys (Fig. 1(b) and (d)), indicating the high thermal stability of the lamellar structure of these alloys. Indeed, the lamellar orientation of the partially melted piece of the Ti–44.5Al–3Nb–0.6Si–0.2C alloy is exactly the same as that of the as-directionally solidified state, as shown in Fig. 2. On the other hand, the formation of new lamellar grains is observed after the partial melting experiment for the other two alloys (Fig. 1(a) and (c)). The results from the partial melting experiments made on TiAl–Nb–Si–C alloys as well as on TiAl–3Nb–C alloys are summarized in Table 2. By adding a small amount of C (0.2C), some alloy compositions with high thermal stability of the lamellar microstructure can be found even with Si content as small as 0.6Si. However, the thermal stability of the lamellar microstructure of these alloys (Ti–44.5Al–3Nb–0.6Si– 0.2C and Ti–45Al–2Nb–0.6Si–0.2C) is easily lost with
Fig. 1. Microstructures of directionally-solidified ingots before (left) and after (right) the partial melting experiment; (a) Ti–43.5Al–3Nb–0.6Si– 0.2C, (b) Ti–44.5Al–3Nb–0.6Si–0.2C, (c) Ti–45Al–3Nb–0.6Si–0.2C and (d) Ti–45Al–2Nb–0.6Si–0.2C. The regions, where new lamellar grains are formed are marked with a circle.
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Fig. 2. Microstructures showing the counterparts of the lamellar microstructure of the Ti–44.5Al–3Nb–0.6Si–0.2C alloy before (a) and after (b) the partial melting experiment.
a small change in the Al, Si, and Nb contents. The lamellar microstructure is always unstable for alloys containing Si at less than 0.6 at.%. 3.2. Directional solidification using a stable seed alloy Of the two alloys with high thermal stability for the lamellar microstructure tabulated in Table 2, we have chosen the Ti–44.5Al–3Nb–0.6Si–0.2C alloy as the seed material for directional solidification experiments. A part of the directionally-solidified ingot of the Ti–44.5Al–3Nb– 0.6Si–0.2C alloy grown without a seed crystal was cut to be used as a seed crystal and was attached to a polycrystalline bar of the same composition by welding the two pieces
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within a floating-zone furnace. Care was taken to ensure that the lamellar boundaries of the seed crystal were oriented parallel to the growth direction. Using the seed crystal, directional solidification was carried out at a rate of 5 mm/h for the first 10 mm growth, and then at a rate of 10 mm/h for the remaining length. Fig. 3 shows a directionally-solidified Ti–44.5Al–3Nb–0.6Si–0.2C ingot grown from the seed material of the same composition. The lamellar microstructure of the grown ingot is continuous with that of the seed material and keeps the same lamellar orientation as the seed material (Fig. 3(b)). Thus, seeding to align the lamellar microstructure parallel to the growth direction is easily accomplished with directional solidification. The lamellar orientation of the Ti–45Al–2Nb–0.6Si–0.2C alloy is also found to be easily aligned parallel to the growth direction by using the Ti–44.5Al–3Nb–0.6Si–0.2C seed crystal. We have thus produced a number of DS ingots of these two compositions for mechanical testing. Fig. 4 shows microstructures of the Ti–44.5Al–3Nb– 0.6Si–0.2C alloy after arc-melting (a) and directional solidification at a growth rate of 5 mm/h (b). The primary solidification phase of this alloy can easily be determined by examining the side arm orientation of the dendrite spines of the arc-melted microstructure, which is cooled rapidly when compared to the other ingots. The primary phase is identified as b by noting that the dendrite arms are orthogonal to one another, as shown in Fig. 4(a) [6,7,15]. The core (bright) regions within the prior b dendrites consist of the B2 phase, which is found by EDS examination to be enriched with Nb, and silicide particles are found in the interdendritic regions. On the other hand, while some silicide particles are observed in the directionally solidified ingots, the B2 phase was never observed (Fig. 4(b)). This was further confirmed by TEM observations as shown in Fig. 5. Neither the B2 nor the carbide phase is observed in the as-solidified microstructures.
Fig. 3. Microstructures of the seed (left) and directionally solidified Ti–44.5Al–3Nb–0.6Si–0.2C ingot (right) showing the aligned g/a2 lamellar microstructure at (a) low and (b) high magnifications.
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Fig. 4. Microstructures (SEM backscattered electron images) of the Ti– 44.5Al–3Nb–0.6Si–0.2C alloy after arc-melting (a) and directional solidification at growth rates of 5 mm/h (b).
3.3. Mechanical properties of DS ingots The room-temperature yield stress and tensile elongation of the directionally solidified Ti–44.5Al–3Nb–0.6Si–0.2C and Ti–45Al–2Nb–0.6Si–0.2C (at.%) alloys with the lamellar microstructure aligned parallel to the loading axis are shown in Fig. 6. When these data are compared with those of the directionally solidified Ti–43Al–3Si and Ti– 46Al–1.5Mo–0.2C alloys in the previous studies [7,15], the yield strength of the present two alloys is not much different from that of the Mo-containing alloy, while the value of tensile elongation of the present two alloys is significantly
Fig. 5. TEM microstructure of the directionally solidified Ti–44.5Al–3Nb– 0.6Si–0.2C alloy (at a growth rate of 10 mm/h) showing the absence of small particles of the B2, silicide and carbide phases.
Fig. 6. Room-temperature yield stress (a) and tensile elongation (b) of directionally solidified ingots of the Ti–44.5Al–3Nb–0.6Si–0.2C and Ti– 45Al–2Nb–0.6Si–0.2C alloys with the lamellar boundaries aligned parallel to the loading axis. The corresponding data from directionally solidified ingots of the Ti–46Al–1.5Mo–0.2C [15] and Ti–43Al–3Si [7] alloys are also shown.
higher than those of the Ti–43Al–3Si and Ti–46Al–1.5Mo– 0.2C alloys. Of the four alloys, the Ti–45Al–2Nb–0.6Si– 0.2C alloy exhibits the best balance of yield stress (660 MPa) and tensile elongation (8.5%). The directionally solidified Ti–43Al–3Si alloy contains very large silicide particles in the interdendritic regions [7]. These brittle silicide particles are known to significantly reduce the tensile ductility of TiAl alloys. On the other hand, the B2 phase contained in the Ti–46Al–1.5Mo–0.2C alloy is known to be detrimental to the tensile ductility [15]. In contrast, both the Ti–44.5Al–3Nb–0.6Si–0.2C and Ti–45Al–2Nb– 0.6Si–0.2C alloys presently investigated contain no particles of the B2 and carbide phase, although they contain a small amount of the silicide phase. We thus believe that the high value of tensile elongation for these two alloys is attributed to the absence of second-phase particles. The yield stress obtained from compression tests conducted for directionally solidified ingots of the Ti– 44.5Al–3Nb–0.6Si–0.2C and Ti–45Al–2Nb–0.6Si–0.2C
J.H. Kim et al. / Intermetallics 13 (2005) 1038–1047
Fig. 7. Yield stress obtained in compression for directionally solidified ingots of the Ti–44.5Al–3Nb–0.6Si–0.2C and Ti–45Al–2Nb–0.6Si–0.2C alloys with the lamellar boundaries aligned parallel to the loading axis. The corresponding data from directionally solidified ingots of the Ti–46Al– 1.5Mo–0.2C [15] alloy are also shown.
alloys are shown in Fig. 7 as a function of temperature. The lamellar orientation for these ingots is aligned parallel to the compression axis. The corresponding data from directionally solidified ingots of the Ti–46Al–1.5Mo–0.2C alloy [15] are also shown in Fig. 7. For all the three alloys, the yield stress at 800 8C is just a little smaller than that at room temperature but it decreases considerably at 1000 8C. While there is no significant difference in yield stress between the Ti–45Al–2Nb–0.6Si–0.2C and Ti–46Al–1.5Mo–0.2C alloys, the yield stress of the Ti–44.5Al–3Nb–0.6Si–0.2C alloy is considerably higher than that for the other two alloys in the whole temperature range. In view of the fact that the yield stress of this alloy in tension is almost the same as those of the other two alloys (Fig. 6), the high yield stress of the Ti–44.5Al–3Nb–0.6Si–0.2C alloy in compression may be due to the variation of microstructure and/ or alloy compositions.
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of new a grains [7,20,21]. To meet the requirement (1), the alloy composition has to be located higher in the Al concentration than the peritectic composition of the LC a/b reaction (hereafter simply called the a-peritectic composition). If the Al content is lower than the a-peritectic composition, the a-phase is formed from the b-phase during cooling, resulting in the loss of the original lamellar microstructure through the Burgers orientation relationship between the two phases. Furthermore, for the requirement (2), it is preferable to have an enough a2-phase volume fraction at room temperature or an enough a-phase volume fraction at eutectoid temperature to avoid a sudden change in the a-phase volume fraction during rapid heating and cooling in the (aCg) two-phase region [21]. The alloy composition has to be well below the (aCg)/g phase boundary composition at the eutectoid temperature for this to occur. However, these requirements cannot be met for any alloy compositions in the Ti–Al binary system and proper alloying with elements of b-stabilizers (Nb, Mo) and a-stabilizers (Si, C) is required to achieve this [14,15]. Alloying b-stabilizer elements such as Nb and Mo promotes b-solidification and shifts the binary Ti–Al phase diagram to the Al-rich side [8,22,23]. Thus, alloying only with bstabilizer elements simply leads to the passage of b singlephase or bCa two-phase region during cooling, which results in the loss of the initial lamellar orientation. However, if the (aCg)/g boundary composition shits to the Al-rich side by alloying with b-stabilizers and at the same time, the a-peritectic composition shifts to the Al-lean side by alloying with a-stabilizers such as Si and C [6,15] upon simultaneous alloying, the appropriate situation for the thermal stability of the lamellar microstructure can be realized, as schematically shown in Fig. 8. The results from the partial melting experiments obtained previously for TiAl–Si ternary [7] and TiAl– 1.5Mo–Si quaternary [14] alloys are summarized in Table 3. It is clear from Table 3 that a wider composition range can be found for the thermal stability of the lamellar
4. Discussion 4.1. Effects of alloying elements on the thermal stability of TiAl alloys For seeding to be successful in directional solidification, the original lamellar orientation of the seed material should be maintained upon heating to and cooling from just below the melting temperature without forming any new grains. The seed material must thus have alloy compositions to meet the following requirements. (1) The b single-phase or bCa two-phase region must not be passed during cooling. (2) The volume fraction of the a-phase must not change significantly during heating and cooling so that the change in the volume fraction occurs by the thickening/thinning of a lamellae in the lamellar structure and not by the formation
Fig. 8. Ti–Al binary phase diagram modified to have a higher thermal stability for the lamellar structure as a result of simultaneous alloying with a- and b-stabilizer elements.
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Table 3 The results from the partial melting experiments obtained previously for TiAl–Si ternary [7] and TiAl–1.5Mo–Si quaternary [14] alloys Nominal composition
Calculated composition
Lamellar stability
Ti–47Al Ti–48Al Ti–43Al–1Si Ti–43Al–3Si Ti–45Al–2Si Ti–46Al–1Si Ti–47Al–1Si Ti–46Al–1.5Mo–3Si Ti–44.5Al–1.5Mo–2Si Ti–46Al–1.5Mo–2Si Ti–46A–1.5Mo–1.5Si Ti–44.5A–1.5Mo–1Si Ti–46Al–1.5Mo–1Si Ti–47Al–1.5Mo–1Si Ti–48Al–1.5Mo–1Si Ti–47Al–1.5Mo–0.5Si Ti–46Al–1.5Mo–0.4Si Ti–46Al–1.5Mo Ti–48Al–1.5Mo
– – – – – – – Ti–46.7Al–3Si Ti–45.2Al–2Si Ti–46.7Al–2Si Ti–46.7Al–1.5Si Ti–45.2Al–1Si Ti–46.7Al–1Si Ti–47.7Al–1Si Ti–48.7Al–1Si Ti–47.7Al–0.5Si Ti–46.7Al–0.4Si Ti–46.7Al Ti–48.7Al
! ! ! B B ! ! ! B B B ! B ! ! ! ! ! !
Open circles and cross marks indicate that the lamellar microstructure is stable and unstable, respectively. The calculated compositions in the middle column are those calculated based on the Eq. (1).
microstructure in the case of simultaneous alloying with Mo and Si than in the case of alloying with Si alone. It can be considered that the Mo additions effectively shift the phase diagram to the Al-rich side around the a-eutectoid temperature while the Si addition additions effectively shift the phase diagram to the Al-lean side around the a-peritectic temperature, modifying the diagram so as for the lamellar microstructure to have a high thermal stability, as schematically shown in Fig. 8. However, the addition of Si above a certain content level is detrimental to the thermal stability of the lamellar microstructure, because such a Si addition shifts not only the primary a-phase region but also the g-phase region to the Al-lean side, and that results in a decreased a2-phase volume fraction at low temperatures. Unlike the addition of Mo (Table 3), the addition of Nb to the TiAl–Si alloys cannot improve the thermal stability of the TiAl–Si ternary alloys (Table 1). Previous studies on liquidus surfaces of the TiAl–Nb and TiAl–Mo systems have indicated that the addition of 1Nb and 1Mo shifts the primary b-phase region to the Al-rich side by the amount corresponding to 0.3Al and 0.6Al, respectively [22,23]. This means that the addition of 3Nb is equivalent to the addition of 1.5Mo in terms of the shift the primary b-phase region. However, the effect of these elements on shifts of the phase boundaries around the a-eutectoid temperature is reported to be very different, as shown in Fig. 9 [24,25]. The addition of a small amount of Mo shifts the aCg/g phase boundary to the Alrich side considerably (Fig. 9(a)), resulting in the increased volume fraction of the a-phase around the a-3eutectoid temperature. Therefore, the change in the a-phase volume fraction during heating or cooling in the (aCg) two-phase
Fig. 9. Isothermal sections of (a) Ti–Al–Mo system at 1100–1200 8C [24] and (b) Ti–Al–Nb system at 1150 8C [25].
region can be reduced, resulting in the increased thermal stability of the lamellar structure of the TiAl–Si alloys upon alloying with Mo. On the other hand, the aCg/g phase boundary remains almost at a constant Al content upon alloying with Nb (Fig. 9(b)), and the volume fraction of the aphase may not change around the a-eutectoid temperature. Thus, Nb addition cannot contribute to reducing the change in a-phase volume fraction during heating and cooling and cannot improve the thermal stability of the TiAl–Si alloys. 4.2. Thermal stability of the lamellar microstructure in TiAl–Nb alloys containing Si and C In this section, we first discuss factors controlling the thermal stability of the lamellar microstructure of TiAl–Si ternary alloys and then extend the discussion to TiAl–Nb–Si and TiAl–Mo–Si quaternary alloys by taking into account of ‘Al-equivalent’ for Nb and Mo [22,23]. Previous studies on liquidus surfaces for the Ti–Al–Si ternary system have indicated that the addition of 1Si shifts the primary a-phase region to the Al-lean side by an amount corresponding to 2.8Al [6,8]. Assuming that the shift of the a-peritectic composition occurs in the same way and by the same amount as the primary a-phase region, the change in the a-peritectic composition (Al content) with Si additions can be indicated as shown in Fig. 10. The results of the partial melting experiments conducted for the Ti–Al–Si ternary alloys (Table 3) are plotted in Fig. 10 together with the line expressing the change in the a-peritectic composition. If the Al content is lower than the corresponding a-peritectic composition, the a-phase is formed from the b-phase during cooling, resulting in the loss of the original lamellar microstructure through the Burgers orientation relationship between the two phases. Therefore, the a-peritectic composition can be considered as the composition limit of
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Fig. 10. Schematic illustration of the composition limits of both the lower and higher Al sides for the lamellar stability and results of the partial melting experiments conducted for the TiAl–Si alloy system [7].
the lower Al side for lamellar stability. However, Ti–46Al– 1Si and Ti–47Al–1Si alloys were unstable even though they have a higher Al content than that in the corresponding a-peritectic composition. The reason for this instability is considered due to the small volume fraction of the a2 phase of these two alloys. In other words, the composition limit of the higher Al side for the lamellar stability is considered to be determined by the critical volume fraction of the a (a2) phase, since a significant change in the a/g phase volume fraction, which leads to the instability of the lamellar structure, is expected to occur during heating and cooling [21], if the volume fraction of the a2 phase is small, i.e. at higher Al contents. The supposed composition limit of the higher Al side for the lamellar stability, which was determined by partial melting experiments, is indicated as a dotted line in Fig. 10. The region delineated by these two limit lines corresponds to the alloy composition range, where the lamellar structure is thermally stable. In order to discuss the lamellar stability of the TiAl– 1.5Mo–Si and TiAl–3Nb–Si quaternary alloys in the same way, the line of the a-peritectic composition in the Ti–Al–Si system (Fig. 10) has to be shifted to the Al-rich side by an amount corresponding to the Al-equivalent for quaternary elements Nb and Mo, which was defined previously [11] as the shift of the a/b primary composition per 1 at.% addition of each of these elements; DCNbZC0.3, DCMoZC0.6 [22, 23]. Moreover, we will convert quaternary compositions (expressed as C) into ternary ones (expressed as C 0 ) corresponding to the TiAl–Si systems with the following equation. Ci0 =100ZCi =ð100KðCX CCY ÞÞ ðiZTi;Al;XZMo;Nb;Y ZCÞ 0 CSi yCSi
ðforsmallSiadditionÞ (1)
Thus, for example, using Eq. (1), the alloy Ti– 46Al–1.5Mo–1Si can be represented in the ternary system as Ti–46.7Al–1Si; C 0 Ti/100ZC Ti/(100KC Mo)Z51.5/ (100K1.5)Z52.3/100, C 0 AlZ46.7 and C 0 Siy1. These converted compositions for alloys of the TiAl–1.5Mo–Si and TiAl–3Nb–Si quaternary systems are shown in Tables 1 and 3, and these data are plotted in Fig. 11(a) and (b) together with the line expressing the change in the a-peritectic composition of those systems. The TiAl–1.5Mo–Si system (Fig. 11(a)) is seen to have a wider composition range for the
Fig. 11. Results of the partial melting experiments conducted for (a) TiAl– 1.5Mo–Si system [14], (b) TiAl–3Nb–Si system, and (c) TiAl–Nb–Si–C system. Quaternary compositions were transformed to ternary composition. Compositions, where the lamellar microstructure is thermally stable and unstable are expressed with closed circles and crosses, respectively. The dashed line in (c) indicates the change in the a-peritectic composition of TiAl–2Nb–Si–C system.
thermal stability of the lamellar microstructure than any other alloys investigated. The composition range for the lamellar stability of the TiAl–3Nb–Si system seems, however, to be narrower than that of the TiAl–Si ternary system, because the addition of Nb hardly affects the high Al limit for the lamellar stability of the TiAl–Si ternary system but only shifts the a-peritectic composition to the Al-rich side. This suggests that the size of the stable composition range depends on how much the b-stabilizers shift the aCg/g phase boundary to the Al-rich side; the range may be wider for the quaternary system with a stronger b-stabilizer. In view of this, the quaternary system containing W, which is a stronger b-stabilizer than Mo [23,26], may have a much wider composition range for lamellar stability. We reduced the Si content to 0.6% to avoid the formation of large silicide particles, which are detrimental to tensile ductility. It is reported that the formation of silicide particles results in a slight reduction in the a2-phase volume fraction [27,28]. Thus, the Si content needs to be reduced to improve both the tensile properties and the thermal stability of TiAl–
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3Nb–Si alloys. We put a small amount of C to solve above problems. Carbon promotes a-solidification more effectively than Si does [15]; the addition of 1C shifts the primary a-phase region to the Al-lean side by approximately 4Al. Thus, we can expect that the addition of 0.2C to shift the a-peritectic composition of the TiAl–3Nb–Si system (Fig. 11(b)) to the Al-lean side by 0.8Al. Alloy compositions in the TiAl–Nb–Si–C system can also be converted into those in the TiAl–Si ternary system by using Eq. (1). These converted compositions are tabulated in Table 2 and they are also plotted in Fig. 11(c). In Fig. 11(c), the dashed line indicates the change in the a-peritectic composition of the TiAl–2Nb–Si–C system. What is noteworthy here is that the stable composition range, which do not exist at a low Si content in the case of the TiAl–3Nb–Si system, is created at 0.6Si by adding 0.2C, although the stable composition range is very narrow. The thermal stability of the unstable Ti–44.5Al–3Nb–1Si alloy cannot be improved by adding 0.2C alone, but a stable composition can be found when the Si content is decreased to 0.6Si. This suggests that Ti–44.5Al–3Nb–1Si(–0.2C) alloys may be unstable due to the small a2-phase volume fraction, and that the reduction of Si contents and silicide particles may result in a slight increase of the a2-phase volume fraction and may contribute to improvement of the thermal stability. Therefore, the excellent thermal stability of Ti–44.5Al–3Nb–0.6Si–0.2C and Ti–45Al–2Nb–0.6Si–0.2C alloys may result from the increased a2-phase volume fraction by reducing the Si content without significantly changing the a-peritectic composition. Consequently, the ideal phase diagram with thermal stability, as shown in Fig. 8, can be created by alloying a small amount of Si and C to the TiAl–Nb system.
5. Conclusions The thermal stability of the lamellar microstructure in TiAl–Nb alloys containing Si and C was investigated by the partial melting experiments. The proper compositions, where the lamellar structure is stable enough to be used as a seed material were found for both TiAl–3Nb–Si and TiAl–Nb–Si–C alloy systems. The lamellar microstructure of the Ti–44.5Al–3Nb–0.6Si–0.2C and Ti–45Al–2Nb– 0.6Si–0.2C alloys was indeed successfully aligned parallel to the growth direction by directional solidification with the seed material from the Ti–44.5Al–3Nb–0.6Si–0.2C alloy. The DS ingots exhibited a promising combination of roomtemperature ductility and high-temperature yield strength. The composition range, where the lamellar structure is thermally stable enough to be used as the seed material for the TiAl–3Nb–Si alloy is narrower than the corresponding range for TiAl–1.5Mo–Si and TiAl–Si alloys. The composition limit for such a region with the lamellar stability was discussed in terms of the critical volume fraction of the a (a2) phase for the Al-rich side limit and the a-peritectic composition for the Al-lean side limit. It was concluded that
Nb was not an effective element to improve the lamellar stability because upon alloying with Nb, no significant change in the volume fraction of the a phase was expected to occur from the shift of the aCg/g phase boundary.
Acknowledgements This work was supported by Korea Research Foundation Grant (KRF-2004-042-D00111). One of the authors (J.H.Kim) is grateful to the BK21 (Brain Korea 21) project of the Korea Ministry of Education for financial support to conduct this research work at Kyoto University for three months. The authors would like to thank Dr David R. Johnson of Purdue University for useful advice and helpful discussion.
References [1] Yamaguchi M, Inui H. TiAl compounds for structural applications. In: Daloria R, Lewwandowski JJ, Liu CT, Martin PL, Miracle DB, Nathal MV, editors. Structural intermetallics. Warrendale, PA: TMS; 1993. p. 127. [2] Yamaguchi M, Inui H, Ito K. Acta Mater 2000;48:307. [3] Kim YW. J Metals 1994;46:30. [4] Lu YH, Zhang YG, Qiao LJ, Wang YB, Chen CQ, Chu WY. Mater Sci Eng 2000;A289:91. [5] Inui H, Nakamura A, Oh MH, Yamaguchi M. Acta Mater 1992;40: 3059. [6] Johnson DR, Inui H, Yamaguchi M. Acta Metall Mater 1996;44:2523. [7] Johnson DR, Masuda Y, Inui H, Yamaguchi M. Acta Metall Mater 1997;45:2523. [8] Muto S, Yamanaka T, Johnson DR, Inui H, Yamaguchi M. Mater Sci Eng 2000;A329–331:424. [9] Blackburn MJ. In: Jaffee RI, Promisel NE, editors. The science, technology, and application of titanium. Oxford: Pergamon Press; 1970. p. 663. [10] Yamanaka T, Johnson DR, Inui H, Yamaguchi M. Intermetallics 1999;7:779. [11] Muto S, Yamanaka T, Lee HN, Johnson DR, Inui H, Yamaguchi M. Adv Eng Mater 2001;3:391. [12] Johnson DR, Masuda Y, Inui H, Yamaguchi M. Mater Sci Eng 1997; A239–240:577. [13] Johnson DR, Lee HN, Muto S, Yamanaka T, Inui H, Yamaguchi M. Intermetallics 2001;9:923. [14] Lee HN, Johnson DR, Inui H, Oh MH, Wee DM, Yamaguchi M. Mater Sci Eng 2002;A329–331:19. [15] Lee HN, Johnson DR, Inui H, Oh MH, Wee DM, Yamaguchi M. Acta Mater 2000;48:3221. [16] Lee HN, Johnson DR, Inui H, Oh MH, Wee DM, Yamaguchi M. Intermetallics 2002;10:841. [17] Kim YW. In: Kim YW, Wagner R, Yamaguchi M, editors. Gamma titanium aluminides. Warrendale, PA: TMS; 1995. p. 637. [18] Kim YW. Mater Sci Eng 1995;A192–193:519. [19] Yoshihara M, Miura K. Intermetallics 1995;3:357. [20] Yamaguchi M, Johnson DR, Lee HN, Inui H. Intermetallics 2000;8: 511. [21] Kim SW, Lee HN, Oh MH, Yamaguchi M, Wee DM. In: George EP, Inui H, Mills MJ, Eggeler G, editors. Defect properties and related phenomena in intermetallic alloys. Warendale (PA): MRS; 2002. p. 249.
J.H. Kim et al. / Intermetallics 13 (2005) 1038–1047 [22] Omiya Y, Muto S, Yamanaka T, Johnson DR, Inui H, Yamaguchi M. In: George EP, Inui H, Mills MJ, Eggeler G, editors. Defect properties and related phenomena in intermetallic alloys. Warendale (PA): MRS; 2002. p. 237. [23] Jung IS, Jang HS, Oh MH, Lee JH, Wee DM. Mater Sci Eng 2002; A329–331:13.
[24] [25] [26] [27]
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Singh AK, Banerjee D. Metall Mater Trans 1997;28A:1735. Dign JJ, Hao SM. Intermetallics 1998;6:329. Hashimoto K, Kimura M, Mizuhara Y. Intermetallics 1998;6:667. Hsu FY, Klaar HJ, Wang GX, Dahms M. Mater Charact 1996; 36:371. [28] Tsuyama S, Mitao S, Minakawa K. Mater Sci Eng 1992;A153:451.