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Effect of Mn on formation of lamellar structure in a γ-TiAl alloy
Materials Letters 58 (2004) 3756 – 3760 www.elsevier.com/locate/matlet
Effect of Mn on formation of lamellar structure in a g-TiAl alloy Yonggang Jina, J.N. Wangb,*, Yong Wangb, Jie Yangb a
Shanghai Qiang Bao Metallurgical Materials Ltd. Co., Shanghai 200901, People’s Republic of China School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200030, People’s Republic of China
b
Received 10 March 2004; received in revised form 27 July 2004; accepted 5 August 2004 Available online 27 August 2004
Abstract A fine cast microstructure can be developed in two TiAl alloys with or without Mn. It can be optimized to be a fine-grained fully lamellar structure (FFLS) by heat treatment. But the presence of Mn was found to affect the development of the lamellar structure. Under furnace cooling, a FFLS can be developed in the alloy without Mn, but a duplex structure in the alloy with Mn. A FFLS can be developed in the alloy with Mn only if the cooling rate is increased. This effect may be a result from the promoted formation of cellular g grains in the presence of Mn. D 2004 Elsevier B.V. All rights reserved. Keywords: Titanium aluminides; Casting; Microstructure; Fine-grained fully lamellar structure; Mn alloying
1. Introduction TiAl-based alloys are of growing interest to the aerospace, automotive and power generation industries since they offer a remarkable combination of low density, high specific modulus and strength, good creep and oxidation resistance [1–3]. However, the practical usage of these alloys still needs the improvement in properties, such as room temperature ductility, fracture toughness and high temperature creep resistance. It is well known that mechanical properties of this class of alloys are strongly microstructure-dependent [4]. Numerous studies have shown that those alloys with a fine-grained fully lamellar structure (FFLS) possess good balanced mechanical properties [5]. So far, much work has been devoted to studying the effects of different alloying elements in g-TiAl. Some of the alloying elements are known to be beta stabilisers such as W, Nb, Mo. Such elements can change the solidification path of g-TiAl to realize cast microstructure control [6]. But other elements are considered to optimise some in-service properties, such as Cr, Mn for improving ductility and Si, C
* Corresponding author. Tel.: +86 2156107316; fax: +86 2156801836. E-mail address: [email protected] (J.N. Wang). 0167-577X/$ - see front matter D 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.matlet.2004.08.009
for raising high-temperature creep resistance [7–10]. When TiAl alloys contain such alloying elements, phase transformation during heat treatment may be significantly changed, and thus the lamellar structure formation is affected. In the present paper, two cast Ti–44Al–7Nb (at.%) alloys with and without Mn were studied. They are shown to have fine cast microstructures that can be optimized to be finegrained fully lamellar structures (FFLS) by heat treatment. But the presence of Mn was found to affect the development of such microstructure.
2. Experimental TiAl based alloys with nominal compositions (at.%) of Ti–44Al–7Nb (alloy I) and Ti–44Al–7Nb–2Mn (alloy II) were used in this study. The alloys were prepared as 80g buttons by arc melting in argon atmosphere in a watercooled copper hearth. The buttons were remelted 4–5 times to ensure chemical homogeneity. Specimens of a size 10108 mm3 were cut by electric discharge machine from the alloy buttons for heat treatment. The specimens were held at 1350 8C for 1 h. They were then cooled at various cooling rates, specifically, furnace cooling (FC), air
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cooling (AC), controlled cooling (CC) and water quenching (WQ). The cooling rate for CC was between those for FC and AC. The phases presented in the alloys were studied by using optical microscopy (OM), scanning electron microscopy (SEM) in the back scattered electron (BSE) mode, transmission electron microscopy (TEM) and X-ray diffraction (XRD). The samples for OM were etched in a solution of modified Kroll’s reagent (5% HF, 15% HNO3, 80% H2O). Those for TEM were prepared by twin-jet polishing in a solution of 8% perchloric acid and 92% ethanol.
3. Results The as-cast microstructures in the two alloys are generally similar, consisting primarily of a lamellar structure with a colony size of ~50 Am. Dendrites were observed having a typical fourfold symmetry feature (Fig. 1a). BSE– SEM (Fig. 1b) examination revealed the presence of untransformed phases both within the lamellar grains and at their boundaries. Detailed TEM study showed that they were g and h phases (Fig. 2). The XRD analysis also revealed the presence of three phases g (L10), a2 (DO19) and B2 (Fig. 3). For the development of a FFLS, the dendrites and the interlaced g and h phases within the lamellar grains and the granular ones at their boundaries should be eliminated by heat treatment. In order to obtain a fully lamellar structure, the alloys with a single a phase field are subject to heat treatment in
Fig. 1. The as-cast microstructure showing the fourfold symmetry of dendrites (alloy II) (a) and BSE–SEM image (b) showing the remaining phases within lamellar grains and their boundaries (alloy I) (bright-beta, dark-gamma, grey-(a+g) lamellar).
Fig. 2. TEM image of g and h phases contained at lamellar grain boundaries (a) and their selected area diffraction patterns (SADPs) (b, c).
this field followed by an appropriate cooling. Those alloys without a single a phase field are instead heat-treated in the a+h regime. During cooling, the a phase transforms to a lamellar structure, whereas h transforms to a (which can subsequently transform to a2+g lamellae) or g. In the present work, the samples were held at temperatures ranging from 1200 to 1400 8C for an hour. This was followed by WQ to investigate the phase constitution of the two alloys at high temperatures. Experimental observation showed that there was no single a phase field but an a+h phase field at high temperature. To obtain a FFLS, both alloys were subject to heat treatment at 1350 8C for 1 h followed by either FC or AC to obtain a FFLS. Fig. 4a shows such a microstructure for alloy I for the FC case. It was made up mainly of a2+g lamellar grains with a size of ~60 Am. Only a small amount of retained h (B2) phase was observed at two-grain boundaries and triple junctions. However, the microstructure developed in alloy II for the FC case was a duplex (DP) one, which consisted mainly of fine equiaxed g grains and lamellar colonies (Fig. 4b). A FFLS was obtained only when the cooling rate was increased to that for CC (R150 8C/min) (Fig. 4c). To understand the different behaviors of the two alloys during FC, some samples were furnace cooled from 1350 8C to various temperatures followed by WQ. Fig. 5a, b, c shows the microstructures in alloy II furnace cooled to and then water quenched from 1250, 1230 and 1220 8C, respectively. As shown in Fig. 5a, the sample quenched from 1250 8C can transform to a two-phase microstructure with equiaxed a2 grains and B2 phases presented at boundaries and triple points. When quenched from 1230
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Fig. 3. X-ray diffraction pattern of the as-cast alloy (alloy II).
8C, the sample exhibits a few black grains within grainboundary regions. They might be g phases based on the XDR pattern (Fig. 6). When quenched from 1220 8C, the sample presents more newly grown g grains at grain boundaries (Fig. 5c). Furthermore, a few gamma lathes start to form inside alpha grains, which suggests that the starting temperature of lamellar formation (TLS) for alloy II is about 1220 8C for the FC case. Fig. 5d shows the microstructure obtained from the sample of alloy I quenched from 1200 8C. As can be seen, there exist no gamma grains at grain boundaries before the lamellae for alloy I begin to form at 1200 8C. In an attempt to understand the microstructural changes from a DP structure to a FFLS for alloy II cooled at the cooling rate above 150 8C/min, some experiments were conducted on the samples which were cooled at the cooling
rate of 150 8C/min from 1350 8C to different temperatures followed by WQ. When the sample was quenched from 1180 8C, the obtained structure is similar to that shown in Fig. 5d. It suggests that the microstructure evolution in alloy II during CC is the same as that in alloy I during FC.
4. Discussion It is interesting to note that in contrast to alloy I, alloy II did not develop a FFLS but a DP structure when it was slowly cooled (i.e. FC) from high temperature. The present results indicate that g grains preferentially form within grain-boundary regions prior to the lamellar formation when alloy II is cooled slowly. However, such experimental phenomenon is not found in alloy I. This difference between
Fig. 4. Optical microstructures after holding at 1350 8C for an hour and then FC or AC. (a) Alloy I, FC; (b) alloy II, FC; (c) alloy II, CC.
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Fig. 5. Optical microstructure obtained from the two alloys after furnace cooled from 1350 8C to various temperatures followed by water quenching. (a) Alloy II, 1250 8C; (b) alloy II, 1230 8C; (c) alloy II, 1220 8C; (d) alloy I, 1200 8C.
them affects the microstructural evolution and results in a DP structure in alloy II. As shown in previous studies [11, 12], an important competitive reaction to lamellar formation is the decomposition of a (or a+h) phase resulting in formation of cellular g grains (a(a+h)Ygc) at grain boundaries in the case of slow cooling. Fig. 7 is the CCT schematic diagram showing the lamellar and cellular start curves for the present two alloys and the previous 47-2-2-1 alloy with a composition of Ti–46.5Al–2Nb–2Cr–0.9Mo– 0.2B [12]. The 47-2-2-1 alloy is an example which provides only a narrow cooling-rate window (gap between lamellar and cellular start curves), and therefore the
cellular reaction in this alloy was so prevalent that fully lamellar microstructures could not be formed. For alloy II with Mn, fully lamellar microstructures cannot also be produced when slowly cooled since such structures are affected by rapid cellular growth of gamma from the grainboundary regions. In contrast to the above two alloys, alloy I has such a wide cooling-rate window that the lamellar microstructures can be produced in the case of FC. In the ternary Ti–Al–Mn system, Mn has been reported to preferentially substitute for Al sublattice site in TiAl [13], thereby reducing the equilibrium aluminum concentration of
Fig. 6. X-ray diffraction patterns observed from alloy II in Fig. 5a and b states.
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minimize the cellular start reaction. Whether lamellar or cellular reaction occurs during cooling is strongly dependent on the chemical compositions of alloys and the cooling rate. This is important for developing lamellar structures in engineered components.
5. Conclusions The as-cast microstructure can be optimized to be a finegrained fully lamellar structure for the present two alloys. But microalloying such as Mn addition has a pronounced effect on the formation of such a structure during slow cooling. Without the Mn addition, this structure can be produced at both FC and AC. But with the Mn addition, this structure can only be developed at the cooling rate above 150 8C/min. A wide cooling-rate window is important for the development of a fully lamellar structure, and thus should be considered for alloy design.
Acknowledgements Fig. 7. Schematic continuous-cooling-transformation (CCT) diagram for the present two alloys and the b47-2-2-1Q alloy [12]. The solid and the dashed lines represent the start curves of lamellar formation and cellular reaction, respectively.
This work was supported by the National Natural Science Foundation of China (Project no. 50271039).
References g phase. This shifts g phase field towards the side of deficient aluminum in TiAl pseudo-binary phase diagram, which increases the ratio of g phase in the three phase a+g+h field, and finally favours the gc formation within grain-boundary regions. Based on the above discussion, it can be deduced that the addition of Mn element in alloy II changes the phase equilibria in alloy I, and finally decreases the driving force required for formation of gc microconstituent. Thus, the reaction a (a+h)Ygc is likely to occur in the case of the low undercooling. In contrast to alloy I, TLS for alloy II is increased by 20 to 1220 8C for the FC case. It can be explained using the hypothesis that the undercooling required for lamellar formation is also reduced due to the effect of Mn element on the TiAl phase equilibria. When the cooling rate is increased to a critical rate (i.e. 150 8C/min), gc transformation is restrained in alloy II. During the consequent cooling, a FFLS could be obtained in the same way with alloy I by FC. As suggested by CCT curves in Fig. 7, the controlled cooling-rate above 150 8C/ min traverses the lamellar start temperature to avoid or
[1] Y.W. Kim, D.M. Dimiduk, JOM 43 (1991) 40. [2] Y.W. Kim, JOM 46 (1994) 30. [3] D.M. Dimiduk, in: Y.W. Kim, R. Wagner, M. Yamaguchi (Eds.), Gamma Titanium Aluminides, TMS, 1995, p. 3. [4] F. Appel, R. Wagner, Mater. Sci. Eng., R 22 (1988) 187. [5] Y.W. Kim, D.M. Dimiduk, in: M.V. Nathal, R. Darolia, C.T. Liu, P.L. Martin, D.B. Miracle, R. Wagner, M. Yamaguchi (Eds.), Structural Intermetallics, TMS, 1997, p. 531. [6] S. Naka, M. Thomas, C. Sanchez, T. Khan, in: M.V. Nathal, R. Darolia, C.T. Liu, P.L. Martin, D.B. Miracle, R. Wagner, M. Yamaguchi (Eds.), Structural Intermetallics, TMS, 1997, p. 313. [7] T. Kawabata, T. Tamura, O. Izumi, Metall. Trans. 24A (1993) 141. [8] S.C. Huang, E.L. Hall, Acta Metall. Mater. 39 (1991) 1053. [9] T. Noda, M. Okable, S. Isobe, M. Sayashi, Mater. Sci. Eng., A 192/ 193 (1995) 774. [10] W.H. Tian, M. Nemoto, in: Y.W. Kim, R. Wagner, M. Yamaguchi (Eds.), Gamma Titanium Aluminides, TMS, 1995, p. 689. [11] D.M. Dimiduk, V.K. Vasudevan, in: Y.W. Kim, D.M. Dimiduk, M.H. Loretto (Eds.), Gamma Titanium Aluminides, TMS, 1999, p. 239. [12] D.M. Dimiduk, P.L. Martin, Y.W. Kim, Mater. Sci. Eng., A 243 (1998) 66. [13] Y.L. Hao, D.S. Xu, Y.Y. Cui, R. Yang, D. Li, Acta Mater. 47 (1999) 1129.
Effect of Mn on formation of lamellar structure in a g-TiAl alloy Yonggang Jina, J.N. Wangb,*, Yong Wangb, Jie Yangb a
Shanghai Qiang Bao Metallurgical Materials Ltd. Co., Shanghai 200901, People’s Republic of China School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200030, People’s Republic of China
b
Received 10 March 2004; received in revised form 27 July 2004; accepted 5 August 2004 Available online 27 August 2004
Abstract A fine cast microstructure can be developed in two TiAl alloys with or without Mn. It can be optimized to be a fine-grained fully lamellar structure (FFLS) by heat treatment. But the presence of Mn was found to affect the development of the lamellar structure. Under furnace cooling, a FFLS can be developed in the alloy without Mn, but a duplex structure in the alloy with Mn. A FFLS can be developed in the alloy with Mn only if the cooling rate is increased. This effect may be a result from the promoted formation of cellular g grains in the presence of Mn. D 2004 Elsevier B.V. All rights reserved. Keywords: Titanium aluminides; Casting; Microstructure; Fine-grained fully lamellar structure; Mn alloying
1. Introduction TiAl-based alloys are of growing interest to the aerospace, automotive and power generation industries since they offer a remarkable combination of low density, high specific modulus and strength, good creep and oxidation resistance [1–3]. However, the practical usage of these alloys still needs the improvement in properties, such as room temperature ductility, fracture toughness and high temperature creep resistance. It is well known that mechanical properties of this class of alloys are strongly microstructure-dependent [4]. Numerous studies have shown that those alloys with a fine-grained fully lamellar structure (FFLS) possess good balanced mechanical properties [5]. So far, much work has been devoted to studying the effects of different alloying elements in g-TiAl. Some of the alloying elements are known to be beta stabilisers such as W, Nb, Mo. Such elements can change the solidification path of g-TiAl to realize cast microstructure control [6]. But other elements are considered to optimise some in-service properties, such as Cr, Mn for improving ductility and Si, C
* Corresponding author. Tel.: +86 2156107316; fax: +86 2156801836. E-mail address: [email protected] (J.N. Wang). 0167-577X/$ - see front matter D 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.matlet.2004.08.009
for raising high-temperature creep resistance [7–10]. When TiAl alloys contain such alloying elements, phase transformation during heat treatment may be significantly changed, and thus the lamellar structure formation is affected. In the present paper, two cast Ti–44Al–7Nb (at.%) alloys with and without Mn were studied. They are shown to have fine cast microstructures that can be optimized to be finegrained fully lamellar structures (FFLS) by heat treatment. But the presence of Mn was found to affect the development of such microstructure.
2. Experimental TiAl based alloys with nominal compositions (at.%) of Ti–44Al–7Nb (alloy I) and Ti–44Al–7Nb–2Mn (alloy II) were used in this study. The alloys were prepared as 80g buttons by arc melting in argon atmosphere in a watercooled copper hearth. The buttons were remelted 4–5 times to ensure chemical homogeneity. Specimens of a size 10108 mm3 were cut by electric discharge machine from the alloy buttons for heat treatment. The specimens were held at 1350 8C for 1 h. They were then cooled at various cooling rates, specifically, furnace cooling (FC), air
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cooling (AC), controlled cooling (CC) and water quenching (WQ). The cooling rate for CC was between those for FC and AC. The phases presented in the alloys were studied by using optical microscopy (OM), scanning electron microscopy (SEM) in the back scattered electron (BSE) mode, transmission electron microscopy (TEM) and X-ray diffraction (XRD). The samples for OM were etched in a solution of modified Kroll’s reagent (5% HF, 15% HNO3, 80% H2O). Those for TEM were prepared by twin-jet polishing in a solution of 8% perchloric acid and 92% ethanol.
3. Results The as-cast microstructures in the two alloys are generally similar, consisting primarily of a lamellar structure with a colony size of ~50 Am. Dendrites were observed having a typical fourfold symmetry feature (Fig. 1a). BSE– SEM (Fig. 1b) examination revealed the presence of untransformed phases both within the lamellar grains and at their boundaries. Detailed TEM study showed that they were g and h phases (Fig. 2). The XRD analysis also revealed the presence of three phases g (L10), a2 (DO19) and B2 (Fig. 3). For the development of a FFLS, the dendrites and the interlaced g and h phases within the lamellar grains and the granular ones at their boundaries should be eliminated by heat treatment. In order to obtain a fully lamellar structure, the alloys with a single a phase field are subject to heat treatment in
Fig. 1. The as-cast microstructure showing the fourfold symmetry of dendrites (alloy II) (a) and BSE–SEM image (b) showing the remaining phases within lamellar grains and their boundaries (alloy I) (bright-beta, dark-gamma, grey-(a+g) lamellar).
Fig. 2. TEM image of g and h phases contained at lamellar grain boundaries (a) and their selected area diffraction patterns (SADPs) (b, c).
this field followed by an appropriate cooling. Those alloys without a single a phase field are instead heat-treated in the a+h regime. During cooling, the a phase transforms to a lamellar structure, whereas h transforms to a (which can subsequently transform to a2+g lamellae) or g. In the present work, the samples were held at temperatures ranging from 1200 to 1400 8C for an hour. This was followed by WQ to investigate the phase constitution of the two alloys at high temperatures. Experimental observation showed that there was no single a phase field but an a+h phase field at high temperature. To obtain a FFLS, both alloys were subject to heat treatment at 1350 8C for 1 h followed by either FC or AC to obtain a FFLS. Fig. 4a shows such a microstructure for alloy I for the FC case. It was made up mainly of a2+g lamellar grains with a size of ~60 Am. Only a small amount of retained h (B2) phase was observed at two-grain boundaries and triple junctions. However, the microstructure developed in alloy II for the FC case was a duplex (DP) one, which consisted mainly of fine equiaxed g grains and lamellar colonies (Fig. 4b). A FFLS was obtained only when the cooling rate was increased to that for CC (R150 8C/min) (Fig. 4c). To understand the different behaviors of the two alloys during FC, some samples were furnace cooled from 1350 8C to various temperatures followed by WQ. Fig. 5a, b, c shows the microstructures in alloy II furnace cooled to and then water quenched from 1250, 1230 and 1220 8C, respectively. As shown in Fig. 5a, the sample quenched from 1250 8C can transform to a two-phase microstructure with equiaxed a2 grains and B2 phases presented at boundaries and triple points. When quenched from 1230
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Fig. 3. X-ray diffraction pattern of the as-cast alloy (alloy II).
8C, the sample exhibits a few black grains within grainboundary regions. They might be g phases based on the XDR pattern (Fig. 6). When quenched from 1220 8C, the sample presents more newly grown g grains at grain boundaries (Fig. 5c). Furthermore, a few gamma lathes start to form inside alpha grains, which suggests that the starting temperature of lamellar formation (TLS) for alloy II is about 1220 8C for the FC case. Fig. 5d shows the microstructure obtained from the sample of alloy I quenched from 1200 8C. As can be seen, there exist no gamma grains at grain boundaries before the lamellae for alloy I begin to form at 1200 8C. In an attempt to understand the microstructural changes from a DP structure to a FFLS for alloy II cooled at the cooling rate above 150 8C/min, some experiments were conducted on the samples which were cooled at the cooling
rate of 150 8C/min from 1350 8C to different temperatures followed by WQ. When the sample was quenched from 1180 8C, the obtained structure is similar to that shown in Fig. 5d. It suggests that the microstructure evolution in alloy II during CC is the same as that in alloy I during FC.
4. Discussion It is interesting to note that in contrast to alloy I, alloy II did not develop a FFLS but a DP structure when it was slowly cooled (i.e. FC) from high temperature. The present results indicate that g grains preferentially form within grain-boundary regions prior to the lamellar formation when alloy II is cooled slowly. However, such experimental phenomenon is not found in alloy I. This difference between
Fig. 4. Optical microstructures after holding at 1350 8C for an hour and then FC or AC. (a) Alloy I, FC; (b) alloy II, FC; (c) alloy II, CC.
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Fig. 5. Optical microstructure obtained from the two alloys after furnace cooled from 1350 8C to various temperatures followed by water quenching. (a) Alloy II, 1250 8C; (b) alloy II, 1230 8C; (c) alloy II, 1220 8C; (d) alloy I, 1200 8C.
them affects the microstructural evolution and results in a DP structure in alloy II. As shown in previous studies [11, 12], an important competitive reaction to lamellar formation is the decomposition of a (or a+h) phase resulting in formation of cellular g grains (a(a+h)Ygc) at grain boundaries in the case of slow cooling. Fig. 7 is the CCT schematic diagram showing the lamellar and cellular start curves for the present two alloys and the previous 47-2-2-1 alloy with a composition of Ti–46.5Al–2Nb–2Cr–0.9Mo– 0.2B [12]. The 47-2-2-1 alloy is an example which provides only a narrow cooling-rate window (gap between lamellar and cellular start curves), and therefore the
cellular reaction in this alloy was so prevalent that fully lamellar microstructures could not be formed. For alloy II with Mn, fully lamellar microstructures cannot also be produced when slowly cooled since such structures are affected by rapid cellular growth of gamma from the grainboundary regions. In contrast to the above two alloys, alloy I has such a wide cooling-rate window that the lamellar microstructures can be produced in the case of FC. In the ternary Ti–Al–Mn system, Mn has been reported to preferentially substitute for Al sublattice site in TiAl [13], thereby reducing the equilibrium aluminum concentration of
Fig. 6. X-ray diffraction patterns observed from alloy II in Fig. 5a and b states.
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minimize the cellular start reaction. Whether lamellar or cellular reaction occurs during cooling is strongly dependent on the chemical compositions of alloys and the cooling rate. This is important for developing lamellar structures in engineered components.
5. Conclusions The as-cast microstructure can be optimized to be a finegrained fully lamellar structure for the present two alloys. But microalloying such as Mn addition has a pronounced effect on the formation of such a structure during slow cooling. Without the Mn addition, this structure can be produced at both FC and AC. But with the Mn addition, this structure can only be developed at the cooling rate above 150 8C/min. A wide cooling-rate window is important for the development of a fully lamellar structure, and thus should be considered for alloy design.
Acknowledgements Fig. 7. Schematic continuous-cooling-transformation (CCT) diagram for the present two alloys and the b47-2-2-1Q alloy [12]. The solid and the dashed lines represent the start curves of lamellar formation and cellular reaction, respectively.
This work was supported by the National Natural Science Foundation of China (Project no. 50271039).
References g phase. This shifts g phase field towards the side of deficient aluminum in TiAl pseudo-binary phase diagram, which increases the ratio of g phase in the three phase a+g+h field, and finally favours the gc formation within grain-boundary regions. Based on the above discussion, it can be deduced that the addition of Mn element in alloy II changes the phase equilibria in alloy I, and finally decreases the driving force required for formation of gc microconstituent. Thus, the reaction a (a+h)Ygc is likely to occur in the case of the low undercooling. In contrast to alloy I, TLS for alloy II is increased by 20 to 1220 8C for the FC case. It can be explained using the hypothesis that the undercooling required for lamellar formation is also reduced due to the effect of Mn element on the TiAl phase equilibria. When the cooling rate is increased to a critical rate (i.e. 150 8C/min), gc transformation is restrained in alloy II. During the consequent cooling, a FFLS could be obtained in the same way with alloy I by FC. As suggested by CCT curves in Fig. 7, the controlled cooling-rate above 150 8C/ min traverses the lamellar start temperature to avoid or
[1] Y.W. Kim, D.M. Dimiduk, JOM 43 (1991) 40. [2] Y.W. Kim, JOM 46 (1994) 30. [3] D.M. Dimiduk, in: Y.W. Kim, R. Wagner, M. Yamaguchi (Eds.), Gamma Titanium Aluminides, TMS, 1995, p. 3. [4] F. Appel, R. Wagner, Mater. Sci. Eng., R 22 (1988) 187. [5] Y.W. Kim, D.M. Dimiduk, in: M.V. Nathal, R. Darolia, C.T. Liu, P.L. Martin, D.B. Miracle, R. Wagner, M. Yamaguchi (Eds.), Structural Intermetallics, TMS, 1997, p. 531. [6] S. Naka, M. Thomas, C. Sanchez, T. Khan, in: M.V. Nathal, R. Darolia, C.T. Liu, P.L. Martin, D.B. Miracle, R. Wagner, M. Yamaguchi (Eds.), Structural Intermetallics, TMS, 1997, p. 313. [7] T. Kawabata, T. Tamura, O. Izumi, Metall. Trans. 24A (1993) 141. [8] S.C. Huang, E.L. Hall, Acta Metall. Mater. 39 (1991) 1053. [9] T. Noda, M. Okable, S. Isobe, M. Sayashi, Mater. Sci. Eng., A 192/ 193 (1995) 774. [10] W.H. Tian, M. Nemoto, in: Y.W. Kim, R. Wagner, M. Yamaguchi (Eds.), Gamma Titanium Aluminides, TMS, 1995, p. 689. [11] D.M. Dimiduk, V.K. Vasudevan, in: Y.W. Kim, D.M. Dimiduk, M.H. Loretto (Eds.), Gamma Titanium Aluminides, TMS, 1999, p. 239. [12] D.M. Dimiduk, P.L. Martin, Y.W. Kim, Mater. Sci. Eng., A 243 (1998) 66. [13] Y.L. Hao, D.S. Xu, Y.Y. Cui, R. Yang, D. Li, Acta Mater. 47 (1999) 1129.