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On the strength of nickel-containing B2 iron aluminides
Materials Science and Engineering A239 – 240 (1997) 245 – 250
On the strength of nickel-containing B2 iron aluminides J.H. Schneibel a,*, E.D. Specht a, P.R. Munroe b a
Metals and Ceramics Di6ision, Oak Ridge National Laboratory, P.O. Box 2008, Oak Ridge, TN 37831 -6115, USA b School of Materials Science and Engineering, Uni6ersity of New South Wales, Sydney, NSW 2052, Australia
Abstract When some of the iron in the B2 iron aluminide Fe-45 at.% Al is replaced by nickel, several important changes occur. Nickel dramatically slows down the removal of quenched-in thermal vacancies during annealing. Also, nickel additions render the hardening due to thermal vacancies less effective than that in binary iron aluminides. Both of these effects may be linked to the formation of nickel-vacancy complexes. Nickel additions also cause substantial solid solution strengthening, and possible reasons for this result are discussed. Our findings are supported by vacancy concentration determinations, by yield strength measurements, and by transmission electron microscope observations. © 1997 Elsevier Science S.A. Keywords: B2 iron aluminide; Nickel; Annealing
1. Introduction Iron aluminides with the B2 structure contain high concentrations of thermal vacancies at elevated temperatures. The origin of these thermal vacancies and their effect on mechanical properties have been discussed in recent publications by Fu et al. [1], Chang et al. [2], Wu¨rschum et al. [3], Munroe [4], Schneibel et al. [5], and Pike et al. (unpublished). Essentially, vacancies in iron aluminides exhibit low enthalpies of formation (e.g. 1 eV), and high enthalpies of migration (e.g. 1.7 eV) [3]. As the Al content of iron aluminides increases towards that of the stoichiometric compound FeAl, the enthalpy of formation decreases (i.e. at a given temperature, the equilibrium vacancy concentration increases) and the enthalpy of migration increases (i.e. it takes longer to reach the vacancy equilibrium concentration corresponding to that temperature) [6]. Since the vacancies cause hardening [2,7], it is difficult to assess the intrinsic mechanical properties of iron aluminides (i.e. the mechanical properties in the absence of vacancies), in particular when the Al concentrations approach 50 at.%. It has therefore become common practice to anneal iron aluminides for prolonged periods of time at relatively
* Corresponding author. Tel.: +1 423 5764644; fax: + 1 423 5747659; e-mail: [email protected] 0921-5093/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved. PII S 0 9 2 1 - 5 0 9 3 ( 9 7 ) 0 0 5 8 8 - 1
low temperatures to reach low vacancy equilibrium concentrations. In this way mechanical data with no or negligible interference from vacancies may be obtained. Nagpal and Baker [7], for example, annealed for 120 h at 673 K, whereas Yoshimi et al. [8] annealed for 150 h at a slightly higher temperature, namely, 713 K. The situation becomes considerably more complicated when ternary solid solution alloying additions are made. In particular in the case of Ni additions, pronounced strengthening is found as evidenced by room temperature hardness and yield strength measurements [4,5,9–13]. The Ni additions influence not only the equilibrium vacancy concentrations, but also the rates with which equilibrium is reached [5]. Some care is therefore required in order to unambiguously determine whether strengthening in Ni containing iron aluminides is primarily due to Ni, or whether it is an indirect effect due to retained vacancies. In the present paper, we extend previous work [5] by studying iron aluminides containing 45 at.% Al and up to 10 at.% Ni. In particular, we measure the vacancy removal kinetics as well as the vacancy concentrations. In this way we will be able to demonstrate more clearly that Ni does indeed cause significant solid solution strengthening in iron aluminides. At the same time, our results are consistent with the previously suggested interaction between Ni and vacancies resulting in Ni-vacancy complexes [9].
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J.H. Schneibel et al. / Materials Science and Engineering A239–240 (1997) 245–250
2. Experimental procedure Iron aluminides containing 45 at.% Al and up to 10 at.% Ni were prepared by arc-melting elemental, 99.99 wt% pure materials into buttons, remelting those buttons several times, and drop-casting into Cu molds. The melting was carried out in argon slightly below atmospheric pressure (i.e. 70 kPa). The ingots were homogenized for 1 h at 1273 K. Parts of the ingots were hot isostatically pressed (HIPed) for 1.8 ks at 1373 K and 200 MPa in an attempt to minimize residual porosity. Unfortunately, sectioning of the HIPed materials revealed some remaining residual porosity, indicating that full densification was not achieved. The reason for this may have been hairline cracks in the castings. The cast as well as the cast and HIPed ingots were then annealed for 3 days at 973 K. According to measurements of density versus time, 3 days was sufficient to reach thermal vacancy equilibrium at 973 K. For the purpose of lattice parameter measurements, ingots quenched from 1273 K were crushed into − 325 mesh ( B 45 mm) powder and annealed in evacuated quartz tubes for 3 days at 973 K, followed by quenching of the quartz tubes into water. Some of the crushed powder was annealed in a dynamic vacuum (10 − 4 Pa) for 3 days at 973 K, followed by 6 days at 773 K. Density measurements on bulk specimens showed that the latter anneal was sufficiently long to reach vacancy equilibrium. The lattice parameters were determined via x-ray powder diffraction using an internal Si standard. Bulk density measurements were carried out by weighing specimens with typical dimensions of 10×20 × 20 mm in distilled water and air, respectively. In order to avoid specimen to specimen variations, one and the same specimen was used for all the density measurements corresponding to a particular composition. Following various heat treatments in air, the specimens were usually quenched in oil. Prior to annealing at a particular temperature, the specimens were always given a 3 day/973 K equilibration anneal. The reproducibility of the densities of the HIPed specimens was 1 part in 104. The density of a Si standard with a size similar to that of the iron aluminide specimens could be reproduced with a standard deviation of 5 parts in 105. Atomic vacancy concentrations were evaluated from the equation cv =(rx − rb)/ rb, where rx is the x-ray density determined from the lattice parameter, the crystal structure, and the nominal composition, and rb the bulk density. Room temperature compression tests were carried out with HIPed specimens having typical dimensions of 6×6 × 10 mm, at a crosshead speed of 10 mm s − 1. The grain sizes were well above 200 mm, making strong Hall-Petch hardening unlikely. Specimens for transmission electron microscope (TEM) examination were annealed at 1223 K for 2 h, air-cooled, and subsequently annealed at 673 K for times of up to 120 h. TEM specimens were prepared
by polishing in a solution of 30% nitric acid in methanol at a temperature of 250 K with a voltage of 10 V and examined in a JEOL 2000FX.
3. Results and discussion
3.1. Kinetics of 6acancy remo6al and microstructural e6olution Cast and homogenized specimens were annealed for 3 days at 973 K, and then annealed for various times at 723 K. Each specimen was occasionally taken out of the furnace in order to determine its bulk density, and then reinserted. The density-time curves were fitted with an exponential function of the form r= r0 +Dr ×(1exp(− t/tr)) in order to evaluate the characteristic time tr for vacancy removal. It is immediately seen from Fig. 1 that tr for the Ni containing specimen, Fe–45Al–3Ni (compositions are stated in at.% unless otherwise noted), is significantly larger than tr for the binary iron aluminide, Fe–45Al. Fig. 2 shows that tr increases linearly with the Ni content. An addition of 10 at.% Ni increases tr by a factor of 27 as compared to the binary alloy. This means that, as compared to binary Fe– 45Al, the Ni-containing specimens require either longer annealing times, or higher annealing temperatures, in order to reach their equilibrium vacancy concentration. This result is in qualitative agreement with our earlier work [5]. In Fig. 3 the (effective) migration enthalpy Hm is evaluated for several alloys by assuming a relationship of the form t8 exp(Hm/(kT)). The value of 1.46 eV found for Fe–45Al is similar to migration enthalpies found with positron annihilation techniques: Wu¨rschum [3] determined a migration enthalpy of 1.7 eV for Fe–39 at.% Al, whereas Wolff et al. [14] (these proceedings) obtained 1.2 eV for Fe–40 at.% Al. Con-
Fig. 1. Bulk density of cast specimens vs. time at 723 K. Prior to the 723 K anneal, the specimens were homogenized at 1273 K and annealed for 3 days at 973 K. Note that these specimens were not HIPed.
J.H. Schneibel et al. / Materials Science and Engineering A239–240 (1997) 245–250
Fig. 2. Characteristic time tr for vacancy removal at 773 K, as a function of Ni content.
sistent with the slowing down of the vacancy removal rates by Ni additions, the activation enthalpies appear to increase with increasing Ni content. However, this finding is somewhat tentative in view of the scatter and the limited number of data points. TEM studies of the alloys following heat treatment at 1223 K and air-cooling showed that they contained a very low density of 001 dislocations. However, a number of microstructural features developed following prolonged annealing at 673 K. Typical transmission electron micrographs are shown in Figs. 4 and 5. Alloys containing less than 10 at.% Ni contained a number of cuboidal voids 50 – 100 nm in diameter, together with 001 edge dislocations. As the Ni content for these alloys increased, the density of the voids decreased and the dislocation density increased. No voids were observed in the alloy with 10 at.% Ni, only 001 dislocations. Both the voids and the 001 dislocations are associated with the removal of thermal vacancies from the lattice.
Fig. 3. Determination of activation enthalpies for migration.
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Fig. 4. TEM micrograph of Fe – 45Al – 3Ni (arc-cast + 2 h/1223 K+ 120 h/673 K) illustrating low density of 001 dislocations and the presence of voids.
3.2. Room temperature yield strength measurements Fig. 6 shows the room temperature yield stresses for several Fe–Al–Ni alloys as a function of the temperature from which the specimens were quenched into oil. In all cases, the anneals prior to the quench were sufficiently long to achieve vacancy equilibrium, as verified by bulk density measurements. Typical times were 3 days at 973 K, 6 days at 873 K, and 12 days at 773 K. The yield stresses are seen to increase with increasing annealing temperature. These yield stress increases are due to quenched-in vacancies [5]. A dominant feature of Fig. 6 is the low slope of the Ni-containing specimens, as compared to binary Fe–45Al. Therefore, the curves for Fe–45Al and Fe– 45Al–6Ni intersect: Fe–45Al–6Ni quenched from 973 K exhibits a strength marginally lower than that of Fe–45Al quenched from the same temperature. If the vacancy contributions were ignored, this result would erro-
Fig. 5. TEM micrograph of Fe – 45Al – 10Ni (arc-cast + 2 h/1223 K +120 h/673 K) illustrating high density of 001 dislocations and absence of voids.
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Fig. 6. Room temperature yield stress as a function of the temperature of prior anneals, for iron aluminides with different Ni contents.
neously suggest solid solution softening due to nickel. If quenching is performed from lower temperatures, such as 773 K, the Ni-containing alloys are substantially stronger than Fe – 45Al, suggesting solid solution strengthening due to Ni.
3.3. Vacancy concentrations and solid solution strengthening The densities of HIPed specimens are shown in Fig. 7 as a function of annealing temperature prior to quenching. Prior to the 773 and 873 K measurements, the specimens were annealed for 3 days at 973 K. Fig. 7 shows that the density decreases with increasing annealing temperature, which is consistent with increasing vacancy concentrations. Fig. 8 shows the lattice parameters measured after annealing at 973 and 773 K, respectively. While sufficiently high vacancy concentrations collapse the lattice slightly and reduce therefore the lattice parameter [2],
Fig. 7. Bulk density of iron aluminides with different Ni contents as a function of annealing temperature. The anneals were sufficiently long to reach equilibrium.
Fig. 8. Lattice parameter of iron aluminides with different Ni contents after reaching equilibrium at 773 and 973 K.
the vacancy concentration difference between 973 and 773 K is too small to cause a noticeable effect on the lattice parameter. Therefore, there is virtually no difference between the two sets of data. From the bulk densities and the lattice parameters vacancy concentrations were calculated. The absolute values of these vacancy concentrations are subject to considerable uncertainty. One reason for this is the residual porosity in the specimens used for the density measurements. Also, small variations in the local specimen composition may cause errors. For example, a difference between the nominal and the actual Al concentration of only 0.1% would change the vacancy concentration by 0.07%. The 973 K vacancy concentrations found here are usually higher, by up to 0.2%, than Pike et al.’s (unpublished, [15]) concentrations for similar alloys. Pike et al. state that their concentration measurements were accurate to within 0.1%. The absolute values of our vacancy concentrations are likely to have a similar uncertainty. However, since in our experiments one and the same specimen was used for all the density measurements corresponding to a particular composition, the vacancy concentration differences, for different annealing conditions of a particular composition, are very reliable. Since differences in the bulk densities could be determined within 1 part in 104, the error in the (relative) vacancy concentrations for one and the same specimen is of similar magnitude. This means that the differences in the slopes of the curves in Fig. 9 showing the yield stresses as a function of the vacancy concentrations are reliable. Consistent with Fig. 6, the Ni-containing alloys exhibit a significantly lower slope than binary Fe–45Al. For low vacancy concentrations such as 0.15 at.%, Fig. 9 suggests strong solid solution hardening in Fe–45Al– 10Ni. For higher vacancy concentrations such as 0.25 at.%, Fe–45Al–6Ni is in fact weaker than Fe–45Al, even if the uncertainty in the absolute vacancy concentration is taken into account. This observation cannot
J.H. Schneibel et al. / Materials Science and Engineering A239–240 (1997) 245–250
be due to Ni-induced solid solution softening, since the Ni-containing iron aluminides with 0.15 at.% vacancies are much stronger than the binary Fe – 45Al. We conclude, therefore, that for a given vacancy concentration, the vacancy hardening in the Ni-containing specimens is not as strong as that in Fe – 45Al. On first sight, the formation of voids might be a reason for this finding. While voids are not seen in binary iron aluminides, they have been found in Fe – Al – Ni [9]. Voids, while reducing the bulk density, are not effective room temperature hardeners. In particular, they remove individual vacancies from the crystal lattice, thus making them unavailable for hardening. While this argument could be valid for Fe–45Al–3Ni, which does contain voids (Fig. 4), it is not valid for Fe – 45Al – 10Ni, which contains no voids, but only dislocations (Fig. 5). Another possibility to explain our result is the formation of Ni-vacancy clusters as they have been suggested by Munroe [9]. Although such clusters may have a higher obstacle strength for dislocations than individual vacancies, the number density of these clusters could be, depending on their size, substantially lower than that of individual vacancies and the net strengthening might be reduced. In Fig. 10 the yield stress is shown as a function of the Ni content, after a 12 day anneal at 773 K. According to Fig. 9, the residual vacancy concentrations are between 0.1 and 0.3 at.%. In view of the uncertainty in the absolute vacancy concentrations, and the uncertainty about the effectiveness of vacancies in Ni-containing FeAl, it is not possible at the present time to correct Fig. 10 for the effect of vacancies. Since vacancy hardening has probably not been eliminated completely, Fig. 10 is likely to represent an overestimate of the solid solution strengthening due to Ni. The present data can also not rule out the influence of Hall–Petch effects due to grain size variations [16]. However, since the vacancies in Ni-containing iron aluminides have been seen to be less effective strengtheners than the
Fig. 9. Room temperature yield strength vs. vacancy concentration, for iron aluminides with different Ni contents.
249
Fig. 10. Room temperature yield strength of iron aluminides with different Ni contents, after a 12 day anneal at 773 K.
vacancies in binary iron aluminides, Ni is likely to be an effective solid solution strengthener in iron aluminides. This would be difficult to justify on the basis of the atomic size difference alone. Nickel has an atomic radius similar to that of Fe, and it is known to occupy the iron sublattice [17]. Elastic modulus measurements also did not show any substantial effect for Ni additions to FeAl [13]. There may be several possibilities, such as electronic or magnetic effects, which may change the dislocation–solute atom interaction and cause the strengthening. One other possibility that has been suggested [9,18] is the following: additions of Ni tend to change the slip behavior of iron aluminides from 111 slip to 100 slip [19]. This transition is illustrated in Fig. 11, in which the Burgers vectors and critical resolved shear stresses (CRSS) experimentally observed for FeAl and NiAl, respectively, are indicated. As Ni is added to FeAl, 111 slip is likely to become more difficult, since the CRSS for 111 slip in NiAl is much higher than that of FeAl. This trend is schematically shown by the solid line in Fig. 11. Also, as Ni is added, the increased tendency for 100 slip may in-
Fig. 11. Schematic illustration of the effect of Ni on the slip systems and strengths of iron aluminides.
J.H. Schneibel et al. / Materials Science and Engineering A239–240 (1997) 245–250
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duce changes in the core structure of the 111 dislocations and reduce their mobility. Thus, the solid solution strengthening by Ni may very well be the result of a dislocation core effect.
4. Conclusions Additions of Ni to Fe – 45Al slow down dramatically the kinetics of vacancy migration and removal. This is likely to be associated with binding between Ni and thermal vacancies, and the formation of Ni-vacancy complexes. While Ni gives rise to high solid solution strengthening, it renders vacancy hardening less effective as verified by yield strength measurements for different nickel and vacancy concentrations. Again, Nivacancy complexes provide a possible explanation of this finding, since such complexes would reduce the number density of defects interacting with the dislocations.
Acknowledgements This research was sponsored by the Division of Materials Sciences, US Department of Energy under contract number DE-AC05-96OR22464 with Lockheed Martin Energy Research Corporation. Valuable discussions with C.L. Fu and Xing-Dong Wang are appreciated, as well as the review of this manuscript by T. Takasugi and L.M. Pike.
.
References [1] C.L. Fu, Y.-Y. Ye, M.H. Yoo, K.M. Ho, Phys. Rev. B 48 (1993) 6712. [2] Y.A. Chang, L.M. Pike, C.T. Liu, A.R. Bilbrey, D.S. Stone, Intermetallics 1 (1993) 107. [3] R. Wu¨rschum, C. Grupp, H.-E. Schaefer, Phys. Rev. Lett. 75 (1995) 97. [4] P.R. Munroe, Processing, properties, and applications, in: S.C. Deevi et al. (Eds.), Int. Symp. Nickel Iron Aluminides, ASM, 1997, 329 pp. [5] J.H. Schneibel, P.R. Munroe, L.M. Pike, in: C.C. Koch et al. (Eds), High-temperature ordered intermetallic alloys VII, Mater. Res. Soc. Proc., 460, 1997, pp. 379 – 384. [6] R. Wu¨rschum, K. Badura-Gergen, E.A. Ku¨mmerle, C. Grupp, H.-E. Schaefer, Phys. Rev. B 54 (1996) 849. [7] P. Nagpal, I. Baker, Metall. Trans. A 21A (1990) 2281. [8] K. Yoshimi, S. Hanada, H. Tokuno, Mater. Trans. JIM 35 (1994) 51. [9] P.R. Munroe, Intermetallics 4 (1996) 5. [10] C.H. Kong, P.R. Munroe, Scr. Metall. Mater. 30 (1994) 1079. [11] P.R. Munroe, C.H. Kong, Intermetallics 4 (1996) 403. [12] J.H. Schneibel, E.P. George, E.D. Specht, J.A. Horton, in J.A. Horton et al. (Eds), High temperature ordered intermetallic alloys, vol. VI, Mater. Res. Soc. Proc., 364, 1995, 73 pp. [13] J.H. Schneibel, E.D. Specht, W.A. Simpson, Intermetallics 4 (1996) 581. [14] J. Wolff, M. Franz, A. Broska, B. Ko¨hler, Th. Hehenkamp, Mater. Sci. Eng. A239/240 (1997) 213 – 219. [15] L.M. Pike, Y.A. Chang, C.T. Liu, Acta Mater. 45 (1997) 3709. [16] J. Winterhoff, E. Nembach, Scr. Mater. 35 (1996) 999. [17] I.M. Anderson, Acta Mater. 45 (1997) 3897. [18] J.H. Schneibel, C.T. Liu, in: G.M. Stocks, P.E.A. Turchi (Eds), Alloy Phase Stability and Design, TMS, Pittsburgh, 1994, 235 pp. [19] D.K. Patrick, K.-M. Chang, D.B. Miracle, H.A. Lipsitt, in H.A. Johnson et al. (Eds.), High-temperature ordered intermetallic alloys, vol. IV, Mater. Res. Soc. Proc., 213, 1991, 267 pp.
On the strength of nickel-containing B2 iron aluminides J.H. Schneibel a,*, E.D. Specht a, P.R. Munroe b a
Metals and Ceramics Di6ision, Oak Ridge National Laboratory, P.O. Box 2008, Oak Ridge, TN 37831 -6115, USA b School of Materials Science and Engineering, Uni6ersity of New South Wales, Sydney, NSW 2052, Australia
Abstract When some of the iron in the B2 iron aluminide Fe-45 at.% Al is replaced by nickel, several important changes occur. Nickel dramatically slows down the removal of quenched-in thermal vacancies during annealing. Also, nickel additions render the hardening due to thermal vacancies less effective than that in binary iron aluminides. Both of these effects may be linked to the formation of nickel-vacancy complexes. Nickel additions also cause substantial solid solution strengthening, and possible reasons for this result are discussed. Our findings are supported by vacancy concentration determinations, by yield strength measurements, and by transmission electron microscope observations. © 1997 Elsevier Science S.A. Keywords: B2 iron aluminide; Nickel; Annealing
1. Introduction Iron aluminides with the B2 structure contain high concentrations of thermal vacancies at elevated temperatures. The origin of these thermal vacancies and their effect on mechanical properties have been discussed in recent publications by Fu et al. [1], Chang et al. [2], Wu¨rschum et al. [3], Munroe [4], Schneibel et al. [5], and Pike et al. (unpublished). Essentially, vacancies in iron aluminides exhibit low enthalpies of formation (e.g. 1 eV), and high enthalpies of migration (e.g. 1.7 eV) [3]. As the Al content of iron aluminides increases towards that of the stoichiometric compound FeAl, the enthalpy of formation decreases (i.e. at a given temperature, the equilibrium vacancy concentration increases) and the enthalpy of migration increases (i.e. it takes longer to reach the vacancy equilibrium concentration corresponding to that temperature) [6]. Since the vacancies cause hardening [2,7], it is difficult to assess the intrinsic mechanical properties of iron aluminides (i.e. the mechanical properties in the absence of vacancies), in particular when the Al concentrations approach 50 at.%. It has therefore become common practice to anneal iron aluminides for prolonged periods of time at relatively
* Corresponding author. Tel.: +1 423 5764644; fax: + 1 423 5747659; e-mail: [email protected] 0921-5093/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved. PII S 0 9 2 1 - 5 0 9 3 ( 9 7 ) 0 0 5 8 8 - 1
low temperatures to reach low vacancy equilibrium concentrations. In this way mechanical data with no or negligible interference from vacancies may be obtained. Nagpal and Baker [7], for example, annealed for 120 h at 673 K, whereas Yoshimi et al. [8] annealed for 150 h at a slightly higher temperature, namely, 713 K. The situation becomes considerably more complicated when ternary solid solution alloying additions are made. In particular in the case of Ni additions, pronounced strengthening is found as evidenced by room temperature hardness and yield strength measurements [4,5,9–13]. The Ni additions influence not only the equilibrium vacancy concentrations, but also the rates with which equilibrium is reached [5]. Some care is therefore required in order to unambiguously determine whether strengthening in Ni containing iron aluminides is primarily due to Ni, or whether it is an indirect effect due to retained vacancies. In the present paper, we extend previous work [5] by studying iron aluminides containing 45 at.% Al and up to 10 at.% Ni. In particular, we measure the vacancy removal kinetics as well as the vacancy concentrations. In this way we will be able to demonstrate more clearly that Ni does indeed cause significant solid solution strengthening in iron aluminides. At the same time, our results are consistent with the previously suggested interaction between Ni and vacancies resulting in Ni-vacancy complexes [9].
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2. Experimental procedure Iron aluminides containing 45 at.% Al and up to 10 at.% Ni were prepared by arc-melting elemental, 99.99 wt% pure materials into buttons, remelting those buttons several times, and drop-casting into Cu molds. The melting was carried out in argon slightly below atmospheric pressure (i.e. 70 kPa). The ingots were homogenized for 1 h at 1273 K. Parts of the ingots were hot isostatically pressed (HIPed) for 1.8 ks at 1373 K and 200 MPa in an attempt to minimize residual porosity. Unfortunately, sectioning of the HIPed materials revealed some remaining residual porosity, indicating that full densification was not achieved. The reason for this may have been hairline cracks in the castings. The cast as well as the cast and HIPed ingots were then annealed for 3 days at 973 K. According to measurements of density versus time, 3 days was sufficient to reach thermal vacancy equilibrium at 973 K. For the purpose of lattice parameter measurements, ingots quenched from 1273 K were crushed into − 325 mesh ( B 45 mm) powder and annealed in evacuated quartz tubes for 3 days at 973 K, followed by quenching of the quartz tubes into water. Some of the crushed powder was annealed in a dynamic vacuum (10 − 4 Pa) for 3 days at 973 K, followed by 6 days at 773 K. Density measurements on bulk specimens showed that the latter anneal was sufficiently long to reach vacancy equilibrium. The lattice parameters were determined via x-ray powder diffraction using an internal Si standard. Bulk density measurements were carried out by weighing specimens with typical dimensions of 10×20 × 20 mm in distilled water and air, respectively. In order to avoid specimen to specimen variations, one and the same specimen was used for all the density measurements corresponding to a particular composition. Following various heat treatments in air, the specimens were usually quenched in oil. Prior to annealing at a particular temperature, the specimens were always given a 3 day/973 K equilibration anneal. The reproducibility of the densities of the HIPed specimens was 1 part in 104. The density of a Si standard with a size similar to that of the iron aluminide specimens could be reproduced with a standard deviation of 5 parts in 105. Atomic vacancy concentrations were evaluated from the equation cv =(rx − rb)/ rb, where rx is the x-ray density determined from the lattice parameter, the crystal structure, and the nominal composition, and rb the bulk density. Room temperature compression tests were carried out with HIPed specimens having typical dimensions of 6×6 × 10 mm, at a crosshead speed of 10 mm s − 1. The grain sizes were well above 200 mm, making strong Hall-Petch hardening unlikely. Specimens for transmission electron microscope (TEM) examination were annealed at 1223 K for 2 h, air-cooled, and subsequently annealed at 673 K for times of up to 120 h. TEM specimens were prepared
by polishing in a solution of 30% nitric acid in methanol at a temperature of 250 K with a voltage of 10 V and examined in a JEOL 2000FX.
3. Results and discussion
3.1. Kinetics of 6acancy remo6al and microstructural e6olution Cast and homogenized specimens were annealed for 3 days at 973 K, and then annealed for various times at 723 K. Each specimen was occasionally taken out of the furnace in order to determine its bulk density, and then reinserted. The density-time curves were fitted with an exponential function of the form r= r0 +Dr ×(1exp(− t/tr)) in order to evaluate the characteristic time tr for vacancy removal. It is immediately seen from Fig. 1 that tr for the Ni containing specimen, Fe–45Al–3Ni (compositions are stated in at.% unless otherwise noted), is significantly larger than tr for the binary iron aluminide, Fe–45Al. Fig. 2 shows that tr increases linearly with the Ni content. An addition of 10 at.% Ni increases tr by a factor of 27 as compared to the binary alloy. This means that, as compared to binary Fe– 45Al, the Ni-containing specimens require either longer annealing times, or higher annealing temperatures, in order to reach their equilibrium vacancy concentration. This result is in qualitative agreement with our earlier work [5]. In Fig. 3 the (effective) migration enthalpy Hm is evaluated for several alloys by assuming a relationship of the form t8 exp(Hm/(kT)). The value of 1.46 eV found for Fe–45Al is similar to migration enthalpies found with positron annihilation techniques: Wu¨rschum [3] determined a migration enthalpy of 1.7 eV for Fe–39 at.% Al, whereas Wolff et al. [14] (these proceedings) obtained 1.2 eV for Fe–40 at.% Al. Con-
Fig. 1. Bulk density of cast specimens vs. time at 723 K. Prior to the 723 K anneal, the specimens were homogenized at 1273 K and annealed for 3 days at 973 K. Note that these specimens were not HIPed.
J.H. Schneibel et al. / Materials Science and Engineering A239–240 (1997) 245–250
Fig. 2. Characteristic time tr for vacancy removal at 773 K, as a function of Ni content.
sistent with the slowing down of the vacancy removal rates by Ni additions, the activation enthalpies appear to increase with increasing Ni content. However, this finding is somewhat tentative in view of the scatter and the limited number of data points. TEM studies of the alloys following heat treatment at 1223 K and air-cooling showed that they contained a very low density of 001 dislocations. However, a number of microstructural features developed following prolonged annealing at 673 K. Typical transmission electron micrographs are shown in Figs. 4 and 5. Alloys containing less than 10 at.% Ni contained a number of cuboidal voids 50 – 100 nm in diameter, together with 001 edge dislocations. As the Ni content for these alloys increased, the density of the voids decreased and the dislocation density increased. No voids were observed in the alloy with 10 at.% Ni, only 001 dislocations. Both the voids and the 001 dislocations are associated with the removal of thermal vacancies from the lattice.
Fig. 3. Determination of activation enthalpies for migration.
247
Fig. 4. TEM micrograph of Fe – 45Al – 3Ni (arc-cast + 2 h/1223 K+ 120 h/673 K) illustrating low density of 001 dislocations and the presence of voids.
3.2. Room temperature yield strength measurements Fig. 6 shows the room temperature yield stresses for several Fe–Al–Ni alloys as a function of the temperature from which the specimens were quenched into oil. In all cases, the anneals prior to the quench were sufficiently long to achieve vacancy equilibrium, as verified by bulk density measurements. Typical times were 3 days at 973 K, 6 days at 873 K, and 12 days at 773 K. The yield stresses are seen to increase with increasing annealing temperature. These yield stress increases are due to quenched-in vacancies [5]. A dominant feature of Fig. 6 is the low slope of the Ni-containing specimens, as compared to binary Fe–45Al. Therefore, the curves for Fe–45Al and Fe– 45Al–6Ni intersect: Fe–45Al–6Ni quenched from 973 K exhibits a strength marginally lower than that of Fe–45Al quenched from the same temperature. If the vacancy contributions were ignored, this result would erro-
Fig. 5. TEM micrograph of Fe – 45Al – 10Ni (arc-cast + 2 h/1223 K +120 h/673 K) illustrating high density of 001 dislocations and absence of voids.
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Fig. 6. Room temperature yield stress as a function of the temperature of prior anneals, for iron aluminides with different Ni contents.
neously suggest solid solution softening due to nickel. If quenching is performed from lower temperatures, such as 773 K, the Ni-containing alloys are substantially stronger than Fe – 45Al, suggesting solid solution strengthening due to Ni.
3.3. Vacancy concentrations and solid solution strengthening The densities of HIPed specimens are shown in Fig. 7 as a function of annealing temperature prior to quenching. Prior to the 773 and 873 K measurements, the specimens were annealed for 3 days at 973 K. Fig. 7 shows that the density decreases with increasing annealing temperature, which is consistent with increasing vacancy concentrations. Fig. 8 shows the lattice parameters measured after annealing at 973 and 773 K, respectively. While sufficiently high vacancy concentrations collapse the lattice slightly and reduce therefore the lattice parameter [2],
Fig. 7. Bulk density of iron aluminides with different Ni contents as a function of annealing temperature. The anneals were sufficiently long to reach equilibrium.
Fig. 8. Lattice parameter of iron aluminides with different Ni contents after reaching equilibrium at 773 and 973 K.
the vacancy concentration difference between 973 and 773 K is too small to cause a noticeable effect on the lattice parameter. Therefore, there is virtually no difference between the two sets of data. From the bulk densities and the lattice parameters vacancy concentrations were calculated. The absolute values of these vacancy concentrations are subject to considerable uncertainty. One reason for this is the residual porosity in the specimens used for the density measurements. Also, small variations in the local specimen composition may cause errors. For example, a difference between the nominal and the actual Al concentration of only 0.1% would change the vacancy concentration by 0.07%. The 973 K vacancy concentrations found here are usually higher, by up to 0.2%, than Pike et al.’s (unpublished, [15]) concentrations for similar alloys. Pike et al. state that their concentration measurements were accurate to within 0.1%. The absolute values of our vacancy concentrations are likely to have a similar uncertainty. However, since in our experiments one and the same specimen was used for all the density measurements corresponding to a particular composition, the vacancy concentration differences, for different annealing conditions of a particular composition, are very reliable. Since differences in the bulk densities could be determined within 1 part in 104, the error in the (relative) vacancy concentrations for one and the same specimen is of similar magnitude. This means that the differences in the slopes of the curves in Fig. 9 showing the yield stresses as a function of the vacancy concentrations are reliable. Consistent with Fig. 6, the Ni-containing alloys exhibit a significantly lower slope than binary Fe–45Al. For low vacancy concentrations such as 0.15 at.%, Fig. 9 suggests strong solid solution hardening in Fe–45Al– 10Ni. For higher vacancy concentrations such as 0.25 at.%, Fe–45Al–6Ni is in fact weaker than Fe–45Al, even if the uncertainty in the absolute vacancy concentration is taken into account. This observation cannot
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be due to Ni-induced solid solution softening, since the Ni-containing iron aluminides with 0.15 at.% vacancies are much stronger than the binary Fe – 45Al. We conclude, therefore, that for a given vacancy concentration, the vacancy hardening in the Ni-containing specimens is not as strong as that in Fe – 45Al. On first sight, the formation of voids might be a reason for this finding. While voids are not seen in binary iron aluminides, they have been found in Fe – Al – Ni [9]. Voids, while reducing the bulk density, are not effective room temperature hardeners. In particular, they remove individual vacancies from the crystal lattice, thus making them unavailable for hardening. While this argument could be valid for Fe–45Al–3Ni, which does contain voids (Fig. 4), it is not valid for Fe – 45Al – 10Ni, which contains no voids, but only dislocations (Fig. 5). Another possibility to explain our result is the formation of Ni-vacancy clusters as they have been suggested by Munroe [9]. Although such clusters may have a higher obstacle strength for dislocations than individual vacancies, the number density of these clusters could be, depending on their size, substantially lower than that of individual vacancies and the net strengthening might be reduced. In Fig. 10 the yield stress is shown as a function of the Ni content, after a 12 day anneal at 773 K. According to Fig. 9, the residual vacancy concentrations are between 0.1 and 0.3 at.%. In view of the uncertainty in the absolute vacancy concentrations, and the uncertainty about the effectiveness of vacancies in Ni-containing FeAl, it is not possible at the present time to correct Fig. 10 for the effect of vacancies. Since vacancy hardening has probably not been eliminated completely, Fig. 10 is likely to represent an overestimate of the solid solution strengthening due to Ni. The present data can also not rule out the influence of Hall–Petch effects due to grain size variations [16]. However, since the vacancies in Ni-containing iron aluminides have been seen to be less effective strengtheners than the
Fig. 9. Room temperature yield strength vs. vacancy concentration, for iron aluminides with different Ni contents.
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Fig. 10. Room temperature yield strength of iron aluminides with different Ni contents, after a 12 day anneal at 773 K.
vacancies in binary iron aluminides, Ni is likely to be an effective solid solution strengthener in iron aluminides. This would be difficult to justify on the basis of the atomic size difference alone. Nickel has an atomic radius similar to that of Fe, and it is known to occupy the iron sublattice [17]. Elastic modulus measurements also did not show any substantial effect for Ni additions to FeAl [13]. There may be several possibilities, such as electronic or magnetic effects, which may change the dislocation–solute atom interaction and cause the strengthening. One other possibility that has been suggested [9,18] is the following: additions of Ni tend to change the slip behavior of iron aluminides from 111 slip to 100 slip [19]. This transition is illustrated in Fig. 11, in which the Burgers vectors and critical resolved shear stresses (CRSS) experimentally observed for FeAl and NiAl, respectively, are indicated. As Ni is added to FeAl, 111 slip is likely to become more difficult, since the CRSS for 111 slip in NiAl is much higher than that of FeAl. This trend is schematically shown by the solid line in Fig. 11. Also, as Ni is added, the increased tendency for 100 slip may in-
Fig. 11. Schematic illustration of the effect of Ni on the slip systems and strengths of iron aluminides.
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duce changes in the core structure of the 111 dislocations and reduce their mobility. Thus, the solid solution strengthening by Ni may very well be the result of a dislocation core effect.
4. Conclusions Additions of Ni to Fe – 45Al slow down dramatically the kinetics of vacancy migration and removal. This is likely to be associated with binding between Ni and thermal vacancies, and the formation of Ni-vacancy complexes. While Ni gives rise to high solid solution strengthening, it renders vacancy hardening less effective as verified by yield strength measurements for different nickel and vacancy concentrations. Again, Nivacancy complexes provide a possible explanation of this finding, since such complexes would reduce the number density of defects interacting with the dislocations.
Acknowledgements This research was sponsored by the Division of Materials Sciences, US Department of Energy under contract number DE-AC05-96OR22464 with Lockheed Martin Energy Research Corporation. Valuable discussions with C.L. Fu and Xing-Dong Wang are appreciated, as well as the review of this manuscript by T. Takasugi and L.M. Pike.
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References [1] C.L. Fu, Y.-Y. Ye, M.H. Yoo, K.M. Ho, Phys. Rev. B 48 (1993) 6712. [2] Y.A. Chang, L.M. Pike, C.T. Liu, A.R. Bilbrey, D.S. Stone, Intermetallics 1 (1993) 107. [3] R. Wu¨rschum, C. Grupp, H.-E. Schaefer, Phys. Rev. Lett. 75 (1995) 97. [4] P.R. Munroe, Processing, properties, and applications, in: S.C. Deevi et al. (Eds.), Int. Symp. Nickel Iron Aluminides, ASM, 1997, 329 pp. [5] J.H. Schneibel, P.R. Munroe, L.M. Pike, in: C.C. Koch et al. (Eds), High-temperature ordered intermetallic alloys VII, Mater. Res. Soc. Proc., 460, 1997, pp. 379 – 384. [6] R. Wu¨rschum, K. Badura-Gergen, E.A. Ku¨mmerle, C. Grupp, H.-E. Schaefer, Phys. Rev. B 54 (1996) 849. [7] P. Nagpal, I. Baker, Metall. Trans. A 21A (1990) 2281. [8] K. Yoshimi, S. Hanada, H. Tokuno, Mater. Trans. JIM 35 (1994) 51. [9] P.R. Munroe, Intermetallics 4 (1996) 5. [10] C.H. Kong, P.R. Munroe, Scr. Metall. Mater. 30 (1994) 1079. [11] P.R. Munroe, C.H. Kong, Intermetallics 4 (1996) 403. [12] J.H. Schneibel, E.P. George, E.D. Specht, J.A. Horton, in J.A. Horton et al. (Eds), High temperature ordered intermetallic alloys, vol. VI, Mater. Res. Soc. Proc., 364, 1995, 73 pp. [13] J.H. Schneibel, E.D. Specht, W.A. Simpson, Intermetallics 4 (1996) 581. [14] J. Wolff, M. Franz, A. Broska, B. Ko¨hler, Th. Hehenkamp, Mater. Sci. Eng. A239/240 (1997) 213 – 219. [15] L.M. Pike, Y.A. Chang, C.T. Liu, Acta Mater. 45 (1997) 3709. [16] J. Winterhoff, E. Nembach, Scr. Mater. 35 (1996) 999. [17] I.M. Anderson, Acta Mater. 45 (1997) 3897. [18] J.H. Schneibel, C.T. Liu, in: G.M. Stocks, P.E.A. Turchi (Eds), Alloy Phase Stability and Design, TMS, Pittsburgh, 1994, 235 pp. [19] D.K. Patrick, K.-M. Chang, D.B. Miracle, H.A. Lipsitt, in H.A. Johnson et al. (Eds.), High-temperature ordered intermetallic alloys, vol. IV, Mater. Res. Soc. Proc., 213, 1991, 267 pp.