- Email: [email protected]
Effect of alloying additions on the κ phase precipitation in austenitic Fe-Mn-Al-C alloys
Scripta METALLURGICA et MATERIALIA
Vol. 30, pp. 499-504, 1994 Printed in the U.S.A.
Pergamon Press Ltd. All rights reserved
EFFECT OF ALLOYING ADDITIONS ON THE lc PHASE PRECIPITATION IN AUSTENITIC Fe-Mn-A1-C ALLOYS H. Huang, D. Gan and P. W. Kao Institute of Materials Science and Engineering National Sun Yat-Sen University Kaohsiung, Taiwan (Received July 12, 1993) (Revised November 15, 1993) Introduction
Austenitic Fe-A1-Mn-C alloys are developed as a possible substitute for the Ni-Cr stainless steel and have shown some promise as engineering alloys for their high strength, low density and low cost [1-4]. Aluminum is added to increase the corrosion and oxidation resistances. Manganese is added to stabilize the austenltic structure but its content should be less than 35 wt% to avoid the precipitation of brittle [~-Mn phase [5]. Carbon is a potent austenite former and strengthener for this alloy system. Wang and Beck [6] suggested a composition range of 20-30 wt% Mn, less than 12 wt% A1 and up to 1 wt% C. Since aluminum is a ferrite former, higher carbon and manganese contents usually accompany an increase in aluminum content in order to stabilize the austenitic structure. A clear composition boundary for single phase austenite has not yet been established in this alloy system. Based on the limited data given in the literature [7-11], however, an approximate boundary for ~'/(x+y may be drawn for Fe-Mn-A1-C alloys with about 30 wt% Mn, which is shown in Figure 1. Aging of austenitic Fe-Mn-Al-C alloys at temperatures from 500 to 750°C generally causes the precipitation of the ic phase, (Fe,Mn)3A1Cx, and has been studied in many reports [2, 12-21]. Aging of Fe-3OMn-10AI-IC-1Si alloy at 550 - 700°C causes extensive precipitation of the 1( phase at austenite grain boundaries as well as within austenite grains [2, 13], which will be named as intergranular and intragranular ~c phase respectively in the following text. The formation of fine intragranular ~c phase is considered to be a possible hardening mechanism for this alloy system. However, it has been shown [2] that the presence of intragranular ~: phase can cause brittle fracture before yielding at subzero temperatures. On the other hand, the intergranular !c phase alone can result in a severe loss in impact energy at both room and subzero temperatures [12]. Owing to the importance of the precipitation of ~: phase on the mechanical properties, a better knowledge of the alloying effect on the precipitation is important for the further development of this alloy system. In the present work, a systematic study of the effect of alloying elements, aluminum and carbon in particular, on the precipitation of ~ phase was carried out over a range of composition within which a single phase austenite is formed upon solution treating and quenching.
499 0956-716X/94 $6.00 + .00 Copyright (c) 1993 Pergamon Press Ltd.
500
PRECIPITATION IN Fe-Mn-A1-C
Vol.
30, No. 4
Exverimental Procedure
In this study, two groups of alloys, A (~ 1 wt% C) and B (- 5 wt% Al), were prepared mainly to investigate the effect of aluminum and carbon, respectively. Because most of the austenitic FeMn-AI-C alloys reported in the literature contain 28-34 wt% Mn, the manganese content is set near 30 wt% in this study except for alloys A3 and AS, which are designed to check the manganese effect by comparing with A2 and A4, respectively. Alloys used in this experiment were prepared from high purity iron, manganese, aluminum, silicon and graphite by induction melting in an argon atmosphere. The molten metal was cast in a steel mould and hot forged at 1200°C into billet. After being homogenized at 1100°C for 2 h, it was hot roiled to a plate 12 mm thick. The chemical compositions of the alloys are shown in Table 1. TABLE 1. Chemical Compositions of the Alloys (wt%) and the Precipitation of the !¢Phase as a Function of Alloy composition. Alloy
A1 A2 A3 A4 A5 A6 B1 B2 B3 B4
Fe
66.12 63.32 66.61 62.90 68.22 58.58 63.81 61.93 60.83 61.21
Mn
30.70 30.40 27.11 29.21 23.46 29.53 30.71 32.07 32.84 32.43
AI
2.28 5.32 5.33 6.95 7.38 10.00 5.08 5.43 5.56 5.32
C
0.90 0.96 0.95 0.94 0.94 1.03 0.40 0.57 0.77 1.04
Si
-----0.86 -~ ~ ~
Intergranular
Intragranular
K phase
K phase
yes
no no no yes yes yes no no no no
yes yes yes yes yes no no yes yes
ao(nm)
0.3628 0.3658 0.3662 0.3678 0.3692 0.3692 0.3647 0.3653 0.3660 0.3667
All specimens except alloy A6 were solution treated at 1150°C for 2 h and oil quenched. Alloy A6 in which the lc phase precipitates very fast was water quenched. A single phase austenitic structure was obtained for all the specimens after solution treatment; this was confirmed by metallographic examination and by x-ray diffraction. A Diano XRD-8536 diffractometer equipped with a diffracted-beam graphite monochromator and Cu-K(x radiation was used for x-ray analysis. Lattice constants were calculated with the Nelson-Riley extrapolation method. The accuracy was found better than + 0.0003 nm. The specimens were aged at 650°C for up to 360 h and the aging temperature was controlled within + 3°C. This temperature was selected because it is near the nose of the time-temperaturetransformation curve. After aging, specimens were polished, etched with 10% nital solution, and studied with an optical microscope and a scanning electron microscope (SEM, JEOL JSM-35CF). Thin specimens were also prepared for observation under a transmission electron microscope (TEM, JEOL JEM-200CX).
Vol.
30, No. 4
PRECIPITATION
IN Fe-Mn-AI-C
501
Results and Discussion The precipitation of ~: phase as a function of alloy chemistry in Fe-Mn-AI-C alloys are shown in Table 1. The lattice constants for the alloys in solution-treated condition, which were determined by x-ray diffraction are also included in Table 1. A. Intergranular ~ phase The intergranular precipitates are identified to be the same as the intragranular ~ phase by electron diffraction in TEM. The int~granular K phase develops originally as discrete particles but soon becomes nearly continuous along the grain boundary. Results of the alloying effect on the precipitation of intergranular ic ptxase at 650°C are shown in Table I. By examining the group A alloys in Table I, it is found that the intergranular i¢ phase appears in all the alloys with about 1 wt% carbon irrespective of the aluminum content. From the results of group B alloys, one may find that for Fe-30Mn-SA1 alloys, 0.77 wt% carbon is necessary for the precipitation of intergranular ic phase at 650oC. Comparing the results of A2 and A3 as well as those of A4 and A5, one can find that the manganese content appears to be not an important factor in determining the precipitation of intergranular ~ phase. B. lntragranular • phase The results of the precipitation of il~ragranular !c phase, identified by electron diffraction in TEM, are shown in Table 1. It has been shown [14, 15, 20] that the intragranular Ic phase maintains a parallel orientation relationship ([100]//[100], [010]//[010]) with the austenite matrix, is essentially cubic in shape, and has {100} type interfaces with the austenlte matrix. The TEM micrograph and diffraction pattern of early stage precipitation of i< phase in A6 specimen are shown in Figure 2. In group A alloys, which contain about 1 wt% carbon, the precipitation of intragranular g phase is observed only in alloys with higher aluminum contents, i.e. > 6.95 wt%. On the other hand, in group B alloys with lower aluminum content of -5.5 wt%, the intragranular i¢ phase does not precipitate in alloys irrespective of the carbon content. By comparing alloys A2 and A3, and A4 and A5, one may find that within the range studied, the content of Mn appears to be not an important factor in determining the precipitation of intragranular ~ phase. By using electrolytical extraction and Debye-Scherrer x-ray diffraction, the lattice constant of intragranular ~ phase extracted from alloy A6 has been determined to be 0.3747 nm, which is larger than the lattice constant of the austenltic Fe-Mn-A1-C alloys in Table 1. Since the intragranular !c phase and the matrix have a parallel orientation relationship and acommon {100} type interfaces [13-15, 20], the lattice misfit between the !¢ phase and the matrix will result in a considerable misfit strain energy. The misfit strain energy should be an important factor in determining the precipitation of intragranular ~ phase. Based upon the results of intra- and intergranular 1¢phase precipitation, it is clear that decreasing the aluminum and carbon content will decrease the chemical driving force for the precipitation of ~ phase. At the same time, this will also decrease the lattice constant of austenite which in turn will increase the misfit between austenite and intragranular 1¢ phase and then the coherent strain energy. Both the decrease in chemical driving force and the increase in strain energy, are unfavorable for the precipitation of intragranular I¢ phase. Therefore, it is possible that the lattice constant of austenite matrix may serve as a simple empirical index in determining the precipitation of intragranular ic phase.
502
PRECIPITATION
IN Fe-Mn-AI-C
Vol. 30, No.
To study this possibility, the measured lattice constants for these Fe-Mn-A1-C alloys are also listed in Table 1. Based on Table 1, a value of -0.3670 nm may be considered as a critical lattice constant for the austenite matrix, below which the intragranular i¢ phase will not develop. The lattice constant ao of austenitic Fe-Mn-A1-C alloys has been determined to be a linear function of chemical composition (in nm and wt%) [22], a o = 0.3574 + 0.000052 Mn + 0.00094 A1 + 0.0020 C - 0.00098 Si. By use of the critical lattice'constant (-0.3670 nm) and eqn. [1], a composition boundary for the precipitation of intragranular !¢ phase in Fe-Mn-AI-C alloys can be expressed as (in wt%) 0.098 AI + 0.208 C = 1 - 0.0054 Mn [2] According to eqn. [2], the composition boundaries for the precipitation of intragranular !c phase in austenitic Fe-Mn-A1-C alloys containing 25-34 wt% Mn are plotted in Figure 1 along with the experimental observations. Figure 1 also includes the results of various austenitic Fe-Mn-AI-C alloys reported in the literature, whose lattice constants were calculated by e q ~ [1]. The predicted boundary agrees quite well with the experimental results. It is to be noted that the change Of manganese content imposes only a small shift on the boundary line.
The precipitation of 1¢phase in austenitic Fe-Mn-A1-C alloys at 650°C was studied on selected alloys. Experimental results indicate that both aluminum and carbon contents are significant in determining the precipitation of these second phases, while the manganese content shows little effect within the range studied (23-33 wt%). An approximate composition boundary may be established for the precipitation of each phase and is summarized as follows. (1) To avoid the precipitation of intergranular • phase, the alloying addition should be limited to within about 5.5 wt% Al and 0.67 wt% C. (2) To avoid the precipitation of intragranular Ic phase, the alloying addition should be limited to within about 6.2 wt% AI and 1.0 wt% C. Furthermore, it is proposed that as an empirical rule, the lattice constant of austenite matrix may be used as a simple empirical index in determining the precipitation of intragranular ~ phase. A critical lattice constant of about 0.3670 nm has been found for the austenite matrix, above which the intragranular ~ phase does not precipitate. Based on this critical lattice constant, a composition boundary for the precipitation of intragranular K phase in Fe-Mn-AI-C alloys can be derived as (in wt%) 0.098 Al + 0.208 C = 10.0054 Mn. Acknowledgements
This work was financially supported by National Science Council, Taiwan, under contract no. NSC-79-0405-E110-10. Thanks are due to Prof. C. M. Wan for providing the materials used in this work. References
1. S.K. Banerji, Metal Prog., 113 (4), 59 (1978).
4
Vol. 30, No. 4
PRECIPITATION IN Fe-Mn-A1-C
2. R.K. You, PoWe Kao and D. Gala, Mater. Sci. Eng., Al17, 141 (1989). 3. J.L. Ham and R. E. Cairns Jr., Product Eng., 29 (12), 50 (1958). 4. D.J. Schmatz, Trans. ASM, 52, 898 (1960). 5. V.G. Rivlin, Int. Met. Rev., 28 (6), 309 (1983). 6. R. Wang and F. H. Beck, Metal Prog., 123 (4), 72 (1983). 7. J.G. Dub, S. H. Huarng and C. M. Wan, Chinese J. Mater. Sci., 16A (1), 14 (1984). 8. C.P. Chou and C. I-L Lee, scripta Metall., 23, 1109 (1989). 9. G.L. Kayak, Metal Sci. and Heat Treatment, 2, 95 (1969). 10. K. Sato, K. Tagawa and Y. Inoue, Metal1. Trans., 21A, 5 (1990). 11. Y. G. Kim, Y. S. Park and J. K. Han, Metall. Trans. A, 16A, 1689 (1985). 12. K.T. Luo, P. W. Kao and D. Gan, Mater. Sci. Eng., A151, L15 (1992). 13. S. M. Chu, P. W. Kao and D. Gan, Scripta Metall. et Mater., 26, 1067 (1992). 14. K. Sato, K. Tagawa and Y. Inoue, Scripta MetaU., 22, 899 (1988). 15. J. E. Krzanowski, Metall. Trans., 19A, 1873 (1988). 16. N. A. Storchak and A. G. Drachinskaya, Phys. Met. Metall., 44 (2), 123 (1978). 17. G. S. Krivonogov, M. F. Alekseyenko and G. S. Solov'yeva, Phys. Met. Metall., 39 (4), 775 (1975). 18. K. H. Han, W. K. Choo and D. E. Laughlin, Scripta Metall., 22, 1873 (1988). 19. K. H. Han and W. K. Choo, MetaU. Trans., 20A, 205 (1989). 20. K. H. Han, J. C. Yoon and W. K. Choo, Scripta Metall., 20, 33 (1986). 21. K. H. Han and W. K. Choo, MetaU. Trans., 14A, 973 (1983). 22. C. M. Chu, H. Huang, P. W. Kao and D. Gan, submitted to Scripta MetaU. et Mater.. 30Mn
34Mn 1.2
~
25Mn
• ~
1.0
o L O
0
0.8 0.6
N ,it 0.4 oO
0.2
!
!
4
6
8
wt%
Aluminum
0.0 0
2
' I !
10
12
FIG. 1. The precipitation of intragranular K phase as a function of alloy composition. The solid and open symbols indicate the compositions with and without precipitation, respectively. The square symbols indicate the results from the literature [10, 12, 15, 18]. The composition boundaries predicted by eqn. [2] are also indicated.
503
504
PRECIPITATION
IN Fe-Mn-AI-C
Vol.
30, No.
FIG. 2. Centered dark field TEM micrograph and diffraction pattern (zone axis [001]) of ahoy A6 aged at 650 °C for 1 h, showing the intragranular K phase.
4
Vol. 30, pp. 499-504, 1994 Printed in the U.S.A.
Pergamon Press Ltd. All rights reserved
EFFECT OF ALLOYING ADDITIONS ON THE lc PHASE PRECIPITATION IN AUSTENITIC Fe-Mn-A1-C ALLOYS H. Huang, D. Gan and P. W. Kao Institute of Materials Science and Engineering National Sun Yat-Sen University Kaohsiung, Taiwan (Received July 12, 1993) (Revised November 15, 1993) Introduction
Austenitic Fe-A1-Mn-C alloys are developed as a possible substitute for the Ni-Cr stainless steel and have shown some promise as engineering alloys for their high strength, low density and low cost [1-4]. Aluminum is added to increase the corrosion and oxidation resistances. Manganese is added to stabilize the austenltic structure but its content should be less than 35 wt% to avoid the precipitation of brittle [~-Mn phase [5]. Carbon is a potent austenite former and strengthener for this alloy system. Wang and Beck [6] suggested a composition range of 20-30 wt% Mn, less than 12 wt% A1 and up to 1 wt% C. Since aluminum is a ferrite former, higher carbon and manganese contents usually accompany an increase in aluminum content in order to stabilize the austenitic structure. A clear composition boundary for single phase austenite has not yet been established in this alloy system. Based on the limited data given in the literature [7-11], however, an approximate boundary for ~'/(x+y may be drawn for Fe-Mn-A1-C alloys with about 30 wt% Mn, which is shown in Figure 1. Aging of austenitic Fe-Mn-Al-C alloys at temperatures from 500 to 750°C generally causes the precipitation of the ic phase, (Fe,Mn)3A1Cx, and has been studied in many reports [2, 12-21]. Aging of Fe-3OMn-10AI-IC-1Si alloy at 550 - 700°C causes extensive precipitation of the 1( phase at austenite grain boundaries as well as within austenite grains [2, 13], which will be named as intergranular and intragranular ~c phase respectively in the following text. The formation of fine intragranular ~c phase is considered to be a possible hardening mechanism for this alloy system. However, it has been shown [2] that the presence of intragranular ~: phase can cause brittle fracture before yielding at subzero temperatures. On the other hand, the intergranular !c phase alone can result in a severe loss in impact energy at both room and subzero temperatures [12]. Owing to the importance of the precipitation of ~: phase on the mechanical properties, a better knowledge of the alloying effect on the precipitation is important for the further development of this alloy system. In the present work, a systematic study of the effect of alloying elements, aluminum and carbon in particular, on the precipitation of ~ phase was carried out over a range of composition within which a single phase austenite is formed upon solution treating and quenching.
499 0956-716X/94 $6.00 + .00 Copyright (c) 1993 Pergamon Press Ltd.
500
PRECIPITATION IN Fe-Mn-A1-C
Vol.
30, No. 4
Exverimental Procedure
In this study, two groups of alloys, A (~ 1 wt% C) and B (- 5 wt% Al), were prepared mainly to investigate the effect of aluminum and carbon, respectively. Because most of the austenitic FeMn-AI-C alloys reported in the literature contain 28-34 wt% Mn, the manganese content is set near 30 wt% in this study except for alloys A3 and AS, which are designed to check the manganese effect by comparing with A2 and A4, respectively. Alloys used in this experiment were prepared from high purity iron, manganese, aluminum, silicon and graphite by induction melting in an argon atmosphere. The molten metal was cast in a steel mould and hot forged at 1200°C into billet. After being homogenized at 1100°C for 2 h, it was hot roiled to a plate 12 mm thick. The chemical compositions of the alloys are shown in Table 1. TABLE 1. Chemical Compositions of the Alloys (wt%) and the Precipitation of the !¢Phase as a Function of Alloy composition. Alloy
A1 A2 A3 A4 A5 A6 B1 B2 B3 B4
Fe
66.12 63.32 66.61 62.90 68.22 58.58 63.81 61.93 60.83 61.21
Mn
30.70 30.40 27.11 29.21 23.46 29.53 30.71 32.07 32.84 32.43
AI
2.28 5.32 5.33 6.95 7.38 10.00 5.08 5.43 5.56 5.32
C
0.90 0.96 0.95 0.94 0.94 1.03 0.40 0.57 0.77 1.04
Si
-----0.86 -~ ~ ~
Intergranular
Intragranular
K phase
K phase
yes
no no no yes yes yes no no no no
yes yes yes yes yes no no yes yes
ao(nm)
0.3628 0.3658 0.3662 0.3678 0.3692 0.3692 0.3647 0.3653 0.3660 0.3667
All specimens except alloy A6 were solution treated at 1150°C for 2 h and oil quenched. Alloy A6 in which the lc phase precipitates very fast was water quenched. A single phase austenitic structure was obtained for all the specimens after solution treatment; this was confirmed by metallographic examination and by x-ray diffraction. A Diano XRD-8536 diffractometer equipped with a diffracted-beam graphite monochromator and Cu-K(x radiation was used for x-ray analysis. Lattice constants were calculated with the Nelson-Riley extrapolation method. The accuracy was found better than + 0.0003 nm. The specimens were aged at 650°C for up to 360 h and the aging temperature was controlled within + 3°C. This temperature was selected because it is near the nose of the time-temperaturetransformation curve. After aging, specimens were polished, etched with 10% nital solution, and studied with an optical microscope and a scanning electron microscope (SEM, JEOL JSM-35CF). Thin specimens were also prepared for observation under a transmission electron microscope (TEM, JEOL JEM-200CX).
Vol.
30, No. 4
PRECIPITATION
IN Fe-Mn-AI-C
501
Results and Discussion The precipitation of ~: phase as a function of alloy chemistry in Fe-Mn-AI-C alloys are shown in Table 1. The lattice constants for the alloys in solution-treated condition, which were determined by x-ray diffraction are also included in Table 1. A. Intergranular ~ phase The intergranular precipitates are identified to be the same as the intragranular ~ phase by electron diffraction in TEM. The int~granular K phase develops originally as discrete particles but soon becomes nearly continuous along the grain boundary. Results of the alloying effect on the precipitation of intergranular ic ptxase at 650°C are shown in Table I. By examining the group A alloys in Table I, it is found that the intergranular i¢ phase appears in all the alloys with about 1 wt% carbon irrespective of the aluminum content. From the results of group B alloys, one may find that for Fe-30Mn-SA1 alloys, 0.77 wt% carbon is necessary for the precipitation of intergranular ic phase at 650oC. Comparing the results of A2 and A3 as well as those of A4 and A5, one can find that the manganese content appears to be not an important factor in determining the precipitation of intergranular ~ phase. B. lntragranular • phase The results of the precipitation of il~ragranular !c phase, identified by electron diffraction in TEM, are shown in Table 1. It has been shown [14, 15, 20] that the intragranular Ic phase maintains a parallel orientation relationship ([100]//[100], [010]//[010]) with the austenite matrix, is essentially cubic in shape, and has {100} type interfaces with the austenlte matrix. The TEM micrograph and diffraction pattern of early stage precipitation of i< phase in A6 specimen are shown in Figure 2. In group A alloys, which contain about 1 wt% carbon, the precipitation of intragranular g phase is observed only in alloys with higher aluminum contents, i.e. > 6.95 wt%. On the other hand, in group B alloys with lower aluminum content of -5.5 wt%, the intragranular i¢ phase does not precipitate in alloys irrespective of the carbon content. By comparing alloys A2 and A3, and A4 and A5, one may find that within the range studied, the content of Mn appears to be not an important factor in determining the precipitation of intragranular ~ phase. By using electrolytical extraction and Debye-Scherrer x-ray diffraction, the lattice constant of intragranular ~ phase extracted from alloy A6 has been determined to be 0.3747 nm, which is larger than the lattice constant of the austenltic Fe-Mn-A1-C alloys in Table 1. Since the intragranular !c phase and the matrix have a parallel orientation relationship and acommon {100} type interfaces [13-15, 20], the lattice misfit between the !¢ phase and the matrix will result in a considerable misfit strain energy. The misfit strain energy should be an important factor in determining the precipitation of intragranular ~ phase. Based upon the results of intra- and intergranular 1¢phase precipitation, it is clear that decreasing the aluminum and carbon content will decrease the chemical driving force for the precipitation of ~ phase. At the same time, this will also decrease the lattice constant of austenite which in turn will increase the misfit between austenite and intragranular 1¢ phase and then the coherent strain energy. Both the decrease in chemical driving force and the increase in strain energy, are unfavorable for the precipitation of intragranular I¢ phase. Therefore, it is possible that the lattice constant of austenite matrix may serve as a simple empirical index in determining the precipitation of intragranular ic phase.
502
PRECIPITATION
IN Fe-Mn-AI-C
Vol. 30, No.
To study this possibility, the measured lattice constants for these Fe-Mn-A1-C alloys are also listed in Table 1. Based on Table 1, a value of -0.3670 nm may be considered as a critical lattice constant for the austenite matrix, below which the intragranular i¢ phase will not develop. The lattice constant ao of austenitic Fe-Mn-A1-C alloys has been determined to be a linear function of chemical composition (in nm and wt%) [22], a o = 0.3574 + 0.000052 Mn + 0.00094 A1 + 0.0020 C - 0.00098 Si. By use of the critical lattice'constant (-0.3670 nm) and eqn. [1], a composition boundary for the precipitation of intragranular !¢ phase in Fe-Mn-AI-C alloys can be expressed as (in wt%) 0.098 AI + 0.208 C = 1 - 0.0054 Mn [2] According to eqn. [2], the composition boundaries for the precipitation of intragranular !c phase in austenitic Fe-Mn-A1-C alloys containing 25-34 wt% Mn are plotted in Figure 1 along with the experimental observations. Figure 1 also includes the results of various austenitic Fe-Mn-AI-C alloys reported in the literature, whose lattice constants were calculated by e q ~ [1]. The predicted boundary agrees quite well with the experimental results. It is to be noted that the change Of manganese content imposes only a small shift on the boundary line.
The precipitation of 1¢phase in austenitic Fe-Mn-A1-C alloys at 650°C was studied on selected alloys. Experimental results indicate that both aluminum and carbon contents are significant in determining the precipitation of these second phases, while the manganese content shows little effect within the range studied (23-33 wt%). An approximate composition boundary may be established for the precipitation of each phase and is summarized as follows. (1) To avoid the precipitation of intergranular • phase, the alloying addition should be limited to within about 5.5 wt% Al and 0.67 wt% C. (2) To avoid the precipitation of intragranular Ic phase, the alloying addition should be limited to within about 6.2 wt% AI and 1.0 wt% C. Furthermore, it is proposed that as an empirical rule, the lattice constant of austenite matrix may be used as a simple empirical index in determining the precipitation of intragranular ~ phase. A critical lattice constant of about 0.3670 nm has been found for the austenite matrix, above which the intragranular ~ phase does not precipitate. Based on this critical lattice constant, a composition boundary for the precipitation of intragranular K phase in Fe-Mn-AI-C alloys can be derived as (in wt%) 0.098 Al + 0.208 C = 10.0054 Mn. Acknowledgements
This work was financially supported by National Science Council, Taiwan, under contract no. NSC-79-0405-E110-10. Thanks are due to Prof. C. M. Wan for providing the materials used in this work. References
1. S.K. Banerji, Metal Prog., 113 (4), 59 (1978).
4
Vol. 30, No. 4
PRECIPITATION IN Fe-Mn-A1-C
2. R.K. You, PoWe Kao and D. Gala, Mater. Sci. Eng., Al17, 141 (1989). 3. J.L. Ham and R. E. Cairns Jr., Product Eng., 29 (12), 50 (1958). 4. D.J. Schmatz, Trans. ASM, 52, 898 (1960). 5. V.G. Rivlin, Int. Met. Rev., 28 (6), 309 (1983). 6. R. Wang and F. H. Beck, Metal Prog., 123 (4), 72 (1983). 7. J.G. Dub, S. H. Huarng and C. M. Wan, Chinese J. Mater. Sci., 16A (1), 14 (1984). 8. C.P. Chou and C. I-L Lee, scripta Metall., 23, 1109 (1989). 9. G.L. Kayak, Metal Sci. and Heat Treatment, 2, 95 (1969). 10. K. Sato, K. Tagawa and Y. Inoue, Metal1. Trans., 21A, 5 (1990). 11. Y. G. Kim, Y. S. Park and J. K. Han, Metall. Trans. A, 16A, 1689 (1985). 12. K.T. Luo, P. W. Kao and D. Gan, Mater. Sci. Eng., A151, L15 (1992). 13. S. M. Chu, P. W. Kao and D. Gan, Scripta Metall. et Mater., 26, 1067 (1992). 14. K. Sato, K. Tagawa and Y. Inoue, Scripta MetaU., 22, 899 (1988). 15. J. E. Krzanowski, Metall. Trans., 19A, 1873 (1988). 16. N. A. Storchak and A. G. Drachinskaya, Phys. Met. Metall., 44 (2), 123 (1978). 17. G. S. Krivonogov, M. F. Alekseyenko and G. S. Solov'yeva, Phys. Met. Metall., 39 (4), 775 (1975). 18. K. H. Han, W. K. Choo and D. E. Laughlin, Scripta Metall., 22, 1873 (1988). 19. K. H. Han and W. K. Choo, MetaU. Trans., 20A, 205 (1989). 20. K. H. Han, J. C. Yoon and W. K. Choo, Scripta Metall., 20, 33 (1986). 21. K. H. Han and W. K. Choo, MetaU. Trans., 14A, 973 (1983). 22. C. M. Chu, H. Huang, P. W. Kao and D. Gan, submitted to Scripta MetaU. et Mater.. 30Mn
34Mn 1.2
~
25Mn
• ~
1.0
o L O
0
0.8 0.6
N ,it 0.4 oO
0.2
!
!
4
6
8
wt%
Aluminum
0.0 0
2
' I !
10
12
FIG. 1. The precipitation of intragranular K phase as a function of alloy composition. The solid and open symbols indicate the compositions with and without precipitation, respectively. The square symbols indicate the results from the literature [10, 12, 15, 18]. The composition boundaries predicted by eqn. [2] are also indicated.
503
504
PRECIPITATION
IN Fe-Mn-AI-C
Vol.
30, No.
FIG. 2. Centered dark field TEM micrograph and diffraction pattern (zone axis [001]) of ahoy A6 aged at 650 °C for 1 h, showing the intragranular K phase.
4