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The effects of interfaces on the mechanical properties of Ni–Al–Fe intermetallics
Materials Chemistry and Physics 75 (2002) 296–300
The effects of interfaces on the mechanical properties of Ni–Al–Fe intermetallics Chun-Huei Tsau Institute of Materials Science and Manufacturing, Chinese Culture University, Taipei, Taiwan ROC
Abstract The casting Ni–25Al–xFe alloys had dendritic microstructures, and the two major phases existing in these alloys were B2 (ordered b.c.c.) phase in dendrites and f.c.c. phase in interdendrites. The B2 phase exhibited a higher hardness and was more brittle by comparison with the f.c.c. phase in interdendrites, and the two phases showed different plastic deformation results during tensile test. Cracks nucleated at the interfaces between these two phases during tensile stress, then propagated and the specimen failed. The room temperature tensile fracture modes were a mixture of dimple and cleavage. Changing the testing temperature or the mole fraction of elements in the alloys changed the fracture mode and mechanical properties. © 2002 Elsevier Science B.V. All rights reserved. Keywords: Interfaces; Dendrites; Ni–Al–Fe intermetallics
1. Introduction The casting Ni–25Al–xFe alloys (x is from 20 to 32.5 at.%) possessed typical dendritic microstructures [1–3], and the major phases existing in the as-annealed Ni–25Al– xFe alloys were B2 (ordered b.c.c.) and f.c.c. phases. The dendrites were B2 phase and thus with intrinsic brittle and hard manner. The f.c.c. phase was matrix of the interdendritic regions. Since f.c.c. phase has 12 slip systems, it supported major plastic deformation during tensile testing and was well studied [2]. In addition, a third phase with b.c.c.-structured coherent precipitates, which were precipitated from B2-dendrites during annealing, was observed in Ni–25Al–xFe with higher Fe-content because it was a Fe-rich phase [3]. These b.c.c. precipitates had no contribution to tensile properties, because they could not support any plastic deformation or change the path of crack propagation. Therefore, these b.c.c. precipitates were not included in the present study. The two major phases (B2 and f.c.c.) had different mechanical behaviors during plastic deformation. Beside their intrinsic characteristics, the interfaces between them also played a very important role on mechanical behaviors. The orientation relationships between the adjacent f.c.c. and B2 phases were common, predominantly the Kurdjumov–Sachs
E-mail address: [email protected] (C.-H. Tsau).
relationship [4] listed by Eqs. (1) and (2): [1 1 1]B2//[1 1 0]f.c.c.
(1)
and (1¯ 1 0)B2//(1¯ 1 1)f.c.c.
(2)
Therefore, the interfaces are very important in controlling and improving the mechanical properties. Moreover, temperature also had strong effects on the fracture mechanism. The present study investigates the relationships among the microstructures, temperature and mechanical behaviors.
2. Experimental The Ni–25Al–xFe alloys, where x is from 20 to 32.5 at.%, used in this study were prepared by arc melting by using the appropriate amounts of 99.9% pure elements, and drop casting into ingots (10 mm × 30 mm × 100 mm) in an argon atmosphere. The castings were annealed at 1000 ◦ C in vacuum (< 3 × 10−4 Pa) for 4 h to eliminate the residual stress and microsegregation. Since the weight losses of the alloys during melting were all less than 0.2 wt.% of the original weight, the chemical compositions were not analyzed. Tensile specimens were machined from the as-annealed ingots with a gauge section of 2 mm ×3 mm ×25 mm. Tensile tests were performed on a Gleeble 2000 high temperature deformation simulation and mechanical testing machine, which
0254-0584/02/$ – see front matter © 2002 Elsevier Science B.V. All rights reserved. PII: S 0 2 5 4 - 0 5 8 4 ( 0 2 ) 0 0 0 7 9 - 2
C.-H. Tsau / Materials Chemistry and Physics 75 (2002) 296–300
were operated at a constant cross-head speed with a strain rate of 1 × 10−3 s−1 . 3. Results and discussion The as-cast Ni–25Al–xFe alloys show typical dendritic microstructures. By referring our previous studies [1–3] and the ternary phase diagrams [5], the three phases existing in the as-cast Ni–25Al–xFe alloys are B2 phase (the dendrites and the particles embedded in the interdendritic regions), f.c.c. phase (the matrix of the interdendrite) and L12 phase (the coherent precipitates in the f.c.c.-phase interdendrite with an average diameter of about 10 nm), followed by annealing to eliminate the residual stress and microsegregation. The annealing treatment decreased the ordering of these alloys [1,2], because L12 -phase precipitates dissolve into f.c.c.-matrix as observed by TEM; and the B2-dendrite also has a lower ordering by comparing their XRD patterns. Therefore, the interdendritic regions of Ni–25Al–xFe alloys changed from a state of three phases (B2-phase particles + f.c.c.-phase matrix + L12 -phase precipitates) to a two-phase state (B2-phase particles + f.c.c.-phase matrix). In addition, the dendrites in these alloys with Fe-content more than 27.5 at.% also changed from a single B2 phase state to a dual-phase state (B2-phase dendrite + b.c.c.-phase coherent precipitates). Consequently, two different interfaces existed in these as-annealed alloys. First is the B2/f.c.c. interface, the interface between B2-dendrite and interdendritic region, and also the interface between f.c.c.-matrix and B2-particle in interdendritic region. Referring to the Kurdjumov–Sachs relationship described by Eqs. (1) and (2), it was observed that the interface regions exhibited a high density of interfacial dislocation, as shown in Fig. 1. Additionally, the second is the coherent b.c.c./B2 interface, the interface between
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B2-dendrite and b.c.c.-precipitate embedded in the dendrites for the alloys with Fe-content more than 27.5 at.% [3]. As mentioned above, this type of interface had no important effect on the mechanical properties. Tensile properties were only tested on the as-annealed Ni–25Al–xFe alloys, because the as-cast alloys possessed the stress concentration and microsegregation. The results of tensile tests performed on each of the as-annealed alloys plotted as a function of temperature are shown in Fig. 2. While the alloys were tested at room temperature, Ni–25Al– 27.5Fe alloys possessed the best fracture elongation and ultimate tensile strength, but it had the lowest yield strength among these five alloys. This phenomenon is caused by the dimensions of f.c.c.-interdendritic regions, i.e. the main plastic zones in these alloys, influenced on the state of stress [2]. The yield strength for all the five alloys decreased with the increasing test temperature, but the ultimate tensile strength and fracture elongation showed a quite different tendency. Both the ultimate tensile strength and elongation increased with increasing test temperature from room temperature to 200 ◦ C, and these increments were due to the softening of B2-dendrites in the alloys. Since the B2 phase is well known to have a brittle manner at room temperature and the ductile-to-brittle temperature (DBTT) of B2–NiAl is around 200–400 ◦ C, depending on the crystal orientation [6]. Therefore, increasing the test temperature results in an increase in fracture elongation. Especially the alloys that were brittle at room temperature show more increments in fracture elongation with increasing temperature. Further increasing the testing temperature, the fracture elongation of these five alloys decreased significantly. The brittle behavior maintained to about 400–600 ◦ C, when the fracture elongation for Ni–25Al–xFe alloys started to increase again up to 700 ◦ C, due to the change in failure mechanism. Therefore,
Fig. 1. A TEM BF image showing the interfacial dislocations arrayed on the B2/f.c.c. interface.
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Fig. 3. The stress–strain curves of Ni–25Al–25Fe alloy tested at variation temperatures.
Fig. 2. Plots of (a) yield strength, (b) ultimate tensile strength, and (c) fracture elongation as a function of temperature for the Ni–25Al–xFe alloys.
the failures were dominated by the strength of phases and interfaces between these phases. The effect of temperature on tensile behaviors is also revealed by the stress–strain curves tested at variation temperature. Fig. 3 is a typical example for Ni–25Al–25Fe alloy. It had a low room temperature tensile elongation of 4.5%. The tensile behavior improves with increasing test temperature. This is because of the softening of B2 phase at
high temperatures. In addition, the stiffness (i.e. the slope of the stress–strain curve) of Ni–25Al–25Fe alloy significantly decreased from room temperature to 400 ◦ C. The decrease of fracture elongation was found over the temperature range 400–600 ◦ C. This mid-temperature brittle manner is a result of the competition of cleavage of B2 phase and the interfacial rupture of B2/f.c.c. interface. But the stiffness almost remained the same from 400 to 600 ◦ C. The fracture is controlled by a high temperature diffusion-controlled failure mechanism at the test temperatures above 600 ◦ C, and the stiffness decreased again when the test temperature was higher than 600 ◦ C. In addition, the stress–strain curves indicate that the specimens were failed at the uniform deformation stage when the test temperature is less than 600 ◦ C. In other words, no necking was found in these specimens after the tensile test at temperatures lower than 600 ◦ C. The specimen became very soft and had good fracture elongation when tested at temperature up to 800 ◦ C. This phenomenon was also observed for the other four alloys. The fracture surfaces of Ni–25Al–25Fe alloy tested at variation temperatures are shown in Fig. 4. Their morphologies show a mixture of dimples and some transgranular cleavage at low testing temperatures. The dimples were formed either from f.c.c. matrix and B2 particles in the interdendritic regions, or from the tearing of interfaces between f.c.c. interdendrites and B2 dendrites. The transgranular cleavage was formed by cracks passing through the B2 dendrite. Because B2 phase has an intrinsic brittle manner at low temperatures and f.c.c. phase has 12 slip systems, the plastic deformation was mainly supported by the f.c.c. phase in the interdendritic regions. During tensile testing, the fracture nucleated at the interfaces between f.c.c. and B2 phases, because of their different mechanical behaviors. The crack then propagated along the interdendritic regions and formed the dimple fracture surfaces. In addition, cracks propagated across the dendrite and formed the cleavage fracture surfaces. The fracture surfaces of Ni–25Al–25Fe tested at room temperature showed a small fraction of dimples and a
C.-H. Tsau / Materials Chemistry and Physics 75 (2002) 296–300
Fig. 4. Fracture surfaces of Ni–25Al–25Fe alloy tested at (a) room temperature; (b) 200 ◦ C; (c) 400 ◦ C; (d) 500 ◦ C; (e) 600 ◦ C; and (f) 700 ◦ C.
Fig. 5. Profiles of fracture surfaces of Ni–25Al–25Fe alloy tested at (a) room temperature; (b) 200 ◦ C; (c) 500 ◦ C; and (d) 600 ◦ C.
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large area of cleavage. This is also supported by the Ni–25Al–25Fe only having 4.5% of fracture elongation. Because B2 phase became softer and exhibited some plastic deformation with increasing test temperature, the fracture surface of Ni–25Al–25Fe alloy tested at 200 ◦ C thus had more dimple fracture areas. This is also shown by the increase in fracture elongation for Ni–25Al–25Fe alloy tested at this temperature range. This fracture surface indicates that the fracture mode of 400 ◦ C is quite different to that of room temperature. The same fracture surface is also observed with the alloy tested at 600 ◦ C. The fracture mode changed again when test temperature is above 700 ◦ C. The tearing at the interfaces is more significant. The interface became weaker with increasing temperature, and the fracture surface showed an apparent dendrite morphology. In other words, crack path propagated almost fully along the interface. The profiles of fracture surfaces of Ni–25Al–25Fe alloy are shown in Fig. 5. Cracks occurred at the interfaces between f.c.c. interdendrites and B2 dendrites and B2 dendrites emerged on the fracture surfaces are apparent. These profiles also reveal the effect of test temperature on the fracture modes. At low temperatures, crack propagates not only along the B2/f.c.c. interfaces, but also across the B2 dendrites. However, crack paths are almost fully along the interfaces at high temperatures. The interfaces between the f.c.c. and B2 phases are always the crack nucleation sites and the major propagation paths. That is, the interfaces between the f.c.c. and B2 phases are the weakness of the alloy system. 4. Conclusions In Ni–25Al–xFe alloy system, the morphologies of microstructure strongly affect the room temperature tensile properties. B2/f.c.c. interfaces are observed as the major
crack nucleation sites during tensile testing due to the incompatibility between these two phases. In addition, temperature exerts a strong influence on the tensile properties of these alloys. Yield strength drops remarkably near 100 ◦ C. Both the ultimate strength and elongation show a significant decrease over the temperature range 400–600 ◦ C. The phenomenon can be explained by the competition of cleavage fracture of B2 dendrite and rupture of B2/f.c.c. interface. Failure of Ni–25Al–xFe alloys occurred at the uniform deformation stage, while necking was observed at test temperatures above 600 ◦ C. The interfaces between the f.c.c. and B2 phases are always crack nucleation sites and the major propagation paths.
Acknowledgements The authors are grateful to the National Science Council of ROC for the financial support under contract No. 892216-E-034-005. References [1] C.-H. Tsau, J.S.-C. Jang, J.-W. Yeh, Mater. Sci. Eng. A 152 (1992) 264. [2] C.-H. Tsau, J.-W. Yeh, Mater. Chem. Phys. 68 (2001) 142. [3] C.-H. Tsau, J.S.-C. Jang, J.-W. Yeh, Scripta Metall. Mater. 34 (1996) 325. [4] S.C. Huang, R.D. Field, D.D. Krueger, Metall. Trans. A 21 (1990) 959. [5] G.V. Raynor, V.G. Rivlin, Phase Equilibrium in Iron Ternary Alloys, No. 4, The Institute of Metals, London, 1988, p. 107. [6] D.F. Lahrman, R.D. Field, R. Darolia, The effect of strain rate on the mechanical properties of single crystal NiAl, in: L.A. Johnson, D.P. Pope, J.O. Stiegler (Eds.), High Temperature Ordered Intermetallic Alloys IV, MRS, Pittsburgh, PA, 1991, p. 603.
The effects of interfaces on the mechanical properties of Ni–Al–Fe intermetallics Chun-Huei Tsau Institute of Materials Science and Manufacturing, Chinese Culture University, Taipei, Taiwan ROC
Abstract The casting Ni–25Al–xFe alloys had dendritic microstructures, and the two major phases existing in these alloys were B2 (ordered b.c.c.) phase in dendrites and f.c.c. phase in interdendrites. The B2 phase exhibited a higher hardness and was more brittle by comparison with the f.c.c. phase in interdendrites, and the two phases showed different plastic deformation results during tensile test. Cracks nucleated at the interfaces between these two phases during tensile stress, then propagated and the specimen failed. The room temperature tensile fracture modes were a mixture of dimple and cleavage. Changing the testing temperature or the mole fraction of elements in the alloys changed the fracture mode and mechanical properties. © 2002 Elsevier Science B.V. All rights reserved. Keywords: Interfaces; Dendrites; Ni–Al–Fe intermetallics
1. Introduction The casting Ni–25Al–xFe alloys (x is from 20 to 32.5 at.%) possessed typical dendritic microstructures [1–3], and the major phases existing in the as-annealed Ni–25Al– xFe alloys were B2 (ordered b.c.c.) and f.c.c. phases. The dendrites were B2 phase and thus with intrinsic brittle and hard manner. The f.c.c. phase was matrix of the interdendritic regions. Since f.c.c. phase has 12 slip systems, it supported major plastic deformation during tensile testing and was well studied [2]. In addition, a third phase with b.c.c.-structured coherent precipitates, which were precipitated from B2-dendrites during annealing, was observed in Ni–25Al–xFe with higher Fe-content because it was a Fe-rich phase [3]. These b.c.c. precipitates had no contribution to tensile properties, because they could not support any plastic deformation or change the path of crack propagation. Therefore, these b.c.c. precipitates were not included in the present study. The two major phases (B2 and f.c.c.) had different mechanical behaviors during plastic deformation. Beside their intrinsic characteristics, the interfaces between them also played a very important role on mechanical behaviors. The orientation relationships between the adjacent f.c.c. and B2 phases were common, predominantly the Kurdjumov–Sachs
E-mail address: [email protected] (C.-H. Tsau).
relationship [4] listed by Eqs. (1) and (2): [1 1 1]B2//[1 1 0]f.c.c.
(1)
and (1¯ 1 0)B2//(1¯ 1 1)f.c.c.
(2)
Therefore, the interfaces are very important in controlling and improving the mechanical properties. Moreover, temperature also had strong effects on the fracture mechanism. The present study investigates the relationships among the microstructures, temperature and mechanical behaviors.
2. Experimental The Ni–25Al–xFe alloys, where x is from 20 to 32.5 at.%, used in this study were prepared by arc melting by using the appropriate amounts of 99.9% pure elements, and drop casting into ingots (10 mm × 30 mm × 100 mm) in an argon atmosphere. The castings were annealed at 1000 ◦ C in vacuum (< 3 × 10−4 Pa) for 4 h to eliminate the residual stress and microsegregation. Since the weight losses of the alloys during melting were all less than 0.2 wt.% of the original weight, the chemical compositions were not analyzed. Tensile specimens were machined from the as-annealed ingots with a gauge section of 2 mm ×3 mm ×25 mm. Tensile tests were performed on a Gleeble 2000 high temperature deformation simulation and mechanical testing machine, which
0254-0584/02/$ – see front matter © 2002 Elsevier Science B.V. All rights reserved. PII: S 0 2 5 4 - 0 5 8 4 ( 0 2 ) 0 0 0 7 9 - 2
C.-H. Tsau / Materials Chemistry and Physics 75 (2002) 296–300
were operated at a constant cross-head speed with a strain rate of 1 × 10−3 s−1 . 3. Results and discussion The as-cast Ni–25Al–xFe alloys show typical dendritic microstructures. By referring our previous studies [1–3] and the ternary phase diagrams [5], the three phases existing in the as-cast Ni–25Al–xFe alloys are B2 phase (the dendrites and the particles embedded in the interdendritic regions), f.c.c. phase (the matrix of the interdendrite) and L12 phase (the coherent precipitates in the f.c.c.-phase interdendrite with an average diameter of about 10 nm), followed by annealing to eliminate the residual stress and microsegregation. The annealing treatment decreased the ordering of these alloys [1,2], because L12 -phase precipitates dissolve into f.c.c.-matrix as observed by TEM; and the B2-dendrite also has a lower ordering by comparing their XRD patterns. Therefore, the interdendritic regions of Ni–25Al–xFe alloys changed from a state of three phases (B2-phase particles + f.c.c.-phase matrix + L12 -phase precipitates) to a two-phase state (B2-phase particles + f.c.c.-phase matrix). In addition, the dendrites in these alloys with Fe-content more than 27.5 at.% also changed from a single B2 phase state to a dual-phase state (B2-phase dendrite + b.c.c.-phase coherent precipitates). Consequently, two different interfaces existed in these as-annealed alloys. First is the B2/f.c.c. interface, the interface between B2-dendrite and interdendritic region, and also the interface between f.c.c.-matrix and B2-particle in interdendritic region. Referring to the Kurdjumov–Sachs relationship described by Eqs. (1) and (2), it was observed that the interface regions exhibited a high density of interfacial dislocation, as shown in Fig. 1. Additionally, the second is the coherent b.c.c./B2 interface, the interface between
297
B2-dendrite and b.c.c.-precipitate embedded in the dendrites for the alloys with Fe-content more than 27.5 at.% [3]. As mentioned above, this type of interface had no important effect on the mechanical properties. Tensile properties were only tested on the as-annealed Ni–25Al–xFe alloys, because the as-cast alloys possessed the stress concentration and microsegregation. The results of tensile tests performed on each of the as-annealed alloys plotted as a function of temperature are shown in Fig. 2. While the alloys were tested at room temperature, Ni–25Al– 27.5Fe alloys possessed the best fracture elongation and ultimate tensile strength, but it had the lowest yield strength among these five alloys. This phenomenon is caused by the dimensions of f.c.c.-interdendritic regions, i.e. the main plastic zones in these alloys, influenced on the state of stress [2]. The yield strength for all the five alloys decreased with the increasing test temperature, but the ultimate tensile strength and fracture elongation showed a quite different tendency. Both the ultimate tensile strength and elongation increased with increasing test temperature from room temperature to 200 ◦ C, and these increments were due to the softening of B2-dendrites in the alloys. Since the B2 phase is well known to have a brittle manner at room temperature and the ductile-to-brittle temperature (DBTT) of B2–NiAl is around 200–400 ◦ C, depending on the crystal orientation [6]. Therefore, increasing the test temperature results in an increase in fracture elongation. Especially the alloys that were brittle at room temperature show more increments in fracture elongation with increasing temperature. Further increasing the testing temperature, the fracture elongation of these five alloys decreased significantly. The brittle behavior maintained to about 400–600 ◦ C, when the fracture elongation for Ni–25Al–xFe alloys started to increase again up to 700 ◦ C, due to the change in failure mechanism. Therefore,
Fig. 1. A TEM BF image showing the interfacial dislocations arrayed on the B2/f.c.c. interface.
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Fig. 3. The stress–strain curves of Ni–25Al–25Fe alloy tested at variation temperatures.
Fig. 2. Plots of (a) yield strength, (b) ultimate tensile strength, and (c) fracture elongation as a function of temperature for the Ni–25Al–xFe alloys.
the failures were dominated by the strength of phases and interfaces between these phases. The effect of temperature on tensile behaviors is also revealed by the stress–strain curves tested at variation temperature. Fig. 3 is a typical example for Ni–25Al–25Fe alloy. It had a low room temperature tensile elongation of 4.5%. The tensile behavior improves with increasing test temperature. This is because of the softening of B2 phase at
high temperatures. In addition, the stiffness (i.e. the slope of the stress–strain curve) of Ni–25Al–25Fe alloy significantly decreased from room temperature to 400 ◦ C. The decrease of fracture elongation was found over the temperature range 400–600 ◦ C. This mid-temperature brittle manner is a result of the competition of cleavage of B2 phase and the interfacial rupture of B2/f.c.c. interface. But the stiffness almost remained the same from 400 to 600 ◦ C. The fracture is controlled by a high temperature diffusion-controlled failure mechanism at the test temperatures above 600 ◦ C, and the stiffness decreased again when the test temperature was higher than 600 ◦ C. In addition, the stress–strain curves indicate that the specimens were failed at the uniform deformation stage when the test temperature is less than 600 ◦ C. In other words, no necking was found in these specimens after the tensile test at temperatures lower than 600 ◦ C. The specimen became very soft and had good fracture elongation when tested at temperature up to 800 ◦ C. This phenomenon was also observed for the other four alloys. The fracture surfaces of Ni–25Al–25Fe alloy tested at variation temperatures are shown in Fig. 4. Their morphologies show a mixture of dimples and some transgranular cleavage at low testing temperatures. The dimples were formed either from f.c.c. matrix and B2 particles in the interdendritic regions, or from the tearing of interfaces between f.c.c. interdendrites and B2 dendrites. The transgranular cleavage was formed by cracks passing through the B2 dendrite. Because B2 phase has an intrinsic brittle manner at low temperatures and f.c.c. phase has 12 slip systems, the plastic deformation was mainly supported by the f.c.c. phase in the interdendritic regions. During tensile testing, the fracture nucleated at the interfaces between f.c.c. and B2 phases, because of their different mechanical behaviors. The crack then propagated along the interdendritic regions and formed the dimple fracture surfaces. In addition, cracks propagated across the dendrite and formed the cleavage fracture surfaces. The fracture surfaces of Ni–25Al–25Fe tested at room temperature showed a small fraction of dimples and a
C.-H. Tsau / Materials Chemistry and Physics 75 (2002) 296–300
Fig. 4. Fracture surfaces of Ni–25Al–25Fe alloy tested at (a) room temperature; (b) 200 ◦ C; (c) 400 ◦ C; (d) 500 ◦ C; (e) 600 ◦ C; and (f) 700 ◦ C.
Fig. 5. Profiles of fracture surfaces of Ni–25Al–25Fe alloy tested at (a) room temperature; (b) 200 ◦ C; (c) 500 ◦ C; and (d) 600 ◦ C.
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large area of cleavage. This is also supported by the Ni–25Al–25Fe only having 4.5% of fracture elongation. Because B2 phase became softer and exhibited some plastic deformation with increasing test temperature, the fracture surface of Ni–25Al–25Fe alloy tested at 200 ◦ C thus had more dimple fracture areas. This is also shown by the increase in fracture elongation for Ni–25Al–25Fe alloy tested at this temperature range. This fracture surface indicates that the fracture mode of 400 ◦ C is quite different to that of room temperature. The same fracture surface is also observed with the alloy tested at 600 ◦ C. The fracture mode changed again when test temperature is above 700 ◦ C. The tearing at the interfaces is more significant. The interface became weaker with increasing temperature, and the fracture surface showed an apparent dendrite morphology. In other words, crack path propagated almost fully along the interface. The profiles of fracture surfaces of Ni–25Al–25Fe alloy are shown in Fig. 5. Cracks occurred at the interfaces between f.c.c. interdendrites and B2 dendrites and B2 dendrites emerged on the fracture surfaces are apparent. These profiles also reveal the effect of test temperature on the fracture modes. At low temperatures, crack propagates not only along the B2/f.c.c. interfaces, but also across the B2 dendrites. However, crack paths are almost fully along the interfaces at high temperatures. The interfaces between the f.c.c. and B2 phases are always the crack nucleation sites and the major propagation paths. That is, the interfaces between the f.c.c. and B2 phases are the weakness of the alloy system. 4. Conclusions In Ni–25Al–xFe alloy system, the morphologies of microstructure strongly affect the room temperature tensile properties. B2/f.c.c. interfaces are observed as the major
crack nucleation sites during tensile testing due to the incompatibility between these two phases. In addition, temperature exerts a strong influence on the tensile properties of these alloys. Yield strength drops remarkably near 100 ◦ C. Both the ultimate strength and elongation show a significant decrease over the temperature range 400–600 ◦ C. The phenomenon can be explained by the competition of cleavage fracture of B2 dendrite and rupture of B2/f.c.c. interface. Failure of Ni–25Al–xFe alloys occurred at the uniform deformation stage, while necking was observed at test temperatures above 600 ◦ C. The interfaces between the f.c.c. and B2 phases are always crack nucleation sites and the major propagation paths.
Acknowledgements The authors are grateful to the National Science Council of ROC for the financial support under contract No. 892216-E-034-005. References [1] C.-H. Tsau, J.S.-C. Jang, J.-W. Yeh, Mater. Sci. Eng. A 152 (1992) 264. [2] C.-H. Tsau, J.-W. Yeh, Mater. Chem. Phys. 68 (2001) 142. [3] C.-H. Tsau, J.S.-C. Jang, J.-W. Yeh, Scripta Metall. Mater. 34 (1996) 325. [4] S.C. Huang, R.D. Field, D.D. Krueger, Metall. Trans. A 21 (1990) 959. [5] G.V. Raynor, V.G. Rivlin, Phase Equilibrium in Iron Ternary Alloys, No. 4, The Institute of Metals, London, 1988, p. 107. [6] D.F. Lahrman, R.D. Field, R. Darolia, The effect of strain rate on the mechanical properties of single crystal NiAl, in: L.A. Johnson, D.P. Pope, J.O. Stiegler (Eds.), High Temperature Ordered Intermetallic Alloys IV, MRS, Pittsburgh, PA, 1991, p. 603.