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Effect of molybdenum substitution on B2 antiphase domain formation in rapidly solidified FeAlMo alloys
588
Materials Science and Engineering, A133 (1991) 588-591
Effect of molybdenum substitution on B2 antiphase domain formation in rapidly solidified Fe-AI-Mo alloys U. Prakash, R. A. Buckley and H. Jones School of Materials, University of Sheffield, Mappin Street, Sheffield (U.K.)
Abstract Studies at Sheffield have shown that the increase in B2 ordering temperature obtained with increase in aluminium content in rapidly solidified Fe-AI-X (X = Cr, Mo) alloys results in an increase in B2 antiphase domain (APD) size. The present contribution concerns the effect of increasing molybdenum substitution for iron in such systems, which is shown to increase the ordering temperature but decrease the B2 APD size. The results are discussed in terms of the extension of solid solubility of molybdenum by rapid solidification, molybdenum segregation to the antiphase boundaries, and molybdenum diffusion in these alloys.
1. Introduction In recent years, ordered intermetallics have been identified as potential high temperature structural materials [1, 2]. Application has been limited in the past by their low room temperature ductility and fracture toughness. Compositional and microstructural control achievable by methods such as judicious alloying, control of stoichiometry, rapid solidification processing (RSP) and thermomechanical processing may lead to considerable improvements in their mechanical properties [3-5]. The present work forms part of an effort [6] to study the effect of RSP on the structure and mechanical properties of ordered Fe-A1-X alloys (X= Cr, Mo, Mn). APD size has been proposed to be one of the factors controlling the ductility of ordered alloys [7, 8]. It has been shown by the authors that B2 antiphase domains form in a limited composition range in rapidly solidified (RS) Fe-A1-X (X = Cr, Mo) alloys, with the APD size increasing with increasing aluminium content of the alloy [9]. The present contribution concerns the effect on B2 APD formation of increasing molybdenum substitution for iron in RS Fe-AI alloys.
2. Experimental procedure Fe-AI-Mo alloys in the composition range (48-78) at.% Fe, (22-37) at.% AI and (0-15) at.% 0921-5093/91/$3.50
Mo were melted in a controlled atmosphere and were rapidly solidified by chill block melt-spinning using a mild steel wheel (wheel speed 27 m s -1, nozzle bore diameter 1 mm) to obtain ribbons 25-30 /zm in thickness. Details of the melting practice and melt-spinning procedure have been given elsewhere [6, 9, 10]. The meltspun ribbons were subjected to X-ray diffractometry in a Philips PW 1730 unit using Co-Ka radiation. Samples for thin foil electron microscopy were prepared in a Struers electropolishing unit using 30-50 V d.c. and 5% perchloric acid in ethanol as electrolyte at 263 K. The electropolished samples were examined in a Philips 301 transmission electron microscope using an accelerating voltage of 100 kV.
3. Results TEM and XRD of the melt-spun ribbons showed that all the ribbons were single phase. The grain size was found to be between 3/zm and 5 /~m and was not significantly dependent on composition. Binary Fe-A1 ribbons with aluminium content up to 22 at.% A1 had a distorted b.c.c, crystal structure [6], whereas those with higher aluminium content had a B2 structure. The ternary Fe-A1-Mo ribbons with aluminium content between 22 at.% and 32 at.% showed the presence of both B2 and DO 3 order in the TEM. © Elsevier Sequoia/Printed in The Netherlands
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No D O 3 order was detected by XRD probably because of excessive line broadening caused by the very fine ( < 0.01/~m) D O 3 domains shown by TEM to be present in these ribbons. Fe-A1-Mo ribbons with higher (37 at.%) aluminium content
showed only B2 order. The reasons for these differences in crystal structure are related to higher B2 and DO 3 ordering temperatures of F e - A l - M o ribbons and have been discussed elsewhere [6]. B2 antiphase domains (APDs) were observed in ribbons of the following compositions: (i) (26-32) at.% A1 for binary Fe-AI alloys. (ii) (22-37) at.% A1 for ternary Fe-AI-Mo alloys. The resulting domain sizes are given in Table 1, the ranges indicating the minimum and maximum sizes observed. The reasons for B2 APD formation in these composition ranges are related to a greater APB mobility at higher ordering temperatures achievable in ribbons with higher aluminium contents and are discussed elsewhere [6, 9]. The B2 APD size was observed to increase
(a)
(b)
TABLE 1 B2 antipbase domain size (in microns) as a function of alloy composition (atomic per cent) in F e - A I - M o melt-spun ribbons
0Mo 5Mo 10Mo 15Mo
22A1
26A1
32A1
37A1
b.c.c. b.c.c. <0.01" <0,01 a
0.2-0.4 0.1-0.5 a 0.1-0.2 a <0.1 a
0.5-1
5.0 b
--
5,0 b
0.1-0.5 a 0.1-0.4 a
0.4-0.7 0.3-0.5
~Also shows fine ( < 0.01 pro) DO 3 domains. bGrain size.
(c) Fig. 1. T E M 100 dark field micrographs showing decreasing B2 antiphase domain size with increasing molybdenum substitution for iron in melt-spun Fe-AI-(Mo) ribbons. (a) APBs in Fe-32A1 ribbon showing fringe contrast. (b) Isotropic APBs in Fe-32AI-15Mo ribbon. (c) APBs in Fe-32A1-15Mo ribbon showing some anisotropy.
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with increasing aluminium content and to decrease with increasing molybdenum substitution for iron at fixed aluminium content in these ribbons. The T E M micrographs in Fig. l(a)-(c) show decreasing APD size with increasing molybdenum substitution for iron in ribbons of fixed aluminium content (32 at.% Al). In general, the APD structure observed was typical of B2 alloys in forming an isotropic (i.e. smoothly curved) array of APBs, the so-called foam structure [11]. Large (15 at.%) molybdenum substitutions in Fe-32at.% Al seem to introduce some APB anisotropy (i.e. blockiness) (Fig. 1(c)). Figure 2 is a TEM micrograph showing anisotropic B2 APBs in Fe-37A1-15Mo ribbon. Trace analysis shows that these lie predominantly on {100} planes. 4. Discussion
The B2 APDs may form either by a nucleation and growth mechanism or by a continuous transformation [12, 13], both of which give rise to a ffoam structure" of APDs which then grow by coalescence. The B2 ordering temperature of Fe-A1 alloys increases with increasing aluminium content [14]. The domains in alloys with high aluminium contents are thus formed at higher temperatures, where mobility of the domain boundaries is expected to be higher, allowing the domains to grow to larger diameters during cooling. These arguments were used to explain the increase in APD size with increasing aluminium content in RS Fe-A1-X alloys [9]. Molybdenum substitution for excess iron in iron-rich Fe-AI alloys also increases the B2 ordering temperature [6, 15], so again the B2 APD size might have been expected to increase
with increasing molybdenum substitution. However, increasing molybdenum substitution leads to a decrease in the B2 domain size at all aluminium contents (Table 1). In Fe-37 at.%A1 ribbons, no APD network was observed because the alloy has a very high ordering temperature (1423 K) and the domain boundaries were probably mobile enough to coalesce and grow across the grains [6, 9]. Under these conditions, each grain represents one domain and so the APB network observed in alloys with lower aluminium contents would be absent. As may be seen from Table 1, APB networks are observed on making large (10 and 15 at.%) substitutions of molybdenum for iron in Fe-37at.%Al. These observations lead us to believe that molybdenum substitution reduces the mobility of APBs in these alloys leading to smaller APD sizes. The mechanisms for this reduced mobility are discussed below. The equilibrium solid solubility of molybdenum in Fe-AI alloys is limited (< 1 at.%) [6, 16]. The structure of these Fe-AI-Mo alloy ribbons after furnace cooling from 1273 K was established to consist of two phases [6, 15]. The primary phase was either D O 3 (for aluminium content between 22 at.% and 32 at.%) or B2 (for an aluminium content of 37 at.%), whereas the secondary phase was invariably molybdenumrich and was either Fe6Mo 7 (for aluminium content between 22 at.% and 32 at.%) or M o 3 m l (for an alurninium content of 37 at.%). Rapid solidification suppresses the formation of these molybdenum-rich phases, giving instead a single phase (DO3 or B2) supersaturated with molybdenum. During domain formation, this excess molybdenum would tend to segregate to domain boundaries to reduce their interfacial energy [17, 18]. It is proposed that this build-up of molybdenum at the APBs exerts a considerable drag on the domain boundaries. This reduced mobility of the APBs results in finer domain sizes in ribbons with high molybdenum content. This concentration of molybdenum atoms at APBs has been proposed to be the cause of formation of anisotropic APBs (Fig. 2) in Fe-37Al-15Mo ribbon
IlOl.
Fig. 2. TEM dark field micrograph showing anisotropic B2 antiphase domains in melt-spun Fe-37A1-15Mo ribbon.
Another possible mechanism for reduced APB mobility may be the substitution behaviour of molybdenum in ordered Fe-AI-Mo alloys. Excess iron in iron-rich B2 Fe-AI alloys substitutes at aluminium sublattice sites. Molybdenum substitution for this excess iron results in preferential molybdenum site occupancy for iron at
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aluminium sublattice sites [6, 15]. As a much heavier atom, molybdenum would be expected to diffuse more slowly than iron or aluminium within the lattice. The additional constraints because of this preferential site occupation would further reduce its rate of diffusion which, in turn, will lower the APB mobility. 5. Conclusions Increasing molybdenum substitution for iron in rapidly solidified Fe-A1-Mo alloys, in the composition range (78-48) at.% Fe, (0-15) at.% Mo, while resulting in an increase in B2 ordering temperature, decreases the B2 antiphase domain size. This reduced domain size is attributed to a reduced antiphase boundary mobility resulting from the expected segregation of molybdenum to domain boundaries during cooline. Acknowledgments One of us (UP) acknowledges financial support from the University of Sheffield and from the ORS award committee of the CVCP (Committee of Vice-Chancellors and Principals, UK).
References C. C. Koch, C. T. Liu and N. S. Stoloff (eds.), High Temperature Ordered lntermetallic Alloys, Mater. Res. Soc. Syrup. Proc., Vol. 39, 1985.
2 N. S. Stoloff, C. C. Koch, C. T. Liu and O. Izumi (eds.), High Temperature Ordered Intermetallic Alloys I1, Mater. Res. Soc. Symp. Proc., Vol. 81, 1987. 3 N.S. Stoloff, lnt. Met. Rev., 24(1984) 123. 4 C.T. Liu and J. O. Steigler, Science, 226 (1984) 636. 5 I. Baker and P. R. Munroe, J. Metals, 40 (2) (1988) 28. 6 U. Prakash, Ph.D. Thesis', University of Sheffield, 1989. 7 A. lnoue, T. Masumoto, H. Tomioka and N. Yano, Int. J. Rapid Solidification, 1 ( 1985 ) 115. 8 R. W. Cahn, in N. S. Stoloff, C. C. Koch, C. T. Liu and O. lsumi (eds.), High-Temp-,ature Ordered lntermetallic Alloys 11, Mater. Res. Soc. Syrup. Proc., VoL 81, 1987, p. 27. 9 U. Prakash. R. A. Buckley and H. Jones, Formation of B2 Antiphase Domains in Rapidly Solidified Fe-A1-X Alloys, in preparation for Phil. Mag. A. 10 U. Prakash, R. A. Buckley and H. Jones, Novel Faulted Structures in Rapidly Solidified Fe-37AI- 15Mo Alloy, in preparation for Acta Met Materiala. 11 A.T. English, Trans. Met. Soc. AIME, 236 (1966) 14. 12 R. A. Buckley and M. Rajkovic, in Phase l'ransJbrmations, Ser. 3, No. 11, Vol. 2, Institution of Metallurgists, London, 1979, p. 11-31. 13 L. E. Tanner and H. J. Leamy, in H. Warlimont (ed.), Order-disorder Transformations in Alloys, SpringerVerlag, Berlin, 1984, p. 180. 14 Anon, in Vol. 1, T. B. Massalski et al. (eds.), Bina o, Alloy Phase Diagrams, Amer. Soc. Met., 1986, p. 112. 15 U. Prakash, R. A. Buckley and H. Jones, Effect of Mo Substitution on Crystal Structure of Ordered Fe-AI-Mo alloys, in preparation. 16 R. H. Titran, K. M. Vedula and G. G. Anderson, in C. C. Koch, C. T. Liu and N. S. Stoloff (eds.), High-Temperature Ordered Intermetallic Alloys, Mater. Res. Soe. Syrup. Proc., 1985, p. 3(t9. 17 S. G. Cupschalk and N. Brown, Acre Metall., 15 (1967) 847. 18 S. G. Cupschalk and N. Brown, Aeta Metall., 16 (1968) 657.
Materials Science and Engineering, A133 (1991) 588-591
Effect of molybdenum substitution on B2 antiphase domain formation in rapidly solidified Fe-AI-Mo alloys U. Prakash, R. A. Buckley and H. Jones School of Materials, University of Sheffield, Mappin Street, Sheffield (U.K.)
Abstract Studies at Sheffield have shown that the increase in B2 ordering temperature obtained with increase in aluminium content in rapidly solidified Fe-AI-X (X = Cr, Mo) alloys results in an increase in B2 antiphase domain (APD) size. The present contribution concerns the effect of increasing molybdenum substitution for iron in such systems, which is shown to increase the ordering temperature but decrease the B2 APD size. The results are discussed in terms of the extension of solid solubility of molybdenum by rapid solidification, molybdenum segregation to the antiphase boundaries, and molybdenum diffusion in these alloys.
1. Introduction In recent years, ordered intermetallics have been identified as potential high temperature structural materials [1, 2]. Application has been limited in the past by their low room temperature ductility and fracture toughness. Compositional and microstructural control achievable by methods such as judicious alloying, control of stoichiometry, rapid solidification processing (RSP) and thermomechanical processing may lead to considerable improvements in their mechanical properties [3-5]. The present work forms part of an effort [6] to study the effect of RSP on the structure and mechanical properties of ordered Fe-A1-X alloys (X= Cr, Mo, Mn). APD size has been proposed to be one of the factors controlling the ductility of ordered alloys [7, 8]. It has been shown by the authors that B2 antiphase domains form in a limited composition range in rapidly solidified (RS) Fe-A1-X (X = Cr, Mo) alloys, with the APD size increasing with increasing aluminium content of the alloy [9]. The present contribution concerns the effect on B2 APD formation of increasing molybdenum substitution for iron in RS Fe-AI alloys.
2. Experimental procedure Fe-AI-Mo alloys in the composition range (48-78) at.% Fe, (22-37) at.% AI and (0-15) at.% 0921-5093/91/$3.50
Mo were melted in a controlled atmosphere and were rapidly solidified by chill block melt-spinning using a mild steel wheel (wheel speed 27 m s -1, nozzle bore diameter 1 mm) to obtain ribbons 25-30 /zm in thickness. Details of the melting practice and melt-spinning procedure have been given elsewhere [6, 9, 10]. The meltspun ribbons were subjected to X-ray diffractometry in a Philips PW 1730 unit using Co-Ka radiation. Samples for thin foil electron microscopy were prepared in a Struers electropolishing unit using 30-50 V d.c. and 5% perchloric acid in ethanol as electrolyte at 263 K. The electropolished samples were examined in a Philips 301 transmission electron microscope using an accelerating voltage of 100 kV.
3. Results TEM and XRD of the melt-spun ribbons showed that all the ribbons were single phase. The grain size was found to be between 3/zm and 5 /~m and was not significantly dependent on composition. Binary Fe-A1 ribbons with aluminium content up to 22 at.% A1 had a distorted b.c.c, crystal structure [6], whereas those with higher aluminium content had a B2 structure. The ternary Fe-A1-Mo ribbons with aluminium content between 22 at.% and 32 at.% showed the presence of both B2 and DO 3 order in the TEM. © Elsevier Sequoia/Printed in The Netherlands
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No D O 3 order was detected by XRD probably because of excessive line broadening caused by the very fine ( < 0.01/~m) D O 3 domains shown by TEM to be present in these ribbons. Fe-A1-Mo ribbons with higher (37 at.%) aluminium content
showed only B2 order. The reasons for these differences in crystal structure are related to higher B2 and DO 3 ordering temperatures of F e - A l - M o ribbons and have been discussed elsewhere [6]. B2 antiphase domains (APDs) were observed in ribbons of the following compositions: (i) (26-32) at.% A1 for binary Fe-AI alloys. (ii) (22-37) at.% A1 for ternary Fe-AI-Mo alloys. The resulting domain sizes are given in Table 1, the ranges indicating the minimum and maximum sizes observed. The reasons for B2 APD formation in these composition ranges are related to a greater APB mobility at higher ordering temperatures achievable in ribbons with higher aluminium contents and are discussed elsewhere [6, 9]. The B2 APD size was observed to increase
(a)
(b)
TABLE 1 B2 antipbase domain size (in microns) as a function of alloy composition (atomic per cent) in F e - A I - M o melt-spun ribbons
0Mo 5Mo 10Mo 15Mo
22A1
26A1
32A1
37A1
b.c.c. b.c.c. <0.01" <0,01 a
0.2-0.4 0.1-0.5 a 0.1-0.2 a <0.1 a
0.5-1
5.0 b
--
5,0 b
0.1-0.5 a 0.1-0.4 a
0.4-0.7 0.3-0.5
~Also shows fine ( < 0.01 pro) DO 3 domains. bGrain size.
(c) Fig. 1. T E M 100 dark field micrographs showing decreasing B2 antiphase domain size with increasing molybdenum substitution for iron in melt-spun Fe-AI-(Mo) ribbons. (a) APBs in Fe-32A1 ribbon showing fringe contrast. (b) Isotropic APBs in Fe-32AI-15Mo ribbon. (c) APBs in Fe-32A1-15Mo ribbon showing some anisotropy.
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with increasing aluminium content and to decrease with increasing molybdenum substitution for iron at fixed aluminium content in these ribbons. The T E M micrographs in Fig. l(a)-(c) show decreasing APD size with increasing molybdenum substitution for iron in ribbons of fixed aluminium content (32 at.% Al). In general, the APD structure observed was typical of B2 alloys in forming an isotropic (i.e. smoothly curved) array of APBs, the so-called foam structure [11]. Large (15 at.%) molybdenum substitutions in Fe-32at.% Al seem to introduce some APB anisotropy (i.e. blockiness) (Fig. 1(c)). Figure 2 is a TEM micrograph showing anisotropic B2 APBs in Fe-37A1-15Mo ribbon. Trace analysis shows that these lie predominantly on {100} planes. 4. Discussion
The B2 APDs may form either by a nucleation and growth mechanism or by a continuous transformation [12, 13], both of which give rise to a ffoam structure" of APDs which then grow by coalescence. The B2 ordering temperature of Fe-A1 alloys increases with increasing aluminium content [14]. The domains in alloys with high aluminium contents are thus formed at higher temperatures, where mobility of the domain boundaries is expected to be higher, allowing the domains to grow to larger diameters during cooling. These arguments were used to explain the increase in APD size with increasing aluminium content in RS Fe-A1-X alloys [9]. Molybdenum substitution for excess iron in iron-rich Fe-AI alloys also increases the B2 ordering temperature [6, 15], so again the B2 APD size might have been expected to increase
with increasing molybdenum substitution. However, increasing molybdenum substitution leads to a decrease in the B2 domain size at all aluminium contents (Table 1). In Fe-37 at.%A1 ribbons, no APD network was observed because the alloy has a very high ordering temperature (1423 K) and the domain boundaries were probably mobile enough to coalesce and grow across the grains [6, 9]. Under these conditions, each grain represents one domain and so the APB network observed in alloys with lower aluminium contents would be absent. As may be seen from Table 1, APB networks are observed on making large (10 and 15 at.%) substitutions of molybdenum for iron in Fe-37at.%Al. These observations lead us to believe that molybdenum substitution reduces the mobility of APBs in these alloys leading to smaller APD sizes. The mechanisms for this reduced mobility are discussed below. The equilibrium solid solubility of molybdenum in Fe-AI alloys is limited (< 1 at.%) [6, 16]. The structure of these Fe-AI-Mo alloy ribbons after furnace cooling from 1273 K was established to consist of two phases [6, 15]. The primary phase was either D O 3 (for aluminium content between 22 at.% and 32 at.%) or B2 (for an aluminium content of 37 at.%), whereas the secondary phase was invariably molybdenumrich and was either Fe6Mo 7 (for aluminium content between 22 at.% and 32 at.%) or M o 3 m l (for an alurninium content of 37 at.%). Rapid solidification suppresses the formation of these molybdenum-rich phases, giving instead a single phase (DO3 or B2) supersaturated with molybdenum. During domain formation, this excess molybdenum would tend to segregate to domain boundaries to reduce their interfacial energy [17, 18]. It is proposed that this build-up of molybdenum at the APBs exerts a considerable drag on the domain boundaries. This reduced mobility of the APBs results in finer domain sizes in ribbons with high molybdenum content. This concentration of molybdenum atoms at APBs has been proposed to be the cause of formation of anisotropic APBs (Fig. 2) in Fe-37Al-15Mo ribbon
IlOl.
Fig. 2. TEM dark field micrograph showing anisotropic B2 antiphase domains in melt-spun Fe-37A1-15Mo ribbon.
Another possible mechanism for reduced APB mobility may be the substitution behaviour of molybdenum in ordered Fe-AI-Mo alloys. Excess iron in iron-rich B2 Fe-AI alloys substitutes at aluminium sublattice sites. Molybdenum substitution for this excess iron results in preferential molybdenum site occupancy for iron at
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aluminium sublattice sites [6, 15]. As a much heavier atom, molybdenum would be expected to diffuse more slowly than iron or aluminium within the lattice. The additional constraints because of this preferential site occupation would further reduce its rate of diffusion which, in turn, will lower the APB mobility. 5. Conclusions Increasing molybdenum substitution for iron in rapidly solidified Fe-A1-Mo alloys, in the composition range (78-48) at.% Fe, (0-15) at.% Mo, while resulting in an increase in B2 ordering temperature, decreases the B2 antiphase domain size. This reduced domain size is attributed to a reduced antiphase boundary mobility resulting from the expected segregation of molybdenum to domain boundaries during cooline. Acknowledgments One of us (UP) acknowledges financial support from the University of Sheffield and from the ORS award committee of the CVCP (Committee of Vice-Chancellors and Principals, UK).
References C. C. Koch, C. T. Liu and N. S. Stoloff (eds.), High Temperature Ordered lntermetallic Alloys, Mater. Res. Soc. Syrup. Proc., Vol. 39, 1985.
2 N. S. Stoloff, C. C. Koch, C. T. Liu and O. Izumi (eds.), High Temperature Ordered Intermetallic Alloys I1, Mater. Res. Soc. Symp. Proc., Vol. 81, 1987. 3 N.S. Stoloff, lnt. Met. Rev., 24(1984) 123. 4 C.T. Liu and J. O. Steigler, Science, 226 (1984) 636. 5 I. Baker and P. R. Munroe, J. Metals, 40 (2) (1988) 28. 6 U. Prakash, Ph.D. Thesis', University of Sheffield, 1989. 7 A. lnoue, T. Masumoto, H. Tomioka and N. Yano, Int. J. Rapid Solidification, 1 ( 1985 ) 115. 8 R. W. Cahn, in N. S. Stoloff, C. C. Koch, C. T. Liu and O. lsumi (eds.), High-Temp-,ature Ordered lntermetallic Alloys 11, Mater. Res. Soc. Syrup. Proc., VoL 81, 1987, p. 27. 9 U. Prakash. R. A. Buckley and H. Jones, Formation of B2 Antiphase Domains in Rapidly Solidified Fe-A1-X Alloys, in preparation for Phil. Mag. A. 10 U. Prakash, R. A. Buckley and H. Jones, Novel Faulted Structures in Rapidly Solidified Fe-37AI- 15Mo Alloy, in preparation for Acta Met Materiala. 11 A.T. English, Trans. Met. Soc. AIME, 236 (1966) 14. 12 R. A. Buckley and M. Rajkovic, in Phase l'ransJbrmations, Ser. 3, No. 11, Vol. 2, Institution of Metallurgists, London, 1979, p. 11-31. 13 L. E. Tanner and H. J. Leamy, in H. Warlimont (ed.), Order-disorder Transformations in Alloys, SpringerVerlag, Berlin, 1984, p. 180. 14 Anon, in Vol. 1, T. B. Massalski et al. (eds.), Bina o, Alloy Phase Diagrams, Amer. Soc. Met., 1986, p. 112. 15 U. Prakash, R. A. Buckley and H. Jones, Effect of Mo Substitution on Crystal Structure of Ordered Fe-AI-Mo alloys, in preparation. 16 R. H. Titran, K. M. Vedula and G. G. Anderson, in C. C. Koch, C. T. Liu and N. S. Stoloff (eds.), High-Temperature Ordered Intermetallic Alloys, Mater. Res. Soe. Syrup. Proc., 1985, p. 3(t9. 17 S. G. Cupschalk and N. Brown, Acre Metall., 15 (1967) 847. 18 S. G. Cupschalk and N. Brown, Aeta Metall., 16 (1968) 657.