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Microstructure and mechanical properties of the hypercooled Ti50Al50 alloy
Pergamon
Materials Research Bulletin 36 (2001) 963–969
Microstructure and mechanical properties of the hypercooled Ti50Al50 alloy Y.C. Liua,*, G.C. Yanga, X.F. Guoa, J.H. Yanga, D.S. Xub, Y.H. Zhoua a
State Key Lab of Solidification Processing, Northwestern Polytechnical University, Xi’an 710072, P. R. China b State Key Lab of Laser Technology, Central University of Sci. & Tech., Wuhan 430074, P. R. China (Refereed) Received 24 July 2000; accepted 14 January 2001
Abstract Hypercooling region of the gamma titanium aluminide alloy Ti50Al50 was achieved by cyclically superheating the alloy melt in a containerless electromagnetic levitation apparatus. The maximum undercooling of the alloy melt amounted to 345 K. XRD, TEM, SEM and optical microscopy techniques were adopted to investigate the microstructure and identify the phase composition. The microstructure of the hypercooled alloy was composed of equaxied gamma phase grains with a mean diameter of 17 nm. The fine microstructure of the gamma phases produced a large strengthening effect, raising the hardness to 545 on the Vickers scale. The aluminum composition distribution in the hypercooled sample was homogenous except for a small difference between first solidified regions and boundaries of cracks. X-ray diffraction results showed that all peaks of the gamma phase shifted slightly in the direction of small angles. This can be explained by the disordering growth pattern caused by the rapid solidification process in the hypercooled melt. © 2001 Elsevier Science Ltd. All rights reserved. Keywords: A. Intermetallic compounds; C. X-ray diffraction; D. Microstructure; D. Mechanical properties
1. Introduction Lightweight gamma titanium aluminide alloys are potential candidate materials in automotive and aerospace industries for high temperature structural applications [1]. However, * Corresponding author. Present address: Max-Planck-Institute for Metals Research, Seestr. 92, D-70174 Stuttgart, Germany. Tel.: ⫹49-711-2095427; fax: ⫹49-711-2095418. E-mail address: [email protected] (Y.C. Liu). 0025-5408/01/$ – see front matter © 2001 Elsevier Science Ltd. All rights reserved. PII: S 0 0 2 5 - 5 4 0 8 ( 0 1 ) 0 0 5 7 7 - 3
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their poor ductility and low fracture toughness at ambient temperature are the major limitations for practical applications [2]. Many efforts to introduce grain boundary deformation mode and suppress crack propagation have been made to improve their ductility and fracture toughness by significantly decreasing the grain size of the material [3–5]. For example, magnetron sputtering and condensation [6] and also mechanical alloying [7] have been used successfully to produce the nanophase of gamma titanium aluminide alloy. Other methods commonly used to refine the grains in the development of high performance materials are the rapid quenching process and thermodynamic undercooling. Due to volume limitation, the rapid quenching process tends to be used only on the laboratory scale. Liquid undercooling is known to significantly affect the solidification process, resulting in changes in the microstructure development and bulky metastable phase formation [8]. Several secondary processes, such as dendritic remelting, structural transformation, and segregation, generally affect solidification at low undercoolings, resulting in a variety of microstructural patterns. In this regard, attainment of hypercooling is of special interest. In recent years, investigations on high melting bulk (Co, Ni, Fe)-Pd and Au-(Cu, Ni) alloys had demonstrated experimentally the achievement of the hypercooling region [9 –11]. Undercooling experiments had been performed on alloy Ti50Al50 involving intermetallic compounds, and a maximum undercooling of 300 K was achieved [12]. Different, even conflicting results about the formation of primary phases in the melt at the lower undercooling have been reported. Similar results had been reported for the highly undercooled region of the melt [12,13].
2. Experimental 2.1. Ingot preparation High-purity (⬎ 99.99%) aluminum and titanium were used in the appropriate amounts to form an alloy with a stoichiometry of Ti50Al50. The process of melting was carried out in a vacuum arc furnace under a high-purity argon atmosphere, to produce a button, approximately 4 cm in diameter, which was remelted four times. The arc-melted button was homogenized in a vacuum heat-treatment furnace at 1273 K for 80 h before undercooling experiments were carried out, in order to achieve complete composition homogenization. 2.2. Undercooling experiments The undercooling experiments were carried out in an electromagnetic levitation melting apparatus manufactured by Edmuld Buhler Co. The device possesses a maximum levitation capacity of 200 mg and a HF-3 high-frequency generator is used to heat the alloy. In the undercooling process, the working chamber was initially evacuated to about 10⫺9 mbar, and then back-filled with high-purity argon gas (⬎ 99.9999%). For the purpose of deactivating heterogeneous nucleation sites, each sample, approximately 50 mg, was cyclically superheated at 300 K for 5 min, then the power was turned off, the superheated sample
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was allowed to cool on an inert zirconia sample holder, which was stretched into the coil in advance. When a desired undercooling of the alloy melt was reached, the sample was cooled naturally to the ambient temperature on the holder. The cooling rate of the process was approximately in the magnitude of 102⬃103 K/s. An infrared pyrometer with an absolute accuracy, relative accuracy, and response time of less than 10 K, 3 K and 5 ms, respectively, was used to record the thermal history of samples [14]. The cooling curve was calibrated with a standard PtRh30-PtRh6 thermal couple, which was encapsulated in a silica tube and then immersed into the melt. The melting temperature and undercooling degree in the cooling curve could be obtained by comparing the absolute temperature recorded by the standard thermal couple. 2.3. Characterization Crystallographic features were mainly examined using Rigaku X-ray power diffractometer with a Cu K␣ source. The slice for transmission microscopy analysis was electrolytically thinned using an electrolyte of 5% HClO4 in ethanol at 243 K and investigated using a JEOL JEM-200cx electron microscope. Microstructure observation was carried out using a JEOL JXA-840 scanning electron microscope. Kroll’s etching regent and a Nikon Epiphot optical microscopy were used for observing the microstructure. Hardness measurements of the polished specimens were carried out on a Vickers hardness indentor. The indentation load and time were 50 g and 10 s, respectively.
3. Experimental results and discussion According to the binary phase diagram of Ti-Al [15], the composition of alloy Al50Ti50 falls in the L3␣ equilibrium solidification region and the equilibrium microstructure of it is in the two-phase (␣2⫹␥) region. Fig. 1a shows the optical microscope image of the arc-melted alloy homogenized at 1273 K for 80 h. The structure is composed of equaxied grains and the lamellar structure could be made out in grains. Fig. 1b shows a more detailed image of the iterative white and black lamella (␣2⫹␥) phases, taken with the transmission electron microscope. The hypercooling regime is achieved when the enthalpy of a liquid is reduced with decreasing temperature to a level equal to that of the solid phase at the melting temperature, Tm. The isenthalpic temperature, Ti, which represents the upper limit of the hypercooling regime, is defined by ⌬H f ⫽
冕
Tm
C pl共T兲dT
Ti
with ⌬Hf the enthalpy of fusion and cpl(T) the specific heat of the undercooled liquid. The enthalpy of fusion of Ti50Al50 can be derived from the calculation based on the sublattice
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Fig. 1. Microstructures of the arc-melted Ti50Al50 alloy homogenized at 1273 K for 80 hours: Optical image and bright field image of lamellar structure.
model [16]. Then, the minimum of undercooling to achieve the hypercooled state of Ti50Al50 can be calculated from ⌬T k ⫽
⌬H m ⫹ 共T l ⫺ T s兲 Cp
where, Tl is the liquid temperature, and Ts, the solid temperature. The minimum calculated undercooling for achieving the hypercooling state is 340 K. Therefore, the sample of Ti50Al50 alloy melt with an undercooling above 340 K should solidify as a hypercooled mode. Fig. 2 gives the optical and scanning electron microscope images of the Ti50Al50 undercooled at 345 K. The optical microscope image (Fig. 2a) shows that the microstructure is composed of white matrix containing some needle-like black regions. From this image, it is difficult to draw a conclusion about whether the dark regions belong to the single phase or
Fig. 2. Microstructures of the hypercooled Ti50Al50 alloy (⌬T ⫽ 345 K). (a) Optical image and (b) SEM image.
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Fig. 3. Composition analysis of the hypercooled Ti50Al50 alloy.
not. A more detailed image of the microstructure of the hypercooled alloy was therefore taken on a scanning microscope (see in Fig. 2b). This image clearly shows that there are some dark crossing cracks in the matrix. Those dark cracks are mini-cracks formed from the rapid solidification process of the undercooled melt and belong to the last solidified regions. The local stress centralization caused by the rapid advancing of solid-liquid interface leads to the formation of cracks in the hypercooled samples. The aluminum composition distribution in the hypercooled samples was analyzed and is illustrated in Fig. 3. All sampling positions are marked in Fig. 2b. The mean aluminum content in the core region of matrix was 50.1, which is a little higher than that (49.9) in boundaries close to cracks. Considering the measurement accuracy of the equipment, the composition difference between the firstly solidified region and the cracks must lie in the systematical error range of the equipment. Therefore, the hypercooling state of the alloy provides us an effective way to get the homogenous microstructure. X-ray diffraction technology was used to identify the phase composition. XRD results of the hypercooled sample surface are given in Fig. 4. By comparing these results with standard diffraction parameters of the gamma phase in the Ti-Al alloy system, all peaks could be indexed as a gamma phase; there was, however, a slight shift on the direction of small angles. The decreasing interplane distances caused the diffraction peak angles to shift to the lower value direction. For the hypercooled sample, the primary phase formed from the melt had no time to perform the transformation from disorder to order sequence, which influenced the ordering process, and thus the interdistance among planes decreased. C.D. Anderson et al. [12] found that the growth kinetics of gamma phase increased discontinuously. This was associated with a change in the primary solidification phase from ordered gamma (Vmax ⫽ 0.5 ms⫺1) to disordered gamma (Vmax ⫽ 10 ms⫺1) at a undercooling of 150 K. If the hypercooled melt solidifies in a disorder pattern, then super-crystalline planes of intermetallic gamma phase such as (003) and (110) would disappear and the interdistance of the crystalline planes would increase. Although no super-crystalline planes
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Fig. 4. X-ray diffraction patterns of the hypercooled Ti50Al50 alloy (⌬T ⫽ 345K).
were detected in the XRD spectrum of the hypercooled sample surface, the disorder growth pattern of the hypercooled gamma phase will be explored further. A typical transmission electron microscopy of the hypercooled sample is presented in Fig. 5. The sample is composed of equaxied gamma phase grains. From the selected area electron diffraction (SAD) patterns (see Fig. 5b), rings for planes (111), (200), (220), and (311) were determined, while no sign of the superlattice rings for planes (001) and (110) was found. The absence of SAD rings indexed as superlattice rings is convincing evidence of the disorder growth mode of the gamma phase in the hypercooled melt. An edge-on orientation method was adopted to determine grain size in the hypercooled samples, from transmission electron microscope observations. The average diameter value of grains in the hypercooled sample was found to be 17 nm. Vickers hardness of the hypercooled samples was measured. For comparison, Fig. 6 gives the hardness values of the arc-melted, hypercooled, and heat-treated alloys. The Vickers hardness value for the arc-melted alloy is 315, while that for the hypercooled sample with a nanometer-
Fig. 5. TEM analysis of the hypercooled Ti50Al50 alloy (⌬T ⫽ 345 K). (a) Bright field image and (b) diffraction rings of dark field image.
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Fig. 6. Vickers hardness values of the arc-melted Ti50Al50, hypercooled Ti50Al50, and fully lamellae Ti-45Al2Cr-2Nb alloy.
sized, gamma phase structure is 561. The latter value is a little larger than that of the nanometersized, fully lamella structure of the heat-treatment Ti-45Al-2Cr-2Nb alloy (545) [17].
Acknowledgments The authors would like to gratitude the National Natural Science Foundation of China granted No. 59671045 and the Open Foundation of State Key Lab of Laser Technology for financial support and express their gratitude to Ms X.R. Liu and Ms S.L. Chen for their help in the manuscript preparation.
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17]
F.H. Froes, C. Suryanarayana, D. Eliezer, J. Mater. Sci. 27 (1992) 5113. Y.-W. Kim, J.O.M. (1994) 30. J. Magnan, G.C. Weatherly, M.-C. Cheynet, Metall. Trans. A 30 (1999) 19. J.C. Williams, Mater. Sci. Eng. A 263 (1999) 107. I. Weiss, S.L. Semiatin, Mater. Sci. Eng. A 263 (1999) 243. H. Chang, J. Hofler, C. Altstetter, R. Averback, Mater. Sci. Eng. A 153 (1992) 676. D.N. Senkov, N. Srisukhumbowornchai, M.L. Ovecoglu, F.H. Froes, Scripta Materialia 39 (1998) 691. D.M. Herlach, Mater. Sci. Eng. A 226 –228 (1997) 348. G. Wilde, G.P. Gorler, R. Willnecker, Appl. Phys. Lett. 69 (1996) 2995. R. Willnecker, G.P. Gorler, G. Wilde, Mater. Sci. Eng. A 226 –228 (1997) 439. G. Wilde, G.P. Gorler, R. Willnecker, J. Non-Crystalline Solids 205–207 (1996) 317. C.D. Anderson, W.H. Hofmeister, R.J. Bayuzick, Metall. Trans. A 23 (1992) 2699. J.J. Valencia, C. McCullough, C.G. Levi, R. Mehrabian, Acta Metall. 37 (1989) 2517. J.F. Li, Y.C. Liu, Y.L. Lu, G.C. Yang, Y.H. Zhou, J. Crystal Growth 192 (1998) 462. H. Okamato, J. Phase Equilibria 14 (1993) 120. U.R. Kattner, J.-C. Lin, Y.A. Chang, Metall. Trans. A 23 (1992) 2081. Y.Q. Sun, Metall. Trans. A 29 (1998) 2679.
Materials Research Bulletin 36 (2001) 963–969
Microstructure and mechanical properties of the hypercooled Ti50Al50 alloy Y.C. Liua,*, G.C. Yanga, X.F. Guoa, J.H. Yanga, D.S. Xub, Y.H. Zhoua a
State Key Lab of Solidification Processing, Northwestern Polytechnical University, Xi’an 710072, P. R. China b State Key Lab of Laser Technology, Central University of Sci. & Tech., Wuhan 430074, P. R. China (Refereed) Received 24 July 2000; accepted 14 January 2001
Abstract Hypercooling region of the gamma titanium aluminide alloy Ti50Al50 was achieved by cyclically superheating the alloy melt in a containerless electromagnetic levitation apparatus. The maximum undercooling of the alloy melt amounted to 345 K. XRD, TEM, SEM and optical microscopy techniques were adopted to investigate the microstructure and identify the phase composition. The microstructure of the hypercooled alloy was composed of equaxied gamma phase grains with a mean diameter of 17 nm. The fine microstructure of the gamma phases produced a large strengthening effect, raising the hardness to 545 on the Vickers scale. The aluminum composition distribution in the hypercooled sample was homogenous except for a small difference between first solidified regions and boundaries of cracks. X-ray diffraction results showed that all peaks of the gamma phase shifted slightly in the direction of small angles. This can be explained by the disordering growth pattern caused by the rapid solidification process in the hypercooled melt. © 2001 Elsevier Science Ltd. All rights reserved. Keywords: A. Intermetallic compounds; C. X-ray diffraction; D. Microstructure; D. Mechanical properties
1. Introduction Lightweight gamma titanium aluminide alloys are potential candidate materials in automotive and aerospace industries for high temperature structural applications [1]. However, * Corresponding author. Present address: Max-Planck-Institute for Metals Research, Seestr. 92, D-70174 Stuttgart, Germany. Tel.: ⫹49-711-2095427; fax: ⫹49-711-2095418. E-mail address: [email protected] (Y.C. Liu). 0025-5408/01/$ – see front matter © 2001 Elsevier Science Ltd. All rights reserved. PII: S 0 0 2 5 - 5 4 0 8 ( 0 1 ) 0 0 5 7 7 - 3
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their poor ductility and low fracture toughness at ambient temperature are the major limitations for practical applications [2]. Many efforts to introduce grain boundary deformation mode and suppress crack propagation have been made to improve their ductility and fracture toughness by significantly decreasing the grain size of the material [3–5]. For example, magnetron sputtering and condensation [6] and also mechanical alloying [7] have been used successfully to produce the nanophase of gamma titanium aluminide alloy. Other methods commonly used to refine the grains in the development of high performance materials are the rapid quenching process and thermodynamic undercooling. Due to volume limitation, the rapid quenching process tends to be used only on the laboratory scale. Liquid undercooling is known to significantly affect the solidification process, resulting in changes in the microstructure development and bulky metastable phase formation [8]. Several secondary processes, such as dendritic remelting, structural transformation, and segregation, generally affect solidification at low undercoolings, resulting in a variety of microstructural patterns. In this regard, attainment of hypercooling is of special interest. In recent years, investigations on high melting bulk (Co, Ni, Fe)-Pd and Au-(Cu, Ni) alloys had demonstrated experimentally the achievement of the hypercooling region [9 –11]. Undercooling experiments had been performed on alloy Ti50Al50 involving intermetallic compounds, and a maximum undercooling of 300 K was achieved [12]. Different, even conflicting results about the formation of primary phases in the melt at the lower undercooling have been reported. Similar results had been reported for the highly undercooled region of the melt [12,13].
2. Experimental 2.1. Ingot preparation High-purity (⬎ 99.99%) aluminum and titanium were used in the appropriate amounts to form an alloy with a stoichiometry of Ti50Al50. The process of melting was carried out in a vacuum arc furnace under a high-purity argon atmosphere, to produce a button, approximately 4 cm in diameter, which was remelted four times. The arc-melted button was homogenized in a vacuum heat-treatment furnace at 1273 K for 80 h before undercooling experiments were carried out, in order to achieve complete composition homogenization. 2.2. Undercooling experiments The undercooling experiments were carried out in an electromagnetic levitation melting apparatus manufactured by Edmuld Buhler Co. The device possesses a maximum levitation capacity of 200 mg and a HF-3 high-frequency generator is used to heat the alloy. In the undercooling process, the working chamber was initially evacuated to about 10⫺9 mbar, and then back-filled with high-purity argon gas (⬎ 99.9999%). For the purpose of deactivating heterogeneous nucleation sites, each sample, approximately 50 mg, was cyclically superheated at 300 K for 5 min, then the power was turned off, the superheated sample
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was allowed to cool on an inert zirconia sample holder, which was stretched into the coil in advance. When a desired undercooling of the alloy melt was reached, the sample was cooled naturally to the ambient temperature on the holder. The cooling rate of the process was approximately in the magnitude of 102⬃103 K/s. An infrared pyrometer with an absolute accuracy, relative accuracy, and response time of less than 10 K, 3 K and 5 ms, respectively, was used to record the thermal history of samples [14]. The cooling curve was calibrated with a standard PtRh30-PtRh6 thermal couple, which was encapsulated in a silica tube and then immersed into the melt. The melting temperature and undercooling degree in the cooling curve could be obtained by comparing the absolute temperature recorded by the standard thermal couple. 2.3. Characterization Crystallographic features were mainly examined using Rigaku X-ray power diffractometer with a Cu K␣ source. The slice for transmission microscopy analysis was electrolytically thinned using an electrolyte of 5% HClO4 in ethanol at 243 K and investigated using a JEOL JEM-200cx electron microscope. Microstructure observation was carried out using a JEOL JXA-840 scanning electron microscope. Kroll’s etching regent and a Nikon Epiphot optical microscopy were used for observing the microstructure. Hardness measurements of the polished specimens were carried out on a Vickers hardness indentor. The indentation load and time were 50 g and 10 s, respectively.
3. Experimental results and discussion According to the binary phase diagram of Ti-Al [15], the composition of alloy Al50Ti50 falls in the L3␣ equilibrium solidification region and the equilibrium microstructure of it is in the two-phase (␣2⫹␥) region. Fig. 1a shows the optical microscope image of the arc-melted alloy homogenized at 1273 K for 80 h. The structure is composed of equaxied grains and the lamellar structure could be made out in grains. Fig. 1b shows a more detailed image of the iterative white and black lamella (␣2⫹␥) phases, taken with the transmission electron microscope. The hypercooling regime is achieved when the enthalpy of a liquid is reduced with decreasing temperature to a level equal to that of the solid phase at the melting temperature, Tm. The isenthalpic temperature, Ti, which represents the upper limit of the hypercooling regime, is defined by ⌬H f ⫽
冕
Tm
C pl共T兲dT
Ti
with ⌬Hf the enthalpy of fusion and cpl(T) the specific heat of the undercooled liquid. The enthalpy of fusion of Ti50Al50 can be derived from the calculation based on the sublattice
966
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Fig. 1. Microstructures of the arc-melted Ti50Al50 alloy homogenized at 1273 K for 80 hours: Optical image and bright field image of lamellar structure.
model [16]. Then, the minimum of undercooling to achieve the hypercooled state of Ti50Al50 can be calculated from ⌬T k ⫽
⌬H m ⫹ 共T l ⫺ T s兲 Cp
where, Tl is the liquid temperature, and Ts, the solid temperature. The minimum calculated undercooling for achieving the hypercooling state is 340 K. Therefore, the sample of Ti50Al50 alloy melt with an undercooling above 340 K should solidify as a hypercooled mode. Fig. 2 gives the optical and scanning electron microscope images of the Ti50Al50 undercooled at 345 K. The optical microscope image (Fig. 2a) shows that the microstructure is composed of white matrix containing some needle-like black regions. From this image, it is difficult to draw a conclusion about whether the dark regions belong to the single phase or
Fig. 2. Microstructures of the hypercooled Ti50Al50 alloy (⌬T ⫽ 345 K). (a) Optical image and (b) SEM image.
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Fig. 3. Composition analysis of the hypercooled Ti50Al50 alloy.
not. A more detailed image of the microstructure of the hypercooled alloy was therefore taken on a scanning microscope (see in Fig. 2b). This image clearly shows that there are some dark crossing cracks in the matrix. Those dark cracks are mini-cracks formed from the rapid solidification process of the undercooled melt and belong to the last solidified regions. The local stress centralization caused by the rapid advancing of solid-liquid interface leads to the formation of cracks in the hypercooled samples. The aluminum composition distribution in the hypercooled samples was analyzed and is illustrated in Fig. 3. All sampling positions are marked in Fig. 2b. The mean aluminum content in the core region of matrix was 50.1, which is a little higher than that (49.9) in boundaries close to cracks. Considering the measurement accuracy of the equipment, the composition difference between the firstly solidified region and the cracks must lie in the systematical error range of the equipment. Therefore, the hypercooling state of the alloy provides us an effective way to get the homogenous microstructure. X-ray diffraction technology was used to identify the phase composition. XRD results of the hypercooled sample surface are given in Fig. 4. By comparing these results with standard diffraction parameters of the gamma phase in the Ti-Al alloy system, all peaks could be indexed as a gamma phase; there was, however, a slight shift on the direction of small angles. The decreasing interplane distances caused the diffraction peak angles to shift to the lower value direction. For the hypercooled sample, the primary phase formed from the melt had no time to perform the transformation from disorder to order sequence, which influenced the ordering process, and thus the interdistance among planes decreased. C.D. Anderson et al. [12] found that the growth kinetics of gamma phase increased discontinuously. This was associated with a change in the primary solidification phase from ordered gamma (Vmax ⫽ 0.5 ms⫺1) to disordered gamma (Vmax ⫽ 10 ms⫺1) at a undercooling of 150 K. If the hypercooled melt solidifies in a disorder pattern, then super-crystalline planes of intermetallic gamma phase such as (003) and (110) would disappear and the interdistance of the crystalline planes would increase. Although no super-crystalline planes
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Fig. 4. X-ray diffraction patterns of the hypercooled Ti50Al50 alloy (⌬T ⫽ 345K).
were detected in the XRD spectrum of the hypercooled sample surface, the disorder growth pattern of the hypercooled gamma phase will be explored further. A typical transmission electron microscopy of the hypercooled sample is presented in Fig. 5. The sample is composed of equaxied gamma phase grains. From the selected area electron diffraction (SAD) patterns (see Fig. 5b), rings for planes (111), (200), (220), and (311) were determined, while no sign of the superlattice rings for planes (001) and (110) was found. The absence of SAD rings indexed as superlattice rings is convincing evidence of the disorder growth mode of the gamma phase in the hypercooled melt. An edge-on orientation method was adopted to determine grain size in the hypercooled samples, from transmission electron microscope observations. The average diameter value of grains in the hypercooled sample was found to be 17 nm. Vickers hardness of the hypercooled samples was measured. For comparison, Fig. 6 gives the hardness values of the arc-melted, hypercooled, and heat-treated alloys. The Vickers hardness value for the arc-melted alloy is 315, while that for the hypercooled sample with a nanometer-
Fig. 5. TEM analysis of the hypercooled Ti50Al50 alloy (⌬T ⫽ 345 K). (a) Bright field image and (b) diffraction rings of dark field image.
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Fig. 6. Vickers hardness values of the arc-melted Ti50Al50, hypercooled Ti50Al50, and fully lamellae Ti-45Al2Cr-2Nb alloy.
sized, gamma phase structure is 561. The latter value is a little larger than that of the nanometersized, fully lamella structure of the heat-treatment Ti-45Al-2Cr-2Nb alloy (545) [17].
Acknowledgments The authors would like to gratitude the National Natural Science Foundation of China granted No. 59671045 and the Open Foundation of State Key Lab of Laser Technology for financial support and express their gratitude to Ms X.R. Liu and Ms S.L. Chen for their help in the manuscript preparation.
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17]
F.H. Froes, C. Suryanarayana, D. Eliezer, J. Mater. Sci. 27 (1992) 5113. Y.-W. Kim, J.O.M. (1994) 30. J. Magnan, G.C. Weatherly, M.-C. Cheynet, Metall. Trans. A 30 (1999) 19. J.C. Williams, Mater. Sci. Eng. A 263 (1999) 107. I. Weiss, S.L. Semiatin, Mater. Sci. Eng. A 263 (1999) 243. H. Chang, J. Hofler, C. Altstetter, R. Averback, Mater. Sci. Eng. A 153 (1992) 676. D.N. Senkov, N. Srisukhumbowornchai, M.L. Ovecoglu, F.H. Froes, Scripta Materialia 39 (1998) 691. D.M. Herlach, Mater. Sci. Eng. A 226 –228 (1997) 348. G. Wilde, G.P. Gorler, R. Willnecker, Appl. Phys. Lett. 69 (1996) 2995. R. Willnecker, G.P. Gorler, G. Wilde, Mater. Sci. Eng. A 226 –228 (1997) 439. G. Wilde, G.P. Gorler, R. Willnecker, J. Non-Crystalline Solids 205–207 (1996) 317. C.D. Anderson, W.H. Hofmeister, R.J. Bayuzick, Metall. Trans. A 23 (1992) 2699. J.J. Valencia, C. McCullough, C.G. Levi, R. Mehrabian, Acta Metall. 37 (1989) 2517. J.F. Li, Y.C. Liu, Y.L. Lu, G.C. Yang, Y.H. Zhou, J. Crystal Growth 192 (1998) 462. H. Okamato, J. Phase Equilibria 14 (1993) 120. U.R. Kattner, J.-C. Lin, Y.A. Chang, Metall. Trans. A 23 (1992) 2081. Y.Q. Sun, Metall. Trans. A 29 (1998) 2679.