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Microstructural refinement and improvement of mechanical properties and oxidation resistance in EPM TiAl-based intermetallics with yttrium addition
Acta Materialia 50 (2002) 1479–1493 www.actamat-journals.com
Microstructural refinement and improvement of mechanical properties and oxidation resistance in EPM TiAl-based intermetallics with yttrium addition Y. Wu a, b, S.K. Hwang a, c,* a Inha University, School of Materials Science and Engineering, Incheon 402-751, South Korea Beijing General Research Institute of Mining and Metallurgy, Beijing 100044, People’s Republic of China Jointly appointed by the Center for Advanced Aerospace Materials, Pohang University of Science and Technology Pohang, South Korea b
c
Received 29 August 2001; received in revised form 7 December 2001; accepted 7 December 2001
Abstract New Y-modified TiAl-based alloys, with nominal chemical compositions of Ti–46.6Al–1.4Mn–2Mo–0.3C–xY (x ⫽ 0.1, 0.33 and 0.6), were developed with the elemental powder metallurgy (EPM) method. The alloys with Y addition had higher ultimate tensile strength, elongation and oxidation resistance than the Y-free alloy. Combined solidsolution strengthening of Mn, Mo, C and Y, and the microstructural refining effect of C and Y, mainly contributed to the strength at room temperature as well as at high temperature up to 800°C, in addition to strengthening due to Y2O3 and carbide precipitation. The room-temperature ductility was also improved by the microstructural refinement and the removal of interstitial oxygen in the α2 and γ phases. Plastic deformation of γ plates during tensile testing occurred through the activation of ordinary dislocations of perfect Burger’s vectors. In an oxidation test in air, the Y-containing alloys showed a significantly low weight gain compared with the Y-free alloy. 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Titanium; Microstructure; Yttrium; Mechanical properties; Oxidation
1. Introduction In the last decade, intermetallic compounds based on TiAl have drawn increasing attention for aerospace structural applications in the intermediate temperature range (600–800°C) due to their
* Corresponding author. Tel.: +82-3286-07537; fax: +823286-25546. E-mail address: [email protected] (S.K. Hwang).
low density of 3.8 g/cm3 and good strength at elevated temperatures [1,2]. Recent research efforts have focused on chemical alloying and thermomechanical processing in order to balance the overall tensile properties at room temperature and at high temperatures [3,4]. Concerning the effects of alloying additions, Mn addition and C addition benefit the room-temperature ductility and the creep resistance, respectively; Mo addition increases the strength as well as the creep and oxidation resistance of the TiAl alloys [5–9]. Mean-
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while, Hwang and co-workers [10–15] have developed Ti–46.6Al–1.4Mn–2Mo–xC alloys of improved high-temperature tensile strength with the elemental powder metallurgy (EPM) method. The authors of these works attributed the improved mechanical properties to the microstructural refining effect, especially the lamellar refinement, of carbon. However, the room-temperature ductility of C-modified TiAl is still rather low, an elongation of approximately 1.6%. It was thought that the high oxygen content, approximately 1500 wt ppm, inherent to the EPM processing, was detrimental to ductility in these alloys. Tonnes et al. [16] reported that an increase in the oxygen content from 1000 wt ppm to 1600 wt ppm resulted in a sharp ductility drop from 2.1% to 0.5% in TiAl. In this respect, Y has an interesting role to play because this element, as an alloying element, is known to enhance the mechanical properties of other alloys [17]. The purpose of present work was to investigate the effect of Y addition on the tensile properties and oxidation resistance of TiAl–Mn–Mo–C intermetallics. Whether Y can scavenge the interstitial oxygen without interfering with the beneficial effects of other alloying elements was the question of primary interest.
2. Experimental The materials investigated for the present work were fabricated by EPM, the details of which are described elsewhere [10–15,18]. Of particular importance are the processing parameters such as powder degassing [10,11], extrusion can design and extrusion ratio [10], heating rate and holding time for the soaking treatment [11,18]. Specimens used for the present study were obtained with the proper control of these variables, thereby chemical homogeneity as well as microstructural uniformity were ensured. The mean size of powder particles of Ti, Al, Mn, Mo, C and Y ranged from 7 µm to 74 µm in experimental alloys. The nominal chemical compositions of the experimental alloys were Ti–46.6Al–1.4Mn–2Mo–0.3C–xY (x ⫽ 0.1, 0.33 and 0.6) (at%). Powders were mixed in a Vblender, compacted into Cr–Mo steel cans and
sealed. The cans were degassed sequentially at room temperature, 250°C and 600°C for 1, 2 and 4 h, respectively, followed by vacuum sealing at 3 × 10⫺5 torr. Degassed cans were hot-extruded at 1250°C with a ratio of 10:1. Prior to the hot extrusion, the cans were pre-heated at the same temperature and held for 12 h. After removal of the steel cans by machining, the specimens were heat-treated in an argon atmosphere. The heat treatment cycles consisted of a solution treatment at 1400°C for 1 h followed by air cooling (AC) and an aging treatment at 800°C. Transmission electron microscopy (TEM) specimens were prepared by standard jet-polishing techniques using a solution of 5% perchloric acid, 30% butanol and 65% methanol at ⫺25°C and 40 mA. Microstructural observation of the specimens was performed by optical microscopy (OM) and an H800 transmission electron microscope (TEM) operated at 200 kV. Quantitative measurements of the lamellar spacing were carried out using an image analyzer attached to the TEM. Chemical compositions of phases and precipitates in the specimens were analyzed by energy-dispersive X-ray spectrometry (EDS) in the TEM. Microhardness of the specimens was measured by a Vickers hardness tester using a load of 100 kg. Tensile specimens were 19 mm in gage length and 4 mm in gage diameter, and tested with a strain rate of 1 × 10⫺3 s⫺1. At least three samples were tested under identical conditions to obtain an average. The oxidation resistance of alloys was tested by isothermal exposure in air at 800°C for up to 350 h. Prior to the exposure, the specimens, 146 mm2 in surface area, were polished with emery papers of up to 1000 mesh finish. A thermo-balance was used to monitor the weight gain.
3. Results 3.1. Effect of Y on the microstructure There were significant differences between the microstructure of the Y-free alloy and that of the Y-added one. Fig. 1 shows optical micrographs of Ti–46.6Al–1.4Mn–2Mo–0.3C–xY alloys solutiontreated at 1400°C for 1 h and air-cooled. Regard-
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Fig. 1. Optical micrographs showing the fully lamellar microstructures of EPM Ti–46.6Al–1.4Mn–2Mo–0.3C alloys with different Y additions solution-treated at 1400°C /1 h / AC: (a) base composition, (b) 0.1Y, (c) 0.33Y and (d) 0.6Y. Each grain consists of colonies of aligned α2 ⫹ γ platelets.
less of Y content, the microstructures of experimental alloys showed a fully lamellar (FL) feature that consisted of colonies of directionally aligned α2 ⫹ γ platelets. A strong grain-refining effect of Y was observed as shown in Fig. 2: addition of 0.33Y to the base composition of Ti–46.6Al– 1.4Mn–2Mo–0.3C reduced the average grain size from 100 µm to approximately 45 µm and further addition of Y (0.6%) decreased it to approximately 30 µm. Y addition also refined the lamellar thickness of the TiAl alloy. An example of the refined lamellar structure is shown in Fig. 3, which compares the α2 ⫹ γ lamellar microstructures of Y-free and 0.33Y-added alloys, each solution-treated at 1400°C for 1 h/AC and aged at 800°C/12 h/AC.
Although Y addition refined the lamellar thickness it did not alter the flat feature of the lamellar interfaces, indicating that the individual morphology of γ and α2 plates was preserved. Twin-related α2 and γ plates were frequently observed, an example of which is shown in Fig. 3(b). In this figure, twinrelated γ-TiAl and α2-Ti3Al plates are shown with the zones of 具110]γ and [112¯ 0]α2, respectively. From the selected-area diffraction pattern (SADP) shown in Fig. 3(c), it was confirmed that the plate interfaces of γ phase and α2 phase were (111) and (0002), respectively. Among the γ plates there were twin relationships. With the rotation axis of [111], the twin plates of γ showed three variants: 120°, 180° and a pseudo-twin relationship. Orientation relationships of (111)γ储(0002)α2 and 具110]γ储
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Fig. 2. Grain-refining effect of Y in EPM Ti–46.6Al–1.4Mn– 2Mo–0.3C alloys α-solution heat-treated at 1400°C / 1 h /AC.
[112¯ 0]α2 were obtained, which are consistent with the Blackburn [19] relationships. Based on the TEM observation, a quantitative relationship between the average thickness of lamellae and Y content was obtained and the result is presented in Fig. 4. The measurement was carried out with an image analyzer attached to the TEM in the zone of 具110]γ储[112¯ 0]α2. At least 200 plates were measured for each condition to ensure reliability of the data. The average thickness of the γ plate decreased from 64 nm to 50 nm by the addition of 0.33% Y to the alloy with the base composition [Fig. 4(a)], while the thickness of the α2 plate also decreased from 28 nm to 22 nm [Fig. 4(b)]. A detailed TEM–EDS analysis showed a slight partitioning of Y in the γ phase. Measurements over 50 spots each for the α2 and γ plates resulted in the average concentration of Y of 0.41 wt% and 0.31 wt% and in the γ and α2 phase, respectively, with a scattering of ± 0.03 wt%. Although this analysis could not separate the contributions of the oxide particles from that of the matrix, it is interesting to note that an earlier study reported a preferential partitioning of oxygen in the α2 phase [20– 22]. In selective terms, this means that the oxygen content in the matrix of the γ plate will be substantially lower than that in the α2 plate due to the oxygen-scavenging effect of Y.
Subsequent extrusion and heat treatment, particularly the solution heat treatment, provide the oxide-forming environments. The relative amount of oxides formed in each stage is not known, although it is surmised that the partitioning of Y and the limited solubility of Y and O favor oxide formation in the γ phase. Due to the strong binding among Y and O atoms [23], the oxides of Y are thought to occur starting from the pre-extrusion heat treatment. Morphological characteristics and the chemical composition of oxide particles, Y2O3, were studied. As shown in Fig. 5(a), the oxide particles in the γ phase of Ti–46.6Al–1.4Mn–2Mo– 0.3C–0.33Y alloy aged at 800°C for 12 h showed an irregularly elongated shape that aligned along crystallographic directions. Oxide particles formed not only homogeneously within α2 and γ plates but also heterogeneously along the α2 / γ boundaries [Fig. 5(b)]. The particles had a hexagonal crystal structure, the lattice parameters determined from SADP being a ⫽ 0.376 nm and c ⫽ 0.582 nm [Fig. 5(c)]. The orientation relationships between the γ matrix and the precipitate were {111}γ储{0001}Y2O3 and 具011典γ储具112¯ 0典Y2O3. As expected, the EDS result revealed that precipitates were enriched with Y and oxygen [Fig. 5(d)]. Beside oxide particles, the aged Ti–46.6Al– 1.4Mn–2Mo–0.3C–0.33Y alloy also showed carbides, which ranged from 40 nm to 300 nm in size, and formed heterogeneously along α2 / γ grain boundaries and on the existing dislocations as shown in Fig. 6. EDS analysis in the TEM showed that the plate-like precipitates in Fig. 6(a) contained mainly Ti and Al, with a ratio of Ti/Al of approximately 3:1, and a small amount of Mn, Mo, C and Y. The crystal structure of the carbide was of the perovskite type (P phase), which was consistent with the earlier observation by Tian and Nemoto [9]. Regardless of their locations, cubictype orientation relationships between Ti3AlC and γ-TiAl phase were found, namely (001)p储(001)γ and [001]p储[001]γ. This result indicates that the crystallography of the precipitates was governed by the crystallography of the growing γ plates. On the other hand, the precipitates shown in Fig. 6(b) were determined to be carbide of Ti2AlC (H phase) type that had an ordered hexagonal crystal structure. The H phase also had strict orientation
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Fig. 3. TEM micrographs of EPM Ti–46.6Al–1.4Mn–2Mo–0.3C alloy with different additions of Y showing the lamellar-refining effect of Y: (a) Y-free alloy, (b) 0.33Y-added alloy and (c) selected-area diffraction pattern (SADP) taken from the area of α2 ⫹ γ plates in (b) with the zone axis of [1¯ 10]γ. The alloys were heat-treated at 1400°C / 1 h /AC ⫹ 800°C / 12 h / AC.
relationships with the matrix of γ plates, which coincided with the results of Tian and Nemoto [9]. According to the Ti–Al–C ternary phase diagram reported by Schuster et al. [24], Ti3AlC is in equilibrium with TiC, α2-Ti3Al and α-Ti, whereas Ti2AlC is in equilibrium with TiC, α2-Ti3Al and γ-TiAl. 3.2. Mechanical properties Evaluation of the mechanical properties of Yadded TiAl-based alloys was made on the basis of hardness and tensile properties. Regardless of heat treatment conditions, the hardness of the experimental alloys increased monotonically over the whole range (0.1 to 0.6%) of Y content studied.
The hardness increase was particularly notable for the Y content of 0.1 to 0.33%: it increased from 454Hv to 513Hv in the aged condition. As expected, the aged alloys showed higher hardness than the as-extruded or as-solution-treated alloys due to extensive precipitation of carbides. Y addition also improved the tensile strength of experimental alloys. The values of tensile strength and elongation of aged Ti–46.6Al–1.4Mn–2Mo– 0.3C–xY alloys measured at room temperature and 800°C are shown in Fig. 7 as a function of the Y content. The room-temperature ultimate tensile strength increased steeply with Y content up to 0.33% and then became saturated up to 0.6%, resulting in an overall increase of 20% compared with the Y-free alloy. This trend was repeated in
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Fig. 4. Quantitative measurement of the effect of Y on the lamellar thickness of EPM Ti–46.6Al–1.4Mn–2Mo–0.3C alloy: (a) γ plate and (b) α2 plate. The alloys were heat-treated at 1400°C / 1 h / AC ⫹ 800°C /12 h / AC.
the 800°C test although there was a gap of about 40 MPa between the strength at room temperature and that at the high temperature. Therefore the results of tensile testing were consistent with those of the hardness test. Because of the effect of Y, the room-temperature ultimate tensile strength reached the highest value of 704 MPa in the alloy containing 0.6% Y. Y addition also affected the ductility of experimental alloys. As shown in Fig. 7(b), a room-tem-
perature tensile elongation ranging from 1.6% to 3.3% was obtained for the experimental alloys of fully lamellar microstructure. Increasing the Y content from 0 to 0.1% promoted an increase in tensile elongation from 1.6% to 3.3%. However, excessive Y addition deteriorated ductility: a room-temperature elongation of 2.4% was obtained for the alloy containing 0.6% Y. The tensile ductility was generally higher at 800°C than the room-temperature values, the gap between the two values broadening with the Y content. Improved ductility in the Y-containing alloys was also evidenced by scanning electron microscopy (SEM) fractographs. In Fig. 8, two sets of fractographs are presented: one from a tensile specimen tested at room temperature [Fig. 8(a) and (b)] and the other from a test at 800°C [Fig. 8(c) and (d)]. Transgranular cleavage features predominated the room-temperature test specimen whereas a mixture of cleavage and intergranular fracture were apparent in the 800°C test specimen. In both cases addition of Y changed the fracture characteristic towards a more ductile mode, converting the brittle fracture to a plastic tear type of separation. The fracture surface of specimens tested at 800°C reflected another effect of Y addition in that the extent of oxidation was reduced. As inferred from the comparison of Fig. 8(c) with (d), the fracture surface of the Y-free alloy revealed a significant oxide layer whereas that of the Y-containing alloy exhibited a rather clean feature. This was an indication of an enhanced oxidation resistance by Y addition, which led to the next phase of investigation. 3.3. Deformed microstructure The slip mode of Y-free TiAl alloy was studied by examining the Burger’s vectors of dislocations in tensile specimens. In Fig. 9, bright-field TEM images of dislocations in a deformed specimen are shown for six different two-beam conditions. There were two different geometries of dislocations, leveled and tilted. Leveled dislocations were visible under all tilt conditions except for g ⫽ 22¯ 0 [Fig. 9(f)], which proved that the Burger’s vector of the dislocations was ordinary 1/2具110]. On the other hand, the Burger’s vector of tilted dislocations was
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Fig. 5. TEM micrographs showing Y2O3 formed in EPM Ti–46.6Al–1.4Mn–2Mo–0.3C alloy heat-treated at 1400°C /1 h / AC ⫹ 800°C /12 h / AC. (a) Y2O3 precipitates formed in a γ grain, (b) Y2O3 precipitates formed along the α2 / γ boundaries, (c) SADPs of the region shown in (a), from which the orientation relationship between the Y2O3 precipitates and the γ-TiAl matrix was determined, and (d) the EDS analysis of the oxides in (b).
super 具101典, some of which were dissociated. Therefore the γ grain deformed through the movement of mainly superdislocations of 具101典 type and some 1/2具110] ordinary dislocations. Compared with the Y-free alloy, dislocations in deformed specimens of Y-added alloys showed different characteristics. A set of TEM micrographs, shown in Fig. 10, was taken from the γ plate of a tensile specimen of 0.33 Y-containing alloy tested at room temperature, for which the Burger’s vectors of dislocations were determined. Like the case for the Y-free alloy (Fig. 9), two groups of dislocations were identified, leveled and tilted. From this analysis it was determined that both types of dislocation were of 1/2具110] Burger’s
vector. These dislocations were mostly straight and aligned along the 具1¯ 10] direction to form screw dislocations. Additionally, there were some cusps and loops along a dislocation line, hinting active cross-slips. Predominance of the ordinary dislocations in Ycontaining alloy was also evident in specimens tensile tested at 800°C. A typical dislocation structure in the γ plate of the alloy containing 0.33Y is shown in Fig. 11. With increase in the tensile test temperature, dislocations became gradually curved compared with those in room-temperature tensile test specimens. Through a series of diffraction contrast analyses, it was concluded that the majority of dislocations in the γ plate were still of 1/2具110]
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Fig. 6. TEM micrographs showing the carbide precipitates (a) along the α2 /γ boundaries and (b) on dislocations within the γ grain of EPM Ti–46.6Al–1.4Mn–2Mo–0.3C–0.33Y alloy aged at 800°C for 12 h. The P phase, Ti3AlC in chemical composition, has a perovskite structure with carbon atoms occupying the octahedral sites of L10 structure. The H phase, Ti2AlC in composition, has an ordered hexagonal structure.
type. Therefore increasing the deformation temperature from room temperature to 800°C promoted cross-slip activities but did not alter the preferential activation of ordinary 1/2具110] dislocations instead of superdislocations.
Fig. 7. Effect of Y content on the tensile properties of EPM Ti–46.6Al–1.4Mn–2Mo–0.3C–xY alloys: (a) tensile strength and (b) elongation. Samples were heat-treated at 1400°C / 1 h / AC ⫹ 800°C / 12 h / AC and tensile-tested at room temperature as well as at 800°C.
3.4. Oxidation resistance of Y-containing TiAl alloys A series of oxidation tests performed in air at 800°C demonstrated that Y addition improved the oxidation resistance of Ti–46.6Al–1.4Mn–2Mo– 0.3C–xY alloys. Shown in Fig. 12 are the curves of mass gains per unit surface area during the oxidation tests, which lasted for up to 350 h. The oxidation kinetic curves consisted of two stages. In
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Fig. 8. SEM fractographs of fractured tensile specimens of EPM Ti–46.6Al–1.4Mn–2Mo–0.3C–xY alloys: (a) base composition and (b) 0.1Y alloy tested at room temperature, and (c) base composition and (d) 0.1Y alloy tested at 800°C.
the first stage of oxidation ( ⬍ 25 h) mass gain increased linearly with time, which was followed by the second stage, in which a parabolic dependence on time was realized. After oxidation for 350 h, the mass gains for Y-free, 0.1Y-, 0.33Y- and 0.6Y-added alloys were 29.0, 19.7, 14.6 and 10.3 g/m2, respectively, reflecting a significant improvement in the oxidation resistance by Y addition.
4. Discussion As reported earlier by Morris and Leboeuf [25], grain refinement is an effective method to improve the tensile properties of TiAl-based alloys. In a previous study Hwang et al. showed that C
addition was beneficial for this purpose [12,15]. In the present study, Y was shown to have an additional effect on refining grains and lamellae. The grain-refining effect of Y was attributed to the precipitation of oxides that formed prior to extrusion. During extrusion and subsequent heat treatment, the oxides provide ample sites of heterogeneous nucleation for the product phases (α2 and γ from high-temperature α) while restraining their growth rate. In this respect, Y2O3 must be stable at 1400°C, the solution heat treatment temperature. According to the thermodynamical data given by Kobayashi and Tsukihashi [23], ⌬Gf° ⫽ 794,000⫺348T (J/mol, T in K) for the decomposition of Y2O3, and it is estimated that the oxide is stable up to 2009°C in TiAl.
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Fig. 9. TEM micrographs showing dislocation structures in a γ plate in EPM Ti–46.6Al–1.4Mn–2Mo–0.3C alloy under different diffraction conditions: (a) g ⫽ 220, (b) g ⫽ 02¯ 0, (c) g ⫽ 200, (d) g ⫽ 1¯ 31¯ , (e) g ⫽ 3¯ 11 and (f) g ⫽ 22¯ 0. Samples were heat-treated at 1400°C / 1 h /AC ⫹ 800°C / 12 h / AC and tensile-tested at room temperature.
Improvement of the tensile strength by Y addition, in addition to the strengthening effect by the perovskite carbides, is a natural consequence of the microstructural refinement. As for the ductility improvement, however, there are additional contributions by Y: first, an oxygen-scavenging effect and, second, the change of the slip mode. Y has a strong affinity to oxygen and, as such, it tends to
absorb the interstitial oxygen atoms from the matrix to form oxides, Y2O3 [17]. Since the interstitial oxygen is detrimental to ductility, this enhances the room-temperature elongation [16]. In general, the dislocation structure in L10-TiAl is complex in that there are 1/2具110] ordinary dislocations, two different types of superdislocation with Burger’s vector of 具011] and 1/2具112],
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Fig. 10. TEM micrographs showing dislocation structures in a γ plate in EPM Ti–46.6Al–1.4Mn–2Mo–0.3C–0.33Y alloy under different diffraction conditions: (a) g ⫽ 1¯ 11, (b) g ⫽ 002, (c) g ⫽ 1¯ 11¯ , (d) g ⫽ 1¯ 1¯ 1. Samples were heat-treated at 1400°C / 1 h / AC ⫹ 800°C /12 h / AC and tensile-tested at room temperature.
respectively, and 1/6具112] twin dislocations. A variety of microdeformation mechanisms in TiAl have been reported in the literature [26–34]. Huang and Hall [28] and Vasudevan et al. [31] found that 1/2具110] slip was the predominant deformation mode in two-phase TiAl alloy, which was also supported by Morris and Lipe in alloys based on Ti– 48Al–2Mn–2Nb [32]. For Ti–48Al alloy with high oxygen content, it was reported that deformation at high temperatures occurred through twinning instead of ordinary dislocations [33,34]. The present studies indicate that increasing concentrations of yttrium have a marked effect on the mobility of 1/2具110] ordinary dislocations regardless of deformation temperature, see Figs. 10 and 11. Thus, if the yttrium concentration is increased,
plasticity of the TiAl is enhanced most readily through the activity of 1/2具110] ordinary dislocations. Superdislocations appear to play a major role in plastic deformation only in the case of Yfree alloy, in which 1/2具110] ordinary dislocations have reduced mobility due to increased oxygen atoms in solution. These results are in concert with the assertion of Aindow et al. [35] and Vasudevan et al. [36], where the increased mobility of 1/2具110] ordinary dislocations in the Er-added TiAl alloys was attributed to the internal gettering of oxygen from solution by formation of Er2O3. Promotion of the activity of ordinary dislocations by Y was attributed to the counteraction to the adverse effect of interstitial elements such as C and O. Previous studies made by Morris et al.
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Fig. 11. TEM micrographs showing dislocation structures in a γ plate in EPM Ti–46.6Al–1.4Mn–2Mo–0.3C–0.33Y under different diffraction conditions: (a) g ⫽ 11¯ 0, (b) g ⫽ 1¯ 11¯ , (c) g ⫽ 1¯ 11, (d) g ⫽ 1¯ 13. Samples were heat-treated at 1400°C / 1 h / AC ⫹ 800°C /12 h / AC and tensile-tested at 800°C.
[37,38] also showed that ordinary dislocations were immobilized in the alloys with high oxygen concentrations. In this respect, the solubility of oxygen needs to be considered. The maximum solubility of oxygen in TiAl phase is 300 ± 150 at ppm, while the oxygen concentration in the experimental alloys was approximately 1500 wt ppm. Therefore it is deduced that a significant portion of the excessive oxygen atoms will be heterogeneously distributed by segregating to dislocations. In the O-rich alloy, it is believed that interstitial oxygen in solid solution is responsible for strong dislocation pinning. This may, of course, apply both to ordinary and to superdislocations. It is difficult to distinguish the oxygen atoms segregated at 1/2具110] ordinary dislocations from those
segregated at superdislocations. The results of observation by the present authors as well as others only hint that the oxygen concentration trapped in the 1/2具110] ordinary dislocations was higher than that in the superdislocations. The question of how oxygen may affect dislocation movement and why 1/2具110] ordinary dislocations may be more affected than 具101典 dislocations have been discussed by Kad and Fraser [39], but remains poorly understood. The higher lattice resistance due to the presence of the oxygen atoms seems to increase the difficulty of motion of ordinary dislocations to the extent that superdislocations are activated instead. In addition, there is a possibility that Y addition decreases the anti-phase boundary (APB) energy of superdislocations, making them more immobile.
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Fig. 12. Effect of Y on the oxidation kinetics of EPM Ti– 46.6Al–1.4Mn–2Mo–0.3C–xY alloys in air at 800°C. Samples were heat-treated at 1400°C /1 h / AC ⫹ 800°C /12 h / AC.
In summary, the Y addition affected the mechanical properties through multiple routes: grain refinement, lamellae refinement, interstitial scavenging, and promoting ordinary dislocation activity. However, it is noteworthy that there existed an optimum range of Y addition, 0.1 to 0.3%, since excessive addition causes embrittlement. The influence of alloying additions on isothermal oxidation behavior of TiAl-based alloys has been investigated by many researchers [40– 44]. A beneficial effect of Y on the oxidation resistance of various alloy systems has been reported [45–48]. As shown in Fig. 12, the alloys containing Y showed a significantly improved resistance to high-temperature oxidation in air. The alloys with Y addition showed smaller mass gains and smaller parabolic rate constants, which is in line with the previous studies mentioned above. It is poorly understood, however, how Y improves the oxidation resistance. There is a good possibility that Y modified the phase constitution and the microstructure of the oxide. Zhang et al. [41] studied the behaviors of the oxidation resistance of a H13 steel implanted with Y or Y plus carbon, and concluded that Y provided a barrier layer that hindered oxygen diffusion during thermal oxidation that occurs primarily along dislocations or grain boundaries. Han and Xiao [45] investigated the
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mechanism of improvement of the high-temperature isothermal oxidation resistance of Ni3Al-based alloys with added Y. They demonstrated that the beneficial effects of Y on high-temperature oxidation resistance. These results, in view of the present observation, suggest that Y addition suppresses the growth of TiO2 by promoting formation of Al2O3 scale. Although direct evidence to support this suggestion is lacking at present, this explanation is plausible for the present system based on TiAl. In the parabolic stage, oxidation in all experimental alloys obeyed the parabolic rate law that was represented by the equation (⌬M)2 ⫽ Kpt, where ⌬M, Kp and t are mass gain, parabolic rate constant and time, respectively. This process was controlled by cation transport through the scale. Because TiO2 oxides grow mainly by inward diffusion of oxygen via a vacancy mechanism [44], the improvement by Y addition in high-temperature oxidation behavior could be explained by its ability to reduce the oxygen vacancy concentration in the TiO2 scale due to the smaller oxide precipitates induced by Y [45–47]. Owing to their strong field strength with oxygen atoms, Y atoms are considered to reduce the oxygen vacancy concentration in TiO2 scale. The results presented in Fig. 12 indicate that this effect is proportional to the Y content, reaching a maximum at 0.6 at%. In addition, the improvement in oxidation resistance of the Y-added alloy could be interpreted by the classical Wagner theory [44], which explains the oxidation behavior in terms of oxygen solubility in the alloy. According to this interpretation, Y enhances the oxidation resistance by lowering the oxygen solubility through the formation of Y2O3 phase.
5. Conclusions From the study of the effect of Y on the mechanical properties and the oxidation resistance of a TiAl-based alloy, the following conclusions were drawn. 1. Partitioning of Y in the lamellar structure was found and quantitatively determined: 0.31 wt%
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Y and 0.41 wt% Y in the α2 and γ phases, respectively. 2. The orientation relationship between the γ matrix and the Y2O3 precipitate was determined and by TEM: {111}γ储{0001}Y2O3 具011典γ储具112¯ 0典Y2O3. 3. A strong grain-refinement effect of Y was found. As a result of the grain refinement, combined with the oxides and carbides precipitation and solid-solution hardening, the tensile strength of the experimental alloys improved significantly with Y addition at test temperatures of up to 800°C. The strength enhancement was accompanied by a concomitant increase in the ductility in the alloys with a moderate amount of Y addition, which was primarily due to the interstitial oxygen scavenging and enhanced activity of 1/2具110] ordinary dislocations. 4. Addition of Y also resulted in a substantial improvement in the oxidation resistance of the intermetallic compound exposed to air at 800°C. Acknowledgements The authors would like to thank Dr H.S. Ryoo, Mr Y.W. Park and Miss J.S. Park for their kind help. Y. Wu also acknowledges Dr W. Liang for TEM experiments in Taiyuan University of Technology. This work was performed under the auspices of the Korean Science and Engineering Fund through the Engineering Research Center program at POSTECH-CAAM, and by the Korean Ministry of Science and Technology through the “2001 National Research Laboratory Program”. References [1] [2] [3] [4] [5]
Kim YW. J Met 1989;41:24. Kim YW, Dimiduk DM. J Met 1991;43:40. Kim YW. J Met 1994;46:30. Kim YW. J Met 1995;47:39. Zhang G, Blenkinsop PA, Wise MLH. Intermetallics 1996;4:447. [6] Worth BD, Jones JW, Allison JE. Metall Trans A 1995;26:2961. [7] Gouma PI, Mills MJ, Kim YW. Phil Mag Lett 1998;78:59. [8] Gouma PI, Subramanian K, Kim YW, Mills MJ. Intermetallics 1998;6:689.
[9] Tian WH, Nemoto M. Intermetallics 1997;5:237. [10] Lee IS, Hwang SK, Park WK, Lee JH, Park DH, Kim HM et al. Scripta Metall Mater 1994;31:57. [11] Lee TK, Mosunov EI, Hwang SK. Mater Sci Eng A 1997;239:540. [12] Park HS, Nam SW, Kim NJ, Hwang SK. Scripta Mater 1999;41:1197. [13] Nam SW, Cho HS, Hwang SK, Kim NJ. Met Mater 2000;6:287. [14] Park HS, Park KL, Hwang SK. Intermetallics, in press. [15] Park HS, Hwang SK, Lee CM, Yoo YC, Nam SW, Kim NJ. Metall Trans A 2001;32:251. [16] Tonnes C, Rosler J, Baumann R, Thumann M. In: Dorolia R, Lewandowski JJ, Liu CT, Martin PL, Mirade DB, Nathal MY, editors. Structural intermetallics. Warrendale (PA): The Minerals, Metals and Materials Society; 1993. p. 241. [17] Yan YQ, Zhang ZQ, Luo GZ, Wang KG, Zhou L. Mater Sci Eng A 2000;280:187. [18] Lee TK, Kim JH, Hwang SK. Metall Trans A 1997;28:2723. [19] Blackburn MJ. In: Jaffee RT, Promised NE, editors. Proceedings of Technology and Application of Titanium, Oxford, UK: Pergamon Press, 1970:633. [20] Larson DJ, Miller MK. Mater Sci Eng A 1998;250:65. [21] Uemori R, Hanamura T, Morikawa H. Scripta Metall Mater 1992;26:969. [22] Menand A, Huguet S, Nerac-partaix A. Acta Mater 1996;44:4729. [23] Kobayashi Y, Tsukihashi F. Metall Trans B 1998;29:1037. [24] Schuster JC, Nowotny H, Vaccaro C. J Solid State Chem 1980;32:213. [25] Morris MA, Leboeuf M. Mater Sci Eng A 1997;224:1. [26] Greenberg BA, Antonova OV, Indenbaum VN, Karkina LE, Notkin AB, Ponomarev MV et al. Acta Mater 1991;39:233. [27] Court SA, Vasudevan VK, Fraser HL. Phil Mag A 1990;61:141. [28] Huang SC, Hall EL. Metall Trans A 1991;22:427. [29] Sastray SML, Lipsitt HA. Metall Trans A 1977;8:299. [30] Inui H, Kishida K, Misakiandetal M. Phil Mag A 1995;72:1609. [31] Vasudevan VK, Stucke MA, Court SA, Fraser HL. Phil Mag Lett 1989;59:299. [32] Morris MA, Lipe T. Intermetallics 1997;5:329. [33] Morris MA. Phil Mag A 1993;68:259. [34] Morris MA, Morris DG. Intermetallics 1999;7:1069. [35] Aindow M, Chaudhuri K, Das S, Fraser HL. Scripta Metall Mater 1990;24:1105. [36] Vasudevan VK, Court SA, Kurath P, Fraser HL. Scripta Metall Mater 1989;23:467. [37] Morris MA. Intermetallics 1996;4:417. [38] Morris DG, Dadras MM, Morris-Munoz MA. Intermetallics 1999;7:589. [39] Kad BK, Fraser HL. Phil Mag Lett 1994;70:211. [40] Becker S, Rahmel A, Schorr M, Schutze M. Oxid Met 1992;38:425.
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[41] Zhang TH, Xie JD, Ji CZ, Chen J, Xu H, Li J et al. Surf Coat Technol 1995;72:93. [42] Bradey M, Brindley WJ, Humphrey DL, Smith J. Acta Mater 1997;45:2371. [43] Kehare SA, Aswath PB. J Mater Sci 1997;32:2485. [44] Yoshihara M, Kim YW. In: Kim YW, Dimiduk DM, Loretto MH, editors. Gamma titanium aluminides. San Diego
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Microstructural refinement and improvement of mechanical properties and oxidation resistance in EPM TiAl-based intermetallics with yttrium addition Y. Wu a, b, S.K. Hwang a, c,* a Inha University, School of Materials Science and Engineering, Incheon 402-751, South Korea Beijing General Research Institute of Mining and Metallurgy, Beijing 100044, People’s Republic of China Jointly appointed by the Center for Advanced Aerospace Materials, Pohang University of Science and Technology Pohang, South Korea b
c
Received 29 August 2001; received in revised form 7 December 2001; accepted 7 December 2001
Abstract New Y-modified TiAl-based alloys, with nominal chemical compositions of Ti–46.6Al–1.4Mn–2Mo–0.3C–xY (x ⫽ 0.1, 0.33 and 0.6), were developed with the elemental powder metallurgy (EPM) method. The alloys with Y addition had higher ultimate tensile strength, elongation and oxidation resistance than the Y-free alloy. Combined solidsolution strengthening of Mn, Mo, C and Y, and the microstructural refining effect of C and Y, mainly contributed to the strength at room temperature as well as at high temperature up to 800°C, in addition to strengthening due to Y2O3 and carbide precipitation. The room-temperature ductility was also improved by the microstructural refinement and the removal of interstitial oxygen in the α2 and γ phases. Plastic deformation of γ plates during tensile testing occurred through the activation of ordinary dislocations of perfect Burger’s vectors. In an oxidation test in air, the Y-containing alloys showed a significantly low weight gain compared with the Y-free alloy. 2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Titanium; Microstructure; Yttrium; Mechanical properties; Oxidation
1. Introduction In the last decade, intermetallic compounds based on TiAl have drawn increasing attention for aerospace structural applications in the intermediate temperature range (600–800°C) due to their
* Corresponding author. Tel.: +82-3286-07537; fax: +823286-25546. E-mail address: [email protected] (S.K. Hwang).
low density of 3.8 g/cm3 and good strength at elevated temperatures [1,2]. Recent research efforts have focused on chemical alloying and thermomechanical processing in order to balance the overall tensile properties at room temperature and at high temperatures [3,4]. Concerning the effects of alloying additions, Mn addition and C addition benefit the room-temperature ductility and the creep resistance, respectively; Mo addition increases the strength as well as the creep and oxidation resistance of the TiAl alloys [5–9]. Mean-
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while, Hwang and co-workers [10–15] have developed Ti–46.6Al–1.4Mn–2Mo–xC alloys of improved high-temperature tensile strength with the elemental powder metallurgy (EPM) method. The authors of these works attributed the improved mechanical properties to the microstructural refining effect, especially the lamellar refinement, of carbon. However, the room-temperature ductility of C-modified TiAl is still rather low, an elongation of approximately 1.6%. It was thought that the high oxygen content, approximately 1500 wt ppm, inherent to the EPM processing, was detrimental to ductility in these alloys. Tonnes et al. [16] reported that an increase in the oxygen content from 1000 wt ppm to 1600 wt ppm resulted in a sharp ductility drop from 2.1% to 0.5% in TiAl. In this respect, Y has an interesting role to play because this element, as an alloying element, is known to enhance the mechanical properties of other alloys [17]. The purpose of present work was to investigate the effect of Y addition on the tensile properties and oxidation resistance of TiAl–Mn–Mo–C intermetallics. Whether Y can scavenge the interstitial oxygen without interfering with the beneficial effects of other alloying elements was the question of primary interest.
2. Experimental The materials investigated for the present work were fabricated by EPM, the details of which are described elsewhere [10–15,18]. Of particular importance are the processing parameters such as powder degassing [10,11], extrusion can design and extrusion ratio [10], heating rate and holding time for the soaking treatment [11,18]. Specimens used for the present study were obtained with the proper control of these variables, thereby chemical homogeneity as well as microstructural uniformity were ensured. The mean size of powder particles of Ti, Al, Mn, Mo, C and Y ranged from 7 µm to 74 µm in experimental alloys. The nominal chemical compositions of the experimental alloys were Ti–46.6Al–1.4Mn–2Mo–0.3C–xY (x ⫽ 0.1, 0.33 and 0.6) (at%). Powders were mixed in a Vblender, compacted into Cr–Mo steel cans and
sealed. The cans were degassed sequentially at room temperature, 250°C and 600°C for 1, 2 and 4 h, respectively, followed by vacuum sealing at 3 × 10⫺5 torr. Degassed cans were hot-extruded at 1250°C with a ratio of 10:1. Prior to the hot extrusion, the cans were pre-heated at the same temperature and held for 12 h. After removal of the steel cans by machining, the specimens were heat-treated in an argon atmosphere. The heat treatment cycles consisted of a solution treatment at 1400°C for 1 h followed by air cooling (AC) and an aging treatment at 800°C. Transmission electron microscopy (TEM) specimens were prepared by standard jet-polishing techniques using a solution of 5% perchloric acid, 30% butanol and 65% methanol at ⫺25°C and 40 mA. Microstructural observation of the specimens was performed by optical microscopy (OM) and an H800 transmission electron microscope (TEM) operated at 200 kV. Quantitative measurements of the lamellar spacing were carried out using an image analyzer attached to the TEM. Chemical compositions of phases and precipitates in the specimens were analyzed by energy-dispersive X-ray spectrometry (EDS) in the TEM. Microhardness of the specimens was measured by a Vickers hardness tester using a load of 100 kg. Tensile specimens were 19 mm in gage length and 4 mm in gage diameter, and tested with a strain rate of 1 × 10⫺3 s⫺1. At least three samples were tested under identical conditions to obtain an average. The oxidation resistance of alloys was tested by isothermal exposure in air at 800°C for up to 350 h. Prior to the exposure, the specimens, 146 mm2 in surface area, were polished with emery papers of up to 1000 mesh finish. A thermo-balance was used to monitor the weight gain.
3. Results 3.1. Effect of Y on the microstructure There were significant differences between the microstructure of the Y-free alloy and that of the Y-added one. Fig. 1 shows optical micrographs of Ti–46.6Al–1.4Mn–2Mo–0.3C–xY alloys solutiontreated at 1400°C for 1 h and air-cooled. Regard-
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Fig. 1. Optical micrographs showing the fully lamellar microstructures of EPM Ti–46.6Al–1.4Mn–2Mo–0.3C alloys with different Y additions solution-treated at 1400°C /1 h / AC: (a) base composition, (b) 0.1Y, (c) 0.33Y and (d) 0.6Y. Each grain consists of colonies of aligned α2 ⫹ γ platelets.
less of Y content, the microstructures of experimental alloys showed a fully lamellar (FL) feature that consisted of colonies of directionally aligned α2 ⫹ γ platelets. A strong grain-refining effect of Y was observed as shown in Fig. 2: addition of 0.33Y to the base composition of Ti–46.6Al– 1.4Mn–2Mo–0.3C reduced the average grain size from 100 µm to approximately 45 µm and further addition of Y (0.6%) decreased it to approximately 30 µm. Y addition also refined the lamellar thickness of the TiAl alloy. An example of the refined lamellar structure is shown in Fig. 3, which compares the α2 ⫹ γ lamellar microstructures of Y-free and 0.33Y-added alloys, each solution-treated at 1400°C for 1 h/AC and aged at 800°C/12 h/AC.
Although Y addition refined the lamellar thickness it did not alter the flat feature of the lamellar interfaces, indicating that the individual morphology of γ and α2 plates was preserved. Twin-related α2 and γ plates were frequently observed, an example of which is shown in Fig. 3(b). In this figure, twinrelated γ-TiAl and α2-Ti3Al plates are shown with the zones of 具110]γ and [112¯ 0]α2, respectively. From the selected-area diffraction pattern (SADP) shown in Fig. 3(c), it was confirmed that the plate interfaces of γ phase and α2 phase were (111) and (0002), respectively. Among the γ plates there were twin relationships. With the rotation axis of [111], the twin plates of γ showed three variants: 120°, 180° and a pseudo-twin relationship. Orientation relationships of (111)γ储(0002)α2 and 具110]γ储
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Fig. 2. Grain-refining effect of Y in EPM Ti–46.6Al–1.4Mn– 2Mo–0.3C alloys α-solution heat-treated at 1400°C / 1 h /AC.
[112¯ 0]α2 were obtained, which are consistent with the Blackburn [19] relationships. Based on the TEM observation, a quantitative relationship between the average thickness of lamellae and Y content was obtained and the result is presented in Fig. 4. The measurement was carried out with an image analyzer attached to the TEM in the zone of 具110]γ储[112¯ 0]α2. At least 200 plates were measured for each condition to ensure reliability of the data. The average thickness of the γ plate decreased from 64 nm to 50 nm by the addition of 0.33% Y to the alloy with the base composition [Fig. 4(a)], while the thickness of the α2 plate also decreased from 28 nm to 22 nm [Fig. 4(b)]. A detailed TEM–EDS analysis showed a slight partitioning of Y in the γ phase. Measurements over 50 spots each for the α2 and γ plates resulted in the average concentration of Y of 0.41 wt% and 0.31 wt% and in the γ and α2 phase, respectively, with a scattering of ± 0.03 wt%. Although this analysis could not separate the contributions of the oxide particles from that of the matrix, it is interesting to note that an earlier study reported a preferential partitioning of oxygen in the α2 phase [20– 22]. In selective terms, this means that the oxygen content in the matrix of the γ plate will be substantially lower than that in the α2 plate due to the oxygen-scavenging effect of Y.
Subsequent extrusion and heat treatment, particularly the solution heat treatment, provide the oxide-forming environments. The relative amount of oxides formed in each stage is not known, although it is surmised that the partitioning of Y and the limited solubility of Y and O favor oxide formation in the γ phase. Due to the strong binding among Y and O atoms [23], the oxides of Y are thought to occur starting from the pre-extrusion heat treatment. Morphological characteristics and the chemical composition of oxide particles, Y2O3, were studied. As shown in Fig. 5(a), the oxide particles in the γ phase of Ti–46.6Al–1.4Mn–2Mo– 0.3C–0.33Y alloy aged at 800°C for 12 h showed an irregularly elongated shape that aligned along crystallographic directions. Oxide particles formed not only homogeneously within α2 and γ plates but also heterogeneously along the α2 / γ boundaries [Fig. 5(b)]. The particles had a hexagonal crystal structure, the lattice parameters determined from SADP being a ⫽ 0.376 nm and c ⫽ 0.582 nm [Fig. 5(c)]. The orientation relationships between the γ matrix and the precipitate were {111}γ储{0001}Y2O3 and 具011典γ储具112¯ 0典Y2O3. As expected, the EDS result revealed that precipitates were enriched with Y and oxygen [Fig. 5(d)]. Beside oxide particles, the aged Ti–46.6Al– 1.4Mn–2Mo–0.3C–0.33Y alloy also showed carbides, which ranged from 40 nm to 300 nm in size, and formed heterogeneously along α2 / γ grain boundaries and on the existing dislocations as shown in Fig. 6. EDS analysis in the TEM showed that the plate-like precipitates in Fig. 6(a) contained mainly Ti and Al, with a ratio of Ti/Al of approximately 3:1, and a small amount of Mn, Mo, C and Y. The crystal structure of the carbide was of the perovskite type (P phase), which was consistent with the earlier observation by Tian and Nemoto [9]. Regardless of their locations, cubictype orientation relationships between Ti3AlC and γ-TiAl phase were found, namely (001)p储(001)γ and [001]p储[001]γ. This result indicates that the crystallography of the precipitates was governed by the crystallography of the growing γ plates. On the other hand, the precipitates shown in Fig. 6(b) were determined to be carbide of Ti2AlC (H phase) type that had an ordered hexagonal crystal structure. The H phase also had strict orientation
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Fig. 3. TEM micrographs of EPM Ti–46.6Al–1.4Mn–2Mo–0.3C alloy with different additions of Y showing the lamellar-refining effect of Y: (a) Y-free alloy, (b) 0.33Y-added alloy and (c) selected-area diffraction pattern (SADP) taken from the area of α2 ⫹ γ plates in (b) with the zone axis of [1¯ 10]γ. The alloys were heat-treated at 1400°C / 1 h /AC ⫹ 800°C / 12 h / AC.
relationships with the matrix of γ plates, which coincided with the results of Tian and Nemoto [9]. According to the Ti–Al–C ternary phase diagram reported by Schuster et al. [24], Ti3AlC is in equilibrium with TiC, α2-Ti3Al and α-Ti, whereas Ti2AlC is in equilibrium with TiC, α2-Ti3Al and γ-TiAl. 3.2. Mechanical properties Evaluation of the mechanical properties of Yadded TiAl-based alloys was made on the basis of hardness and tensile properties. Regardless of heat treatment conditions, the hardness of the experimental alloys increased monotonically over the whole range (0.1 to 0.6%) of Y content studied.
The hardness increase was particularly notable for the Y content of 0.1 to 0.33%: it increased from 454Hv to 513Hv in the aged condition. As expected, the aged alloys showed higher hardness than the as-extruded or as-solution-treated alloys due to extensive precipitation of carbides. Y addition also improved the tensile strength of experimental alloys. The values of tensile strength and elongation of aged Ti–46.6Al–1.4Mn–2Mo– 0.3C–xY alloys measured at room temperature and 800°C are shown in Fig. 7 as a function of the Y content. The room-temperature ultimate tensile strength increased steeply with Y content up to 0.33% and then became saturated up to 0.6%, resulting in an overall increase of 20% compared with the Y-free alloy. This trend was repeated in
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Fig. 4. Quantitative measurement of the effect of Y on the lamellar thickness of EPM Ti–46.6Al–1.4Mn–2Mo–0.3C alloy: (a) γ plate and (b) α2 plate. The alloys were heat-treated at 1400°C / 1 h / AC ⫹ 800°C /12 h / AC.
the 800°C test although there was a gap of about 40 MPa between the strength at room temperature and that at the high temperature. Therefore the results of tensile testing were consistent with those of the hardness test. Because of the effect of Y, the room-temperature ultimate tensile strength reached the highest value of 704 MPa in the alloy containing 0.6% Y. Y addition also affected the ductility of experimental alloys. As shown in Fig. 7(b), a room-tem-
perature tensile elongation ranging from 1.6% to 3.3% was obtained for the experimental alloys of fully lamellar microstructure. Increasing the Y content from 0 to 0.1% promoted an increase in tensile elongation from 1.6% to 3.3%. However, excessive Y addition deteriorated ductility: a room-temperature elongation of 2.4% was obtained for the alloy containing 0.6% Y. The tensile ductility was generally higher at 800°C than the room-temperature values, the gap between the two values broadening with the Y content. Improved ductility in the Y-containing alloys was also evidenced by scanning electron microscopy (SEM) fractographs. In Fig. 8, two sets of fractographs are presented: one from a tensile specimen tested at room temperature [Fig. 8(a) and (b)] and the other from a test at 800°C [Fig. 8(c) and (d)]. Transgranular cleavage features predominated the room-temperature test specimen whereas a mixture of cleavage and intergranular fracture were apparent in the 800°C test specimen. In both cases addition of Y changed the fracture characteristic towards a more ductile mode, converting the brittle fracture to a plastic tear type of separation. The fracture surface of specimens tested at 800°C reflected another effect of Y addition in that the extent of oxidation was reduced. As inferred from the comparison of Fig. 8(c) with (d), the fracture surface of the Y-free alloy revealed a significant oxide layer whereas that of the Y-containing alloy exhibited a rather clean feature. This was an indication of an enhanced oxidation resistance by Y addition, which led to the next phase of investigation. 3.3. Deformed microstructure The slip mode of Y-free TiAl alloy was studied by examining the Burger’s vectors of dislocations in tensile specimens. In Fig. 9, bright-field TEM images of dislocations in a deformed specimen are shown for six different two-beam conditions. There were two different geometries of dislocations, leveled and tilted. Leveled dislocations were visible under all tilt conditions except for g ⫽ 22¯ 0 [Fig. 9(f)], which proved that the Burger’s vector of the dislocations was ordinary 1/2具110]. On the other hand, the Burger’s vector of tilted dislocations was
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Fig. 5. TEM micrographs showing Y2O3 formed in EPM Ti–46.6Al–1.4Mn–2Mo–0.3C alloy heat-treated at 1400°C /1 h / AC ⫹ 800°C /12 h / AC. (a) Y2O3 precipitates formed in a γ grain, (b) Y2O3 precipitates formed along the α2 / γ boundaries, (c) SADPs of the region shown in (a), from which the orientation relationship between the Y2O3 precipitates and the γ-TiAl matrix was determined, and (d) the EDS analysis of the oxides in (b).
super 具101典, some of which were dissociated. Therefore the γ grain deformed through the movement of mainly superdislocations of 具101典 type and some 1/2具110] ordinary dislocations. Compared with the Y-free alloy, dislocations in deformed specimens of Y-added alloys showed different characteristics. A set of TEM micrographs, shown in Fig. 10, was taken from the γ plate of a tensile specimen of 0.33 Y-containing alloy tested at room temperature, for which the Burger’s vectors of dislocations were determined. Like the case for the Y-free alloy (Fig. 9), two groups of dislocations were identified, leveled and tilted. From this analysis it was determined that both types of dislocation were of 1/2具110] Burger’s
vector. These dislocations were mostly straight and aligned along the 具1¯ 10] direction to form screw dislocations. Additionally, there were some cusps and loops along a dislocation line, hinting active cross-slips. Predominance of the ordinary dislocations in Ycontaining alloy was also evident in specimens tensile tested at 800°C. A typical dislocation structure in the γ plate of the alloy containing 0.33Y is shown in Fig. 11. With increase in the tensile test temperature, dislocations became gradually curved compared with those in room-temperature tensile test specimens. Through a series of diffraction contrast analyses, it was concluded that the majority of dislocations in the γ plate were still of 1/2具110]
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Fig. 6. TEM micrographs showing the carbide precipitates (a) along the α2 /γ boundaries and (b) on dislocations within the γ grain of EPM Ti–46.6Al–1.4Mn–2Mo–0.3C–0.33Y alloy aged at 800°C for 12 h. The P phase, Ti3AlC in chemical composition, has a perovskite structure with carbon atoms occupying the octahedral sites of L10 structure. The H phase, Ti2AlC in composition, has an ordered hexagonal structure.
type. Therefore increasing the deformation temperature from room temperature to 800°C promoted cross-slip activities but did not alter the preferential activation of ordinary 1/2具110] dislocations instead of superdislocations.
Fig. 7. Effect of Y content on the tensile properties of EPM Ti–46.6Al–1.4Mn–2Mo–0.3C–xY alloys: (a) tensile strength and (b) elongation. Samples were heat-treated at 1400°C / 1 h / AC ⫹ 800°C / 12 h / AC and tensile-tested at room temperature as well as at 800°C.
3.4. Oxidation resistance of Y-containing TiAl alloys A series of oxidation tests performed in air at 800°C demonstrated that Y addition improved the oxidation resistance of Ti–46.6Al–1.4Mn–2Mo– 0.3C–xY alloys. Shown in Fig. 12 are the curves of mass gains per unit surface area during the oxidation tests, which lasted for up to 350 h. The oxidation kinetic curves consisted of two stages. In
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Fig. 8. SEM fractographs of fractured tensile specimens of EPM Ti–46.6Al–1.4Mn–2Mo–0.3C–xY alloys: (a) base composition and (b) 0.1Y alloy tested at room temperature, and (c) base composition and (d) 0.1Y alloy tested at 800°C.
the first stage of oxidation ( ⬍ 25 h) mass gain increased linearly with time, which was followed by the second stage, in which a parabolic dependence on time was realized. After oxidation for 350 h, the mass gains for Y-free, 0.1Y-, 0.33Y- and 0.6Y-added alloys were 29.0, 19.7, 14.6 and 10.3 g/m2, respectively, reflecting a significant improvement in the oxidation resistance by Y addition.
4. Discussion As reported earlier by Morris and Leboeuf [25], grain refinement is an effective method to improve the tensile properties of TiAl-based alloys. In a previous study Hwang et al. showed that C
addition was beneficial for this purpose [12,15]. In the present study, Y was shown to have an additional effect on refining grains and lamellae. The grain-refining effect of Y was attributed to the precipitation of oxides that formed prior to extrusion. During extrusion and subsequent heat treatment, the oxides provide ample sites of heterogeneous nucleation for the product phases (α2 and γ from high-temperature α) while restraining their growth rate. In this respect, Y2O3 must be stable at 1400°C, the solution heat treatment temperature. According to the thermodynamical data given by Kobayashi and Tsukihashi [23], ⌬Gf° ⫽ 794,000⫺348T (J/mol, T in K) for the decomposition of Y2O3, and it is estimated that the oxide is stable up to 2009°C in TiAl.
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Fig. 9. TEM micrographs showing dislocation structures in a γ plate in EPM Ti–46.6Al–1.4Mn–2Mo–0.3C alloy under different diffraction conditions: (a) g ⫽ 220, (b) g ⫽ 02¯ 0, (c) g ⫽ 200, (d) g ⫽ 1¯ 31¯ , (e) g ⫽ 3¯ 11 and (f) g ⫽ 22¯ 0. Samples were heat-treated at 1400°C / 1 h /AC ⫹ 800°C / 12 h / AC and tensile-tested at room temperature.
Improvement of the tensile strength by Y addition, in addition to the strengthening effect by the perovskite carbides, is a natural consequence of the microstructural refinement. As for the ductility improvement, however, there are additional contributions by Y: first, an oxygen-scavenging effect and, second, the change of the slip mode. Y has a strong affinity to oxygen and, as such, it tends to
absorb the interstitial oxygen atoms from the matrix to form oxides, Y2O3 [17]. Since the interstitial oxygen is detrimental to ductility, this enhances the room-temperature elongation [16]. In general, the dislocation structure in L10-TiAl is complex in that there are 1/2具110] ordinary dislocations, two different types of superdislocation with Burger’s vector of 具011] and 1/2具112],
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Fig. 10. TEM micrographs showing dislocation structures in a γ plate in EPM Ti–46.6Al–1.4Mn–2Mo–0.3C–0.33Y alloy under different diffraction conditions: (a) g ⫽ 1¯ 11, (b) g ⫽ 002, (c) g ⫽ 1¯ 11¯ , (d) g ⫽ 1¯ 1¯ 1. Samples were heat-treated at 1400°C / 1 h / AC ⫹ 800°C /12 h / AC and tensile-tested at room temperature.
respectively, and 1/6具112] twin dislocations. A variety of microdeformation mechanisms in TiAl have been reported in the literature [26–34]. Huang and Hall [28] and Vasudevan et al. [31] found that 1/2具110] slip was the predominant deformation mode in two-phase TiAl alloy, which was also supported by Morris and Lipe in alloys based on Ti– 48Al–2Mn–2Nb [32]. For Ti–48Al alloy with high oxygen content, it was reported that deformation at high temperatures occurred through twinning instead of ordinary dislocations [33,34]. The present studies indicate that increasing concentrations of yttrium have a marked effect on the mobility of 1/2具110] ordinary dislocations regardless of deformation temperature, see Figs. 10 and 11. Thus, if the yttrium concentration is increased,
plasticity of the TiAl is enhanced most readily through the activity of 1/2具110] ordinary dislocations. Superdislocations appear to play a major role in plastic deformation only in the case of Yfree alloy, in which 1/2具110] ordinary dislocations have reduced mobility due to increased oxygen atoms in solution. These results are in concert with the assertion of Aindow et al. [35] and Vasudevan et al. [36], where the increased mobility of 1/2具110] ordinary dislocations in the Er-added TiAl alloys was attributed to the internal gettering of oxygen from solution by formation of Er2O3. Promotion of the activity of ordinary dislocations by Y was attributed to the counteraction to the adverse effect of interstitial elements such as C and O. Previous studies made by Morris et al.
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Fig. 11. TEM micrographs showing dislocation structures in a γ plate in EPM Ti–46.6Al–1.4Mn–2Mo–0.3C–0.33Y under different diffraction conditions: (a) g ⫽ 11¯ 0, (b) g ⫽ 1¯ 11¯ , (c) g ⫽ 1¯ 11, (d) g ⫽ 1¯ 13. Samples were heat-treated at 1400°C / 1 h / AC ⫹ 800°C /12 h / AC and tensile-tested at 800°C.
[37,38] also showed that ordinary dislocations were immobilized in the alloys with high oxygen concentrations. In this respect, the solubility of oxygen needs to be considered. The maximum solubility of oxygen in TiAl phase is 300 ± 150 at ppm, while the oxygen concentration in the experimental alloys was approximately 1500 wt ppm. Therefore it is deduced that a significant portion of the excessive oxygen atoms will be heterogeneously distributed by segregating to dislocations. In the O-rich alloy, it is believed that interstitial oxygen in solid solution is responsible for strong dislocation pinning. This may, of course, apply both to ordinary and to superdislocations. It is difficult to distinguish the oxygen atoms segregated at 1/2具110] ordinary dislocations from those
segregated at superdislocations. The results of observation by the present authors as well as others only hint that the oxygen concentration trapped in the 1/2具110] ordinary dislocations was higher than that in the superdislocations. The question of how oxygen may affect dislocation movement and why 1/2具110] ordinary dislocations may be more affected than 具101典 dislocations have been discussed by Kad and Fraser [39], but remains poorly understood. The higher lattice resistance due to the presence of the oxygen atoms seems to increase the difficulty of motion of ordinary dislocations to the extent that superdislocations are activated instead. In addition, there is a possibility that Y addition decreases the anti-phase boundary (APB) energy of superdislocations, making them more immobile.
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Fig. 12. Effect of Y on the oxidation kinetics of EPM Ti– 46.6Al–1.4Mn–2Mo–0.3C–xY alloys in air at 800°C. Samples were heat-treated at 1400°C /1 h / AC ⫹ 800°C /12 h / AC.
In summary, the Y addition affected the mechanical properties through multiple routes: grain refinement, lamellae refinement, interstitial scavenging, and promoting ordinary dislocation activity. However, it is noteworthy that there existed an optimum range of Y addition, 0.1 to 0.3%, since excessive addition causes embrittlement. The influence of alloying additions on isothermal oxidation behavior of TiAl-based alloys has been investigated by many researchers [40– 44]. A beneficial effect of Y on the oxidation resistance of various alloy systems has been reported [45–48]. As shown in Fig. 12, the alloys containing Y showed a significantly improved resistance to high-temperature oxidation in air. The alloys with Y addition showed smaller mass gains and smaller parabolic rate constants, which is in line with the previous studies mentioned above. It is poorly understood, however, how Y improves the oxidation resistance. There is a good possibility that Y modified the phase constitution and the microstructure of the oxide. Zhang et al. [41] studied the behaviors of the oxidation resistance of a H13 steel implanted with Y or Y plus carbon, and concluded that Y provided a barrier layer that hindered oxygen diffusion during thermal oxidation that occurs primarily along dislocations or grain boundaries. Han and Xiao [45] investigated the
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mechanism of improvement of the high-temperature isothermal oxidation resistance of Ni3Al-based alloys with added Y. They demonstrated that the beneficial effects of Y on high-temperature oxidation resistance. These results, in view of the present observation, suggest that Y addition suppresses the growth of TiO2 by promoting formation of Al2O3 scale. Although direct evidence to support this suggestion is lacking at present, this explanation is plausible for the present system based on TiAl. In the parabolic stage, oxidation in all experimental alloys obeyed the parabolic rate law that was represented by the equation (⌬M)2 ⫽ Kpt, where ⌬M, Kp and t are mass gain, parabolic rate constant and time, respectively. This process was controlled by cation transport through the scale. Because TiO2 oxides grow mainly by inward diffusion of oxygen via a vacancy mechanism [44], the improvement by Y addition in high-temperature oxidation behavior could be explained by its ability to reduce the oxygen vacancy concentration in the TiO2 scale due to the smaller oxide precipitates induced by Y [45–47]. Owing to their strong field strength with oxygen atoms, Y atoms are considered to reduce the oxygen vacancy concentration in TiO2 scale. The results presented in Fig. 12 indicate that this effect is proportional to the Y content, reaching a maximum at 0.6 at%. In addition, the improvement in oxidation resistance of the Y-added alloy could be interpreted by the classical Wagner theory [44], which explains the oxidation behavior in terms of oxygen solubility in the alloy. According to this interpretation, Y enhances the oxidation resistance by lowering the oxygen solubility through the formation of Y2O3 phase.
5. Conclusions From the study of the effect of Y on the mechanical properties and the oxidation resistance of a TiAl-based alloy, the following conclusions were drawn. 1. Partitioning of Y in the lamellar structure was found and quantitatively determined: 0.31 wt%
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Y and 0.41 wt% Y in the α2 and γ phases, respectively. 2. The orientation relationship between the γ matrix and the Y2O3 precipitate was determined and by TEM: {111}γ储{0001}Y2O3 具011典γ储具112¯ 0典Y2O3. 3. A strong grain-refinement effect of Y was found. As a result of the grain refinement, combined with the oxides and carbides precipitation and solid-solution hardening, the tensile strength of the experimental alloys improved significantly with Y addition at test temperatures of up to 800°C. The strength enhancement was accompanied by a concomitant increase in the ductility in the alloys with a moderate amount of Y addition, which was primarily due to the interstitial oxygen scavenging and enhanced activity of 1/2具110] ordinary dislocations. 4. Addition of Y also resulted in a substantial improvement in the oxidation resistance of the intermetallic compound exposed to air at 800°C. Acknowledgements The authors would like to thank Dr H.S. Ryoo, Mr Y.W. Park and Miss J.S. Park for their kind help. Y. Wu also acknowledges Dr W. Liang for TEM experiments in Taiyuan University of Technology. This work was performed under the auspices of the Korean Science and Engineering Fund through the Engineering Research Center program at POSTECH-CAAM, and by the Korean Ministry of Science and Technology through the “2001 National Research Laboratory Program”. References [1] [2] [3] [4] [5]
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