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The effect of ternary additions on vacancy hardening in near stoichiometric FeAl
Inrermetailics 4 (1996) 403-415 0 1996 Elsevier Science Limited Printed in Great Britain. All rights reserved 0966-9795/96/$15.00
0966-9795(95)00057-7 ELSEVIER
The effect of ternary additions on vacancy hardening in near stoichiometric FeAl P. R. Munroe & C. H. Kong School of Materials
Science and Engineering, (Received
University of New South Wales, Sydney,
21 September
1995; accepted
15 November
NS W 2052, Australia
1995)
The effect of transition metal ternary additions which replace iron on the variations in hardness and defect morphology with heat treatment of near stoichiometric FeAl has been studied. It was generally found that ternary additions of 1% have little effect on the vacancy-related hardening and softening which occurs in binary FeAl. The notable exception to this behaviour is the addition of nickel, which strongly hardens this alloy and dramatically changes the defect microstructure that develops during low temperature annealing. The hardening effect of nickel additions may be attributed to either this changed defect microstructure or the interaction of nickel atoms with the cores of glide dislocations. In contrast, ternary additions of 5% significantly harden FeAl, and it is believed that this is associated with the interaction of ternary atoms with thermal vacancies, in addition to conventional solute hardening. These ternary additions also strongly affected the defect microstructures which develop in FeAl during low temperature, vacancy-relieving heat treatments. Ternary additions resulted in the dislocation structures becoming much wavier, and cobalt and vanadium additions induced the formation of ~11 l> dislocations, in addition to the expected defects. Furthermore, low-temperature heat treatment of Fe,,Al,,Cr, led not only to the formation of FeAl, particles, but also to localised disordered regions. Copyright 0 1996 Elsevier Science Ltd Key words: A. iron aluminides based on FeAl, C. heat treatment, D. defects: point defects. defects: dislocation geometry and arrangement, F. electron microscopy, transmission.
1 INTRODUCTION
of edge dislocations form.3~s~9 It has been previously shown that small (< 10 at%) nickel additions, which replace iron, may suppress the softening which occurs during heat treatment at 400°C in both iron-rich and near-stoichiometric FeAl.‘“i2 Consequently, these additions also produce significant increases in the equilibrium hardness following these low temperature vacancy-relieving heat treatments. Microstructural studies of these alloys following low temperature annealing also showed that nickel additions, as low as 0.1 ato/, dramatically affected the dislocation structures which formed during this heat treatment.‘2,‘3 The aim of this work is to examine the effect of other transition metal ternary additions on vacancy hardening in FeAl. Two sets of alloys with nominal compositions of Fe,,Al,&, (X = Cu, Ni, Co, Mn, Cr, V and Ti) and Fe,,Al,,X, (X = Cu, Ni, Mn, Co and Cr) were prepared (unless otherwise stated all compositions given will be in atomic per cent). These alloys were subjected to both high temperature annealing and subsequent low temperature vacancy-
The B2-structured compound FeAl is seen as a candidate material for elevated temperature strucespecially in hostile environtural applications, ments, because of its good oxidation resistance, low density and relatively low cost. Although FeAl is ductile at high temperatures, its commercial application has been inhibited because of its low room temperature ductility and limited strength at However, FeAl is temperatures above -500°C.’ stable over a wide range of compositions,2 which provides an opportunity to improve the mechanical properties of this compound through alloying. Following heat treatments at temperatures greater than about 700°C FeAl retains large concentrations of vacancies, even when furnacecooled, which leads to increases in hardness and decreases in ductility.” 7 This anomalous hardening can be relieved by low temperature (400°C) heat treatments during which excess vacancies are removed from the lattice and high concentrations 403
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relieving heat treatment. Their behaviour was monitored by microhardness measurements and their microstructures were examined by both optical microscopy and transmission electron micro- scopy (TEM).
q
Air-cooled
Mn
Co
4OOW 120h
2 EXPERIMENTAL To investigate the effects of transition metal additions on the vacancy retention behaviour of FeAl, two sets of alloys with nominal compositions of Fe,9A1,,X, (X = Cu, Ni, Co, Mn, Cr, V and Ti) and Fe,,Al,,X, (X = Cu, Ni, Mn, Co and Cr) were prepared. All of the samples were annealed in air at 950°C for 2 h and air-cooled. Subsequent low temperature annealing was performed on the air-cooled alloys at 400°C for 120 h. Air-cooled, rather than the water-quenched samples, were chosen for study as they were slightly less brittle and easier to handle than water-quenched samples. Preliminary experiments have confirmed that the behaviour and microstructures of the waterquenched and air-cooled samples were very similar.14 Vickers microhardness testing was performed on each sample using a 300 g load. Each data point obtained was the average of at least 10 measurements. Standard deviations were less than 5’%, but error bars were omitted from Figs 1 and 2 to improve clarity. The microstructures of all the materials were examined by both optical and transmission electron microscopy (TEM). Thin foils for TEM study were prepared using methods described elsewhere,15 and examined in a JEOL 2000FX transmission electron microscope, furnished with an energy dispersive X-ray spectrometer (EDS), operating at 200 kV.
q
4OOW120h
Air-cooled
550 $500 3
450
@-JO 350 300 250 Ti
V
Cr
Mn
Co
Ni
Cu
FeAFe-49AI
Fig. 1. The microhardness of Fe,,AI,,X, (X = Cu, Ni, Co, Mn, Cr, V and Ti) following air-cooling from 950°C and subsequent heat treatment at 400°C for 120 h, data for nominally stoichiometric FeAl and Fe5,A1,, is also included for comparison.
Cr
Ni
0.1
FeAl
Fig. 2. The microhardness of Fe,,Al,,X, (X = Cu, Ni, Co, Mn and Cr) following air-cooling from 950°C and subsequent heat treatment at 400°C for 120 h, data for nominally stoichiometric FeAl is also included for comparison.
The atom site location of the ternary elements was determined by atom sitelocation by channelling enhanced microanalysis (ALCHEMI) techniques; experimental details have been given elsewhere.‘(j
3 RESULTS Figure 1 shows the comparison of microhardness of the ternary alloys Fe,9Al,,X, (X = Cu, Ni, Co, Mn, Cr, V and Ti) with that of both nominally stoichiometric FeAl and Fe-49Al following aircooling from 950°C and subsequent low temperature annealing at 400°C for 120 h. The air-cooled samples (all the ternary alloys and both FeAl alloys) exhibited generally similar hardness values. Following subsequent annealing at 400°C for 120 h the alloys containing Cu, Co, Mn, Cr and V additions again exhibited hardness levels similar to both FeAl and Fe-49A1, although all these ternary alloys were slightly harder than the binary alloys. In contrast, the alloys with additions of either Ni or Ti exhibited much less softening when heat treated at 400°C indicating that the equilibrium hardness values of these ternary alloys were much higher than that of binary FeAl. Figure 2 shows the comparison of microhardness of the ternary alloys Fe,,Al,,X, (X = Cu, Ni, Mn, Co and Cr) with that of nominally stoichiometric FeAl following both air-cooling from 950°C and subsequent annealing at 400°C for 120 h. Following air-cooling from 950°C the ternary alloys containing Cu, Co, Mn and Cr additions showed a comparable or even slightly lower hardness than the binary FeAl, whereas the Ni-containing alloy was harder. Following the vacancy
Vacancy hardening in near stoichiometric FeAl
relieving heat treatment, a small reduction in hardness was observed for all the ternary alloys. In comparison, the reduction in hardness for FeAl was substantial. Thus, these ternary additions not only substantially increased the equilibrium hardness of binary FeAl, but also appeared to reduce the extent of any vacancy-related softening. Optical microscopy showed that following annealing at 950°C and air-cooling, all the ternary alloys and the binary FeAl alloys appeared to be single phase with a comparable grain size of about 200 pm. TEM investigations confirmed that all the alloys exhibited a similar microstructure; single phase with a low density (-lOI m-“) of residual dislocations with cOOI> Burgers vectors. These dislocations are presumably induced through thermal stresses associated with the melting process and subsequent high temperature annealing. The atom site location of the ternary elements added to FeAl was assessed using ALCHEMI methods on the samples air-cooled from 950°C. The results for the Fed9Al,,X, alloys are shown in Fig. 3. Each data point is the average of five measurements. For Cu, Ni, Co, Mn and Ti it was found that more than 95% ternary of these atoms were located on the iron sublattice, although no corrections were made for delocalisation effects. In contrast, V atoms were found to preferentially occupy the Al sublattice, and Cr atoms occupy both sublattices in approximately equal proportions. Similar lattice site preferences were observed for the Fe,SAl,,X, alloys. Following annealing at 400°C for 120 h, TEM studies on the Fe,,Al,,X, alloys revealed that all the ternary alloys, except the titanium-containing alloy, were still single phase, but contained a much higher density (-5 X lOi mm’) of dislocations compared to that following high temperature .‘.’ :: ::
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heat treatment. For example, Fig. 4 shows the microstructure of Fe,Al,,Cr, following annealing at 950°C and air-cooling, and also following subsequent annealing at 400°C for 120 h. In this alloy, following both heat treatments, the dislocations were determined to have Burgers vectors and were edge in character and lay on (001 > planes. Similar microstructures were observed in the alloys containing Mn and Cu. The microstructure of Fe,,Al,Ni, following this heat treatment consisted of both dislocations and cuboidal voids and this has been described in detail elsewhere.13 However, in the ternary alloys containing either vanadium or cobalt following annealing at 400°C for 120 h a number of dislocations with Burgers vectors were frequently observed, although edge dislocations were still predominant. Figure 5 shows a series of micrographs of the dislocation structure in Fe,,Al,,V, following low temperature annealing. The majority of the dislocations have Burgers vectors. For example, dislocations such as those labelled a, exhibited invisibility when the operating reflection, g, was a
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Fig. 3. Atom site location of ternary additions in Fe,AI,,X, (X = Cu, Ni, Co, Mn, Cr, V and Ti) following air-cooling from 950°C as determined by ALCHEMI methods.
Fig. 4. Bright field transmission electron Fe,Al,,Cr, following annealing at 950°C for cooling, (b) subsequent annealing at 400°C both micrographs the electron beam direction and diffraction vector is [ 1lo].
micrographs of 2 h and (a) airfor 120 h. For is close to [OOl]
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Fig. 5. Bright field transmission electron micrographs of Fe,,Al,,V, following ann_ealing at 4OO’C for 120 h. The electron-beam direction is close to (a-d) [OOl] and (e, Q [ill], and diffraction vectors are: (a) [l IO]; (b) [l lo]; (c) [200]; (d) [020]; (e) [loll and
[200] and [lOi], indicating that their Burgers vector was [OlO]. However, dislocations, such as the one marked _b, exhibited invisibility when g was [I lO]_and [Ol 11, indicating that its Burgers vector was [l Ill. One of these ~111~ dislocations (arrowed in Fig. 5(a)) exhibited a M-shape morphology where it appeared to be pinned by a [OOl]dislocation marked d. The dislocation d which also exhibited an abrupt change in line direction from [OIO] to [IOO] where it had interacted with this
~11 I> dislocation. In addition, some of the ~001~ dislocations exhibited a C-shape morphology, for example those labelled c in Fig. 5(d), inferring that they had experienced some form of pinning force during the heat treatment. Furthermore, in Fe,,Al,,Vr some of the ~11 l> dislocations were observed to exist as dislocation dipoles. Dislocation pairs marked d in Fig. 6 exhibited a change in spacing between the two dislocations when they were imaged using +g and -g under two beam
Vacancy hardening in near stoichiometric FeAl
Fig. 6. Bright field transmission electron micrographs of ~11 I> dislocation dipoles, labelled d, in Fe,,A&,,V, following ianneal ling at 400°C for 120 h. The electron beam direction is close to -.10011 for all three micrographs and the diffraction vectors are (a) [21001, (b) [200] and (4 w201. 1
diffraction conditions, contrast characteristic of dislocation dipoles. l7 It can also be seen in Fig. 6(c) that one of the dipoles has been pulled out from a [OlO] edge dislocation (arrowed) again suggesting that the < 111> defects were formed through the interaction of dislocations and some pinning force, presumably vacancy-related. Similar dislocation arrangements were observed in Fe4,A1,,Co,. That is, a small number of < 11l> dislocations were noted in addition to a high density of edge dislocations. Figure 7 shows a series of micrographs showing dislocation structures in Fe,Al,,Co,. The majority of dislocations were ~001~ edge dislocations, similar to those described for Fe,,Al,,V,, but dislocations such as those labelled-a were out of contrast when g was [l IO] and [Ol I], indicating that their Burgers vector was [il I]. Again these ~11 I> dislocations were often observed to be pinned by ~001~ dislocations. For example, in Fig. 8 the dislocation labelled V, has a ~111~ Burgers vector and exhibits a V-shape morphology. It appears to have been pulled out from a ~001~ dislocation, labelled Z, and be
pinned at point J. This arrangement again implies that the formation of such < 11l> dislocations is associated with the interaction of dislocations with some vacancy related pinning force during the low temperature heat treatment. Second phase particles were observed in two alloys, Fe,,Al,, and Fe,Al,,Ti,, following low temperature annealing at 400°C for 120 h, although these two alloys exhibited single phase microstructures prior to this heat treatment. The microstructures of these alloys have been described elsewhere in detail.‘* The particles in the stoichiometric binary alloy were suggested to be FeAl,, and from the hardness data in Fig. 1 they would appear to have no significant hardening effect on this alloy. The particles observed in Fe,,Al,,Ti, have been suggested to be either FeAl, or TiFeAl,, it is possible that the large equilibrium hardness observed in Fe,,Al,,-,Ti, may arise from the second phase hardening, although it is possible that the Ti-containing alloy is intrinsically much harder than the binary FeAl because of either solid solution strengthening or vacancy-related effects.
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Fig. 7. Bright field transmission electron micrographs of ~11 I> dislocations, labelled a, in Fe,,Al,,Co, following annealing at 400°C for 120 h. The electron beam direction is close to (a, b) [OOl]_and (c) [ill]. The diffraction vectors are (a) [l lo], (b) [IlO] and (c) LO111.
Following air-cooling from 950°C and subsequent annealing at 400°C for 120 h, TEM observations revealed that all the Fe,SAl,OX, alloys except the chromium-containing alloy, remained single phase, but all again contained a significantly higher density (-5 X 1013 me2) of dislocations
Fig. 8. Bright field transmission
electron micrograph showing the intersection (arrowed) of a < 11 l> dislocation, labelled V, and a dislocation, labelled I, in Fe,,Al&o, following annealing at 400°C for 120 h. The electron beam-direction is close to [OOl]. The diffraction vector is [l lo].
compared with that observed following air-cooling from 950°C. Figure 9(a) shows the microstructure of Fe,,Al,,Co, following air-cooling from 95O”C, it is clear that this alloy contains a low dislocation density. Figures 9(bd) show the microstructures Fe,,Al,,Cu,, FeJSAISOMnSand Fe4,A1,,Co, respectively following subsequent annealing at 400°C. In Fe,SA1,,Co, (Fig. 9(b)) the dislocations, all of which had Burgers vectors, lay in specific directions, similar to those observed in binary FeAl. However, in Fe4SAl,,MnS (Fig. 9(c)) the dislocations were noted to be slightly wavier in nature, although all the dislocations still had a Burgers vector of . In contrast, the dislocations observed in Fe,,Al,,Co, following annealing at 400°C exhibited considerable waviness and were often tangled, and no favoured line direction could be clearly identified. In this latter case, although dislocations were still predominant, a number of ~1 1l> dislocations were again observed. Figure 10 shows a more detailed analysis of the dislocations in Fe,SA1,,Co, following annealing at 400°C. It was found that the dislocations such as those labelled a exhibited invisibility when
Vacancy hardening in near stoichiometric FeAl
Fig. 9. Bright field transmission
electron micrograph of the microstructure of: (a) Fe,,Al,,Co, following air-cooling from 95O’C; (b) Fe,5Al,,CuS; (c) Fe,,Al,OMn, and (d) Fe,5A1,,Co, following subsequent annealing at 400°C for 120 h. For all the micrographs the electron beam direction is close to [OOl] and diffraction vector is [ 1lo].
g was both [OZO]and [Oli], so the Burgers vector was [loo], and dislocations, such as the one labelled b exhibited invisibility when g was both -[I lo] and [lOi], so the Burgers vector was [lil]. It was also noted that a weak fringe contrast was observed to be associated with these dislocations. This suggests that these defects have the Burgers vector l/2< 11l> and their motion has created disorder and left a region of matrix which is out of phase with the rest of the matrix. In addition, some of the ~1113 dislocations, for example those labelled c, which exhibited contrast consistent with a [ll l] Burgers vector were in the form of tangled dislocation loops. However, the dislocation labelled d exhibited invisibility when g was both [OZO]and [lOi], indicating that its Burgers vector was [loll. This defect is connected to two other dislocations, each of which exhibited contrast consistent with a ~11 l> Burgers vectors, suggesting that it had formed through the interaction of two dislocations. The microstructure observed in Fe~SA1,,Cr, following air-cooling from 950°C and subsequent annealing at 400°C for 120 h was much more
complex than the other Fe,Al,&, alloys studied. Bands of slightly elongated, plate-like second phase precipitates (arrowed), about 50 nm in diameter, were observed, which trace analysis showed were parallel to (112) planes (Fig. 11). If this region was imaged using (001) superlattice reflections, lines parallel to the { 112) planes were observed which exhibited weak contrast relative to the blocks of matrix enclosed between them (Fig. 12(a)). Such images are strongly reminiscent of antiphase boundary (APB) related contrast in disorderable Fe3Al-based alloys following deformation.‘9v20It appeared that the apparently disordered lines were connected, or pinned, to the second phase particles, and their line directions were, in general, parallel to the longitudinal axis of the precipitates. If these regions were imaged using any fundamental reflections, such as f 1101, these lines became invisible. This is consistent with the suggestion that this contrast was associated with APBs. To study the precipitates and APB-type contrast observed in Fed5Als0Cr5in more detail selected area electron diffraction (SAD) patterns were taken
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Fig. 10. Bright field transmission electron micrographs of Fe,,A1&05 following annealing at 400°C for 18-h. The electron beam directions are close to (aa) [OOl] and (e, f) [ill] and the diffraction_vectors are: (a) [200]; (b) [llO]; (c) [l lo]; (d) [020]; (e) [loll and (f) [Ol I].
from a number of low index axes of the matrix. Figure 12 shows the SAD patterns recorded from the [l lo] and [II l] zone axes. Some weak extra spots, marked with a subscript p, together with the expected FeAl matrix reflections, marked with a subscript m, were observed. Analysis of the diffraction patterns showed that the precipitate spots were consistent with the second phase being were observed in FeAl,. Similar precipitates Fe,5A1,,V,, and the analysis of these patterns has
been described in detail elsewhere.21 The orientation relationship between the FeAI, precipitates and FeAl-based matrix in Fe,,Al,,Cr, was also the same as that observed in Fe45Al,0V5.21That is: [10I JFeAl, // [ 11 l]FeAl and (040)FeAl, // (0 1i)FeAI. It is also worth noting that the precipitate reflections arising from the second phase in the SAD pattern taken from the [l 1 I] zone axis of the matrix were elongated in the direction of the ( 112) planes
Vacancy hardening in near stoichiometric FeAl
Fig. 11. Bright
field transmission electron micrograph of Fe,,Al,,Cr, following annealing at 400°C for 120 h. The electron beam direction is close to [ 10 l] and the diffraction vector is [020].
Fig. 13. Selected area electron diffraction patterns from matrix phases in Fe,,Al&r, following air-cooling from 950°C and subsequent annealing at 400°C for 120 h. Zone axes of the FeAl-based matrix are (a) [l IO] and (b) [I 111.
4 DISCUSSION
Fig. 12. Transmission electron micrographs of Fe,,A1,,Cr5 following annealing at 400°C for 120 h imaged using (a) a (010) superlattice reflection and (b) a (101) fundamental reflection.
of the matrix. These elongated diffraction spots may result from the precipitates exhibiting a small spread in orientation relationship with the matrix. Such elongation was not noted in the diffraction patterns obtained from Fe4SA150V5.21
Following heat treatment at 950°C and air-cooling all the Fed9A15,,X, alloys and Fe,SAl,,X, alloys exhibited approximately equal hardness values, which were comparable to that of binary FeAl. Following such heat treatments it has been shown that a large concentration of vacancies is retained in the lattice in FeAl and this strongly increases hardness.6,7 The similar hardness of these ternary alloys to the binary alloy following this heat treatment suggests that they also retain significant concentrations of vacancies, which increase hardness. Furthermore, these data would suggest that any solute hardening effects associated with the ternary additions are largely obscured by the presence of these thermal vacancies. However, subsequent annealing of these alloys at 400°C revealed larger differences in hardness behaviour between the ternary alloys and FeAl. For those alloys with the generic composition Fe,Al,,X, considerable softening occurred during
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P. R. Munroe, C. H. Kong
this heat treatment, except for Fed9Al,,Ni, and Feb9Al,,Ti,. That is, the hardness of these alloys was comparable to, but marginally harder than, binary FeAl. However, it is worth noting that the marginally higher hardness values obtained after this heat treatment increased with increasing difference in atomic number between iron and the ternary addition. This hardening behaviour is presumably related to solid solution hardening associated with the ternary additions. The increasing hardness with increasing difference in atomic number between iron and the ternary addition is presumably associated with atomic size misfit as suggested by Schneibel et ~1.~~The exception to this behaviour was Fe,Al,Ni,, which exhibited a much higher hardness than either the copper and cobalt containing alloys, suggesting that a more complex hardening behaviour was in operation in this alloy. This behaviour may be related to the defect structure evolved during heat treatment at 400°C. Voids, in addition to dislocations, were observed in Fe,,Al,,Ni, following this heat treatment, and it is possible that these voids induce additional hardening in this alloy.‘*,13 In addition, it has been suggested that the nickel atoms in this alloy induce core changes around glide dislocations which inhibit their motion and so harden the alloy. 23A very high hardness was also observed in the Fe49Al,0Ti, following the 400°C anneal, although some of this hardening may be attributable to the atomic size misfit effect, this alloy contained a large number of fine second phase particles which may also have contributed a hardening effect. In contrast, the Fe,,Al,,X, alloys softened to a much lower degree during annealing at 400°C compared with FeAl. That is, their equilibrium hardness values at this temperature were much higher than that of the binary composition. It is possible that the higher concentration of ternary additions in these alloys leads to greater vacancysolute interactions, which thus inhibit the removal of vacancies from the matrix during low temperature annealing. That is, the high hardnesses of these alloys following annealing at 400°C are still attributable to the presence of thermal vacancies trapped in the lattice. Alternatively, the higher hardness of these alloys may be associated with solid solution hardening effects. Consistent with this was the observation that the hardness levels increased with increasing atomic number difference between iron and the ternary addition, with the exception of the Nicontaining alloy, which was again much harder
than any of the other single-phase alloys. The high hardness of the Cr-containing alloy would presumably be due in part to the high density of fine particles which precipitated following the low temperature heat treatment. Schneibel et al. found similar trends in the yield strengths of Fe45Al-5X alloys (where X = Cu, Ni, Co, Fe, Mn, Cr, V, Ti) with small B and Zr additions.22,24 In that case, a linear relationship between the yield strength and the atomic size misfit for all ternary alloys was established with the exception of nickel-containing alloy, which again showed a much stronger strengthening effect. However, these workers noted that the manganese alloy was marginally softer than the binary alloy. The results of ALCHEMI studies revealed that Cu, Ni, Co, Mn and Ti atoms preferentially occupied the iron sublattice, whereas V atoms preferentially occupied the aluminium sublattice and the Cr atoms occupied both sublattice sites in approximately equal proportions. Recently Kao et al.” predicted the atom site preference of ternary additions of transition metals in a range of B2-structured intermetallic compounds, including FeAl, based on the calculation of standard enthalpies for binary compound formation. For FeAl it was predicted that ternary elements with atomic numbers greater than Fe, such as Co, Ni and Cu, would prefer to occupy the iron sublattice, and ternary elements with atomic numbers less than Fe, such as Cr, Ti, V and Mn would prefer to occupy the iron sublattice. However, these predictions for Cu, Mn, Cr and Ti were limited by being subject to significant statistical error. Schneibel and co-workers used ALCHEMI to determine the lattice site occupancy of ternary additions in a range of Fee45Al-5X alloys. 22 Their results were broadly in agreement with those obtained here, except that both titanium and chromium were observed to exhibit a strong preference for the aluminium sublattice. Earlier work by Munroe and Baker also showed that chromium occupied the aluminium sublattice in alloys based on Fe-40Al.26 The differences in behaviour for these results with the observations made here may arise because the alloys previously examined were iron-rich, and thus in these alloys the location of the ternary additions on the aluminium sublattice would be expected to reduce the number of iron antisite defects. In the present study, however, the location of the titanium atoms on the iron sublattice would minimise the number of anti-site defects. The even distribution of chromium atoms between both iron and aluminium sublattices in these alloys, would lead to
Vacancy hardening in near stoichiometric FeAl no iron anti-site defects, but the presence of the chromium atoms on the aluminium sublattice would produce constitutional vacancies. Presumably, for Fe,,Al,,X, it is energetically more favourable to contain a small number of constitutional vacancies than to occupy the iron sublattice in significant numbers. The microstructures of all the alloys following air-cooling from 950°C were observed to contain very low density of edge dislocations. Following subsequent annealing at 400°C the density of dislocations in all these alloys was significantly increased, but was still mainly composed of edge dislocations, on (001 } planes. This is consistent with a recent study which has shown that for FeAl dislocations on (001 } planes are most stable in the edge orientation.27 However, a small number of dislocations were observed in Fe,,Al,,VI, Fe,,Al,,Co, and Fe,,Al,,Co,. The origin of these < 111~ dislocations is not clear. It is well known that binary FeAl deforms by < 111~ slip at room temperature, irrespective of the iron:aluminium ratio.28 However, these samples were not subjected to deformation, and these dislocations clearly formed during low temperature heat treatment. It is possible that these ~1 1 l> dislocations may be formed through the interaction of ~00 1> dislocations. During high temperature deformation of NiAl two dislocations may interact to form a ~011~ dislocation, which then further react with another defect to form a dislocation with a ~11 l> Burgers vector.2” iI However, if this mechanism was in operation, it would be expected that a higher density of dislocations with ~011~ Burgers vectors would be observed. However, no ~01 l> dislocations were noted in either Fe,,Al,,V, or Fe,,Al,,Co,. and only a very small number of ~1 lO> dislocations were observed in Fe,,Al,,Co,, and in this instance it is likely that such ~01 l> dislocations formed through the interaction of pairs of ~111~ dislocations. Fourdeaux and Lesbats observed loops with l/2< Ill> Burgers vectors in Fe40Al following low temperature vacancy-relieving annealing.’ These were shown to arise from the dissociation of ~001~ dislocations into pairs of 1/2<111> dislocations by the reaction, which was followed by re-combination of these loops to form ~001~ dislocations. However, in this study no such loops were noted in Fe,,Al,,V, or Fe,,Al,,Co, and although some ~1 1 l> loops were observed in Fe,,Al,,Co,, dislocation interactions similar to those noted by Fourdeaux and Lesbats were not seen here.
413
It has been shown that many of these ~1 1l> dislocations were in the form of dipoles. ~1112 dislocation dipoles have been observed elsewhere in extruded iron-rich FeA1.27,32 In these cases, it was assumed that their formation was associated with dislocation interactions during high temperature deformation. Recently, Baker and Horton observed in stoichiometric FeAl following room temperature deformation dislocations with < Ill> Burgers vectors which were in the form of either dipoles or exhibited C-shape morphologies, similar to that observed in this study.33 It was suggested that these dipoles and C-shape dislocations arose not through mutual dislocation interaction but through interaction of gliding dislocations with quenched-in vacancies. It is likely that the dipoles formed here are also associated with presence of vacancies. The M-shape dislocation in Figure 5 appears to be formed by the pinning action of a ~11 l> dipole with a dislocation. Presumably, further motion of the ~11 l> dislocation, whilst being pinned by the dislocation would lead to dipole formation. The observation of ~11 I> dislocations in the alloys containing either vanadium or cobalt, associated with dipole formation, suggests that the mobility of the dislocations during vacancy-assisted climb is limited in these alloys, although the precise mechanism for < 111~ dislocation formation remains unclear. In relation to the microhardness measurements, it seems that the presence of ~11 I> dislocation dipoles in Fe,,Al,,V, and Fe,,Al,,Co, did not appear to affect the equilibrium hardness significantly. Similarly, the hardness of Fe,SA1,,Co, was little different to the other Fe,,AI,,X, alloys following heat treatment at 400°C which also suggests that these ~1 I1 > defects do not greatly affect the hardness of this alloy. Differences in dislocation morphology in the Fe,,Al,,XS alloys became evident and the dislocations became wavier as the ternary addition changed from Cu to Mn to Co. This suggests that in these alloys the ternary additions may alter the line energy of the dislocations. In addition, some fringe contrast was observed to be associated with some of the ~111~ dislocation, this fringe contrast may arise through the passage of a l/2< 111> dislocation through the lattice, which thus trailed disorder, resulting in the formation of fringe contrast. It is interesting to note that for Fe,,Al&r, following annealing at 400°C a high density of precipitates formed and contrast from the matrix phase consistent with that expected from APBs
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was observed. Neither these precipitates nor the APB-type contrast were observed in the samples air-cooled from high temperature. This indicates that they were formed during subsequent low temperature annealing. This alloy is significantly harder than near-stoichiometric binary FeAl, following the same heat treatment, suggesting that in addition to solid solution effects these precipitates may provide some contribution to hardness. It is also clear that the solubility of chromium in nearstoichiometric FeAl at 400°C is greater than lo/o but less than 5%, but at higher temperatures the solubility of chromium in FeAl is greater than 5%. Examination of the Fe-Al-Cr alloy phase diagram 873 K isotherm (the lowest temperature available) shows that Fe,,Al,,Cr, lies in the (FeAl) B2 plus disordered (a-Fe, Cr) phase field.34 This is somewhat contrary to the observed microstructure, where FeAl, particles were also observed. The APB-like contrast is therefore related to the formation of the disordered (a-Fe, Cr) phase. It is, however, interesting to note that these disordered regions are associated with the FeAl, particles. This suggests that localised chemical variations occur in these regions, and the local decrease in aluminium content in the disordered regions, leads to adjacent regions locally rich in aluminium, and so the formation of FeAl,. However, these areas were sufficiently fine to preclude detailed microanalytical studies. Further studies, using high resolution microanalytical techniques are underway to investigate the development of the microstructure in these FeAl+Cr alloys.
5 CONCLUSIONS
(1) The
variation in hardness with heat treatment for alloys with the composition Fe49Al,,X, is little affected by ternary additions. The exceptions to this behaviour are those alloys where either titanium or nickel are the ternary additions. Significant hardening in Fe,,Al,,Ti, is associated with presence of second phase particles. In contrast, the hardening in FeJyA1,,Ni, is believed to be associated with the interaction of nickel additions with the cores of glide dislocations and the significant effects which nickel has on the defect microstructure which evolves in FeAl-based alloys during low temperature heat treatment. (2) Ternary additions are highly effective in hardening alloys with the composition Fe,,Al,,X,. This is believed to be associated with the inter-
C. H. Kong
action of vacancies with the solute atoms as well as solid solution hardening effects. Ternary additions affect the dislocation structures (3) which evolve in FeAl during low temperature vacancy relieving heat treatment. Vanadium and cobalt additions lead to the formation of < 11 1> dislocations in addition to dislocations usually observed. heat treatment of Fe,,Al,,Cr, (4) Low temperature leads not only to the precipitation of FeAl, particles but also to the formation of localised disordered regions.
REFERENCES 1. Baker, I., Processing, Properties und Applicution of Iron Aluminides, eds M. A. Crimp & J. H. Schneibel. TMS, Warrendale, USA, 1994. p. 101. 2. Massalski, T. B., Okamato, H., Subramanian, P. R. & Kacprazak, K. K., Binary Alloy Phase Diugrums, ASM, Materials Park, OH, 1986, p. 112. 3. Rieu, J. & Goux, C., Mem. Sci. Rev. Met., 65 (1969) 869. 4. Weber, D., Meurtin, M., Paris, D.. Fourdeux, A. & Lesbats, P., J. de Phps. Colloque, C7 (1977) 332. 5. Nagpal. P. & Baker, I., Metall. Trans, 21A (1990) 2281. 6. Chang, Y. A., Pike, L. M., Liu, C. T., Bilbrey, A. R. & Stone, D. S., Intermetal/ics, 1 (1993) 107. 7. Xiao, H. & Baker, I., Acto Metall. et Muter., 43 (1995) 391. 8. Morton, A. J. & West, G. W. PPOC. 8th inter. Gong. on Electron Microscopy VI, eds J. V. Sanders & D. J. Goodchild, The Australian Academy of Science, Canberra, 1976, p. 460. 9. Fourdeux, A. & Lesbats, P., Phil. Msg., A45 (1982) 81. 10. Kong, C. H. & Munroe, P. R., Scripta Metall. et Muter., 28 (1993) 1241. 11. Kong, C. H. & Munroe, P. R., Scripta Metull. et Muter., 30 (1994) 1079. 12. Munroe, P. R., Intermetnllics. in press (1995). 13. Kong, C. H. & Munroe, P. R., Intermetallics, 2 (1994) 333. of New South 14. Kong, C. H., Ph.D. Thesis, University Wales, Sydney, 1995. 15. Baker, I. & Gaydosh, D., Phys. Stat. Solidi, 96 (1986) 185. 16. Munroe, P. R. & Baker, I., Microbeam Analysis ~ 1990, eds J. R. Michael & P. Ingram, San Francisco Press, San Francisco, USA, 1990, p. 297. G. & Goringe, M. J., Transmission Electron 17. Thomas, Microscopy qf Muteriuls, Wiley, New York, 1979, p. 157. 18. Kong, C. H. & Munroe, P. R., Processing, Properties and Application of Iron Aluminides, eds M. A. Crimp & J. H. Schneibel. TM’S, Warrendale, USA, 1994, p. 23 1. 19. Jung, H. C., Park, C. R., Lee, C. S. & Park, C. G., Processing, Properties and Application of Iron Aluminides, eds M. A. Crimp & J. H. Schneibel. TMS, Warrendale, USA, 1994, p. 171. 20. Morris, D. G., Dadras, M. M. & Morris, M. A., Acta Metall. et Mater., 41 (1993) 97. 21. Kong, C. H. & Munroe, P. R., Scriptu Metall. et Mater., 32 (1995) 469. 22. Schneibel, J. H., George, E. P., Specht, E. D. & Horton, .I. A., High Temperature Ordered Intermetallic Allo~~s 6, MRS Symp., in press. 23. Schneibel, J. H, private communication, 1995.
Vacancy
hardening
in near stoichiometric
24. Schneibel, J. H. & Maziasz, P. J., Processing, Properties und Application of Iron Aluminides. eds M. A. Crimp & J. H. Schneibel. TMS, Warrendale, USA, 1994, p. 183. 25. Kao, C. R., Pike, L. M., Chen. S.-L. & Chang, Y. A., InterrnetuNics, 2 (1994) 235. 26. Munroe, P. R. & Baker, I., Scriptu MetaU. et Muter., 24 (1990) 2273. 27. Feng C. R. & Sadananda, K., Scripta Met&. et Mater., 24 (1990) 2107. 28. Baker, I. & Munroe, P. R., High Temperature Alurninides and Intermetullics. eds S. H. Whang, et al., TMS, Warrendale PA, 1990 p. 425.
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29. Zaluzec, N. & Fraser, H. L., Scripta Metall., 8 (1974) 1049. 30. Lloyd, C. H. & Loretto, M. H., Phys. Status Solidi, 39 (1970) 163. 3 1. Lasalmonie, A., .I. Mater. Sci., 17 (1982) 2419. 32. Munroe, P. R. & Baker, I., .J. Mater. Sci., 24 (1989) 4246. 33. Baker, I. & Horton, J. A., Philos. Msg. A, 67 (1993) 479. 34. Petzow, G. & Effenberg, G., Ternary Alloys: a Cornprehensive Conlpendium of Evaluated Constitutional Data and Phase Diagrams, Vol. 5, VCH Publishers, New York, 1991, pp. 324-43.
0966-9795(95)00057-7 ELSEVIER
The effect of ternary additions on vacancy hardening in near stoichiometric FeAl P. R. Munroe & C. H. Kong School of Materials
Science and Engineering, (Received
University of New South Wales, Sydney,
21 September
1995; accepted
15 November
NS W 2052, Australia
1995)
The effect of transition metal ternary additions which replace iron on the variations in hardness and defect morphology with heat treatment of near stoichiometric FeAl has been studied. It was generally found that ternary additions of 1% have little effect on the vacancy-related hardening and softening which occurs in binary FeAl. The notable exception to this behaviour is the addition of nickel, which strongly hardens this alloy and dramatically changes the defect microstructure that develops during low temperature annealing. The hardening effect of nickel additions may be attributed to either this changed defect microstructure or the interaction of nickel atoms with the cores of glide dislocations. In contrast, ternary additions of 5% significantly harden FeAl, and it is believed that this is associated with the interaction of ternary atoms with thermal vacancies, in addition to conventional solute hardening. These ternary additions also strongly affected the defect microstructures which develop in FeAl during low temperature, vacancy-relieving heat treatments. Ternary additions resulted in the
1 INTRODUCTION
of
The B2-structured compound FeAl is seen as a candidate material for elevated temperature strucespecially in hostile environtural applications, ments, because of its good oxidation resistance, low density and relatively low cost. Although FeAl is ductile at high temperatures, its commercial application has been inhibited because of its low room temperature ductility and limited strength at However, FeAl is temperatures above -500°C.’ stable over a wide range of compositions,2 which provides an opportunity to improve the mechanical properties of this compound through alloying. Following heat treatments at temperatures greater than about 700°C FeAl retains large concentrations of vacancies, even when furnacecooled, which leads to increases in hardness and decreases in ductility.” 7 This anomalous hardening can be relieved by low temperature (400°C) heat treatments during which excess vacancies are removed from the lattice and high concentrations 403
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C. H. Kong
relieving heat treatment. Their behaviour was monitored by microhardness measurements and their microstructures were examined by both optical microscopy and transmission electron micro- scopy (TEM).
q
Air-cooled
Mn
Co
4OOW 120h
2 EXPERIMENTAL To investigate the effects of transition metal additions on the vacancy retention behaviour of FeAl, two sets of alloys with nominal compositions of Fe,9A1,,X, (X = Cu, Ni, Co, Mn, Cr, V and Ti) and Fe,,Al,,X, (X = Cu, Ni, Mn, Co and Cr) were prepared. All of the samples were annealed in air at 950°C for 2 h and air-cooled. Subsequent low temperature annealing was performed on the air-cooled alloys at 400°C for 120 h. Air-cooled, rather than the water-quenched samples, were chosen for study as they were slightly less brittle and easier to handle than water-quenched samples. Preliminary experiments have confirmed that the behaviour and microstructures of the waterquenched and air-cooled samples were very similar.14 Vickers microhardness testing was performed on each sample using a 300 g load. Each data point obtained was the average of at least 10 measurements. Standard deviations were less than 5’%, but error bars were omitted from Figs 1 and 2 to improve clarity. The microstructures of all the materials were examined by both optical and transmission electron microscopy (TEM). Thin foils for TEM study were prepared using methods described elsewhere,15 and examined in a JEOL 2000FX transmission electron microscope, furnished with an energy dispersive X-ray spectrometer (EDS), operating at 200 kV.
q
4OOW120h
Air-cooled
550 $500 3
450
@-JO 350 300 250 Ti
V
Cr
Mn
Co
Ni
Cu
FeAFe-49AI
Fig. 1. The microhardness of Fe,,AI,,X, (X = Cu, Ni, Co, Mn, Cr, V and Ti) following air-cooling from 950°C and subsequent heat treatment at 400°C for 120 h, data for nominally stoichiometric FeAl and Fe5,A1,, is also included for comparison.
Cr
Ni
0.1
FeAl
Fig. 2. The microhardness of Fe,,Al,,X, (X = Cu, Ni, Co, Mn and Cr) following air-cooling from 950°C and subsequent heat treatment at 400°C for 120 h, data for nominally stoichiometric FeAl is also included for comparison.
The atom site location of the ternary elements was determined by atom sitelocation by channelling enhanced microanalysis (ALCHEMI) techniques; experimental details have been given elsewhere.‘(j
3 RESULTS Figure 1 shows the comparison of microhardness of the ternary alloys Fe,9Al,,X, (X = Cu, Ni, Co, Mn, Cr, V and Ti) with that of both nominally stoichiometric FeAl and Fe-49Al following aircooling from 950°C and subsequent low temperature annealing at 400°C for 120 h. The air-cooled samples (all the ternary alloys and both FeAl alloys) exhibited generally similar hardness values. Following subsequent annealing at 400°C for 120 h the alloys containing Cu, Co, Mn, Cr and V additions again exhibited hardness levels similar to both FeAl and Fe-49A1, although all these ternary alloys were slightly harder than the binary alloys. In contrast, the alloys with additions of either Ni or Ti exhibited much less softening when heat treated at 400°C indicating that the equilibrium hardness values of these ternary alloys were much higher than that of binary FeAl. Figure 2 shows the comparison of microhardness of the ternary alloys Fe,,Al,,X, (X = Cu, Ni, Mn, Co and Cr) with that of nominally stoichiometric FeAl following both air-cooling from 950°C and subsequent annealing at 400°C for 120 h. Following air-cooling from 950°C the ternary alloys containing Cu, Co, Mn and Cr additions showed a comparable or even slightly lower hardness than the binary FeAl, whereas the Ni-containing alloy was harder. Following the vacancy
Vacancy hardening in near stoichiometric FeAl
relieving heat treatment, a small reduction in hardness was observed for all the ternary alloys. In comparison, the reduction in hardness for FeAl was substantial. Thus, these ternary additions not only substantially increased the equilibrium hardness of binary FeAl, but also appeared to reduce the extent of any vacancy-related softening. Optical microscopy showed that following annealing at 950°C and air-cooling, all the ternary alloys and the binary FeAl alloys appeared to be single phase with a comparable grain size of about 200 pm. TEM investigations confirmed that all the alloys exhibited a similar microstructure; single phase with a low density (-lOI m-“) of residual dislocations with cOOI> Burgers vectors. These dislocations are presumably induced through thermal stresses associated with the melting process and subsequent high temperature annealing. The atom site location of the ternary elements added to FeAl was assessed using ALCHEMI methods on the samples air-cooled from 950°C. The results for the Fed9Al,,X, alloys are shown in Fig. 3. Each data point is the average of five measurements. For Cu, Ni, Co, Mn and Ti it was found that more than 95% ternary of these atoms were located on the iron sublattice, although no corrections were made for delocalisation effects. In contrast, V atoms were found to preferentially occupy the Al sublattice, and Cr atoms occupy both sublattices in approximately equal proportions. Similar lattice site preferences were observed for the Fe,SAl,,X, alloys. Following annealing at 400°C for 120 h, TEM studies on the Fe,,Al,,X, alloys revealed that all the ternary alloys, except the titanium-containing alloy, were still single phase, but contained a much higher density (-5 X lOi mm’) of dislocations compared to that following high temperature .‘.’ :: ::
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heat treatment. For example, Fig. 4 shows the microstructure of Fe,Al,,Cr, following annealing at 950°C and air-cooling, and also following subsequent annealing at 400°C for 120 h. In this alloy, following both heat treatments, the dislocations were determined to have
::::
::::
::::
405
::::
q
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Fig. 3. Atom site location of ternary additions in Fe,AI,,X, (X = Cu, Ni, Co, Mn, Cr, V and Ti) following air-cooling from 950°C as determined by ALCHEMI methods.
Fig. 4. Bright field transmission electron Fe,Al,,Cr, following annealing at 950°C for cooling, (b) subsequent annealing at 400°C both micrographs the electron beam direction and diffraction vector is [ 1lo].
micrographs of 2 h and (a) airfor 120 h. For is close to [OOl]
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P. R. Muwoe,
C. ff. Kong
Fig. 5. Bright field transmission electron micrographs of Fe,,Al,,V, following ann_ealing at 4OO’C for 120 h. The electron-beam direction is close to (a-d) [OOl] and (e, Q [ill], and diffraction vectors are: (a) [l IO]; (b) [l lo]; (c) [200]; (d) [020]; (e) [loll and
[200] and [lOi], indicating that their Burgers vector was [OlO]. However, dislocations, such as the one marked _b, exhibited invisibility when g was [I lO]_and [Ol 11, indicating that its Burgers vector was [l Ill. One of these ~111~ dislocations (arrowed in Fig. 5(a)) exhibited a M-shape morphology where it appeared to be pinned by a [OOl]dislocation marked d. The dislocation d which also exhibited an abrupt change in line direction from [OIO] to [IOO] where it had interacted with this
~11 I> dislocation. In addition, some of the ~001~ dislocations exhibited a C-shape morphology, for example those labelled c in Fig. 5(d), inferring that they had experienced some form of pinning force during the heat treatment. Furthermore, in Fe,,Al,,Vr some of the ~11 l> dislocations were observed to exist as dislocation dipoles. Dislocation pairs marked d in Fig. 6 exhibited a change in spacing between the two dislocations when they were imaged using +g and -g under two beam
Vacancy hardening in near stoichiometric FeAl
Fig. 6. Bright field transmission electron micrographs of ~11 I> dislocation dipoles, labelled d, in Fe,,A&,,V, following ianneal ling at 400°C for 120 h. The electron beam direction is close to -.10011 for all three micrographs and the diffraction vectors are (a) [21001, (b) [200] and (4 w201. 1
diffraction conditions, contrast characteristic of dislocation dipoles. l7 It can also be seen in Fig. 6(c) that one of the dipoles has been pulled out from a [OlO] edge dislocation (arrowed) again suggesting that the < 111> defects were formed through the interaction of
pinned at point J. This arrangement again implies that the formation of such < 11l> dislocations is associated with the interaction of
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Fig. 7. Bright field transmission electron micrographs of ~11 I> dislocations, labelled a, in Fe,,Al,,Co, following annealing at 400°C for 120 h. The electron beam direction is close to (a, b) [OOl]_and (c) [ill]. The diffraction vectors are (a) [l lo], (b) [IlO] and (c) LO111.
Following air-cooling from 950°C and subsequent annealing at 400°C for 120 h, TEM observations revealed that all the Fe,SAl,OX, alloys except the chromium-containing alloy, remained single phase, but all again contained a significantly higher density (-5 X 1013 me2) of dislocations
Fig. 8. Bright field transmission
electron micrograph showing the intersection (arrowed) of a < 11 l> dislocation, labelled V, and a
compared with that observed following air-cooling from 950°C. Figure 9(a) shows the microstructure of Fe,,Al,,Co, following air-cooling from 95O”C, it is clear that this alloy contains a low dislocation density. Figures 9(bd) show the microstructures Fe,,Al,,Cu,, FeJSAISOMnSand Fe4,A1,,Co, respectively following subsequent annealing at 400°C. In Fe,SA1,,Co, (Fig. 9(b)) the dislocations, all of which had
Vacancy hardening in near stoichiometric FeAl
Fig. 9. Bright field transmission
electron micrograph of the microstructure of: (a) Fe,,Al,,Co, following air-cooling from 95O’C; (b) Fe,5Al,,CuS; (c) Fe,,Al,OMn, and (d) Fe,5A1,,Co, following subsequent annealing at 400°C for 120 h. For all the micrographs the electron beam direction is close to [OOl] and diffraction vector is [ 1lo].
g was both [OZO]and [Oli], so the Burgers vector was [loo], and dislocations, such as the one labelled b exhibited invisibility when g was both -[I lo] and [lOi], so the Burgers vector was [lil]. It was also noted that a weak fringe contrast was observed to be associated with these dislocations. This suggests that these defects have the Burgers vector l/2< 11l> and their motion has created disorder and left a region of matrix which is out of phase with the rest of the matrix. In addition, some of the ~1113 dislocations, for example those labelled c, which exhibited contrast consistent with a [ll l] Burgers vector were in the form of tangled dislocation loops. However, the dislocation labelled d exhibited invisibility when g was both [OZO]and [lOi], indicating that its Burgers vector was [loll. This defect is connected to two other dislocations, each of which exhibited contrast consistent with a ~11 l> Burgers vectors, suggesting that it had formed through the interaction of two dislocations. The microstructure observed in Fe~SA1,,Cr, following air-cooling from 950°C and subsequent annealing at 400°C for 120 h was much more
complex than the other Fe,Al,&, alloys studied. Bands of slightly elongated, plate-like second phase precipitates (arrowed), about 50 nm in diameter, were observed, which trace analysis showed were parallel to (112) planes (Fig. 11). If this region was imaged using (001) superlattice reflections, lines parallel to the { 112) planes were observed which exhibited weak contrast relative to the blocks of matrix enclosed between them (Fig. 12(a)). Such images are strongly reminiscent of antiphase boundary (APB) related contrast in disorderable Fe3Al-based alloys following deformation.‘9v20It appeared that the apparently disordered lines were connected, or pinned, to the second phase particles, and their line directions were, in general, parallel to the longitudinal axis of the precipitates. If these regions were imaged using any fundamental reflections, such as f 1101, these lines became invisible. This is consistent with the suggestion that this contrast was associated with APBs. To study the precipitates and APB-type contrast observed in Fed5Als0Cr5in more detail selected area electron diffraction (SAD) patterns were taken
410
P. R. Munroe,
C. H. Kong
Fig. 10. Bright field transmission electron micrographs of Fe,,A1&05 following annealing at 400°C for 18-h. The electron beam directions are close to (aa) [OOl] and (e, f) [ill] and the diffraction_vectors are: (a) [200]; (b) [llO]; (c) [l lo]; (d) [020]; (e) [loll and (f) [Ol I].
from a number of low index axes of the matrix. Figure 12 shows the SAD patterns recorded from the [l lo] and [II l] zone axes. Some weak extra spots, marked with a subscript p, together with the expected FeAl matrix reflections, marked with a subscript m, were observed. Analysis of the diffraction patterns showed that the precipitate spots were consistent with the second phase being were observed in FeAl,. Similar precipitates Fe,5A1,,V,, and the analysis of these patterns has
been described in detail elsewhere.21 The orientation relationship between the FeAI, precipitates and FeAl-based matrix in Fe,,Al,,Cr, was also the same as that observed in Fe45Al,0V5.21That is: [10I JFeAl, // [ 11 l]FeAl and (040)FeAl, // (0 1i)FeAI. It is also worth noting that the precipitate reflections arising from the second phase in the SAD pattern taken from the [l 1 I] zone axis of the matrix were elongated in the direction of the ( 112) planes
Vacancy hardening in near stoichiometric FeAl
Fig. 11. Bright
field transmission electron micrograph of Fe,,Al,,Cr, following annealing at 400°C for 120 h. The electron beam direction is close to [ 10 l] and the diffraction vector is [020].
Fig. 13. Selected area electron diffraction patterns from matrix phases in Fe,,Al&r, following air-cooling from 950°C and subsequent annealing at 400°C for 120 h. Zone axes of the FeAl-based matrix are (a) [l IO] and (b) [I 111.
4 DISCUSSION
Fig. 12. Transmission electron micrographs of Fe,,A1,,Cr5 following annealing at 400°C for 120 h imaged using (a) a (010) superlattice reflection and (b) a (101) fundamental reflection.
of the matrix. These elongated diffraction spots may result from the precipitates exhibiting a small spread in orientation relationship with the matrix. Such elongation was not noted in the diffraction patterns obtained from Fe4SA150V5.21
Following heat treatment at 950°C and air-cooling all the Fed9A15,,X, alloys and Fe,SAl,,X, alloys exhibited approximately equal hardness values, which were comparable to that of binary FeAl. Following such heat treatments it has been shown that a large concentration of vacancies is retained in the lattice in FeAl and this strongly increases hardness.6,7 The similar hardness of these ternary alloys to the binary alloy following this heat treatment suggests that they also retain significant concentrations of vacancies, which increase hardness. Furthermore, these data would suggest that any solute hardening effects associated with the ternary additions are largely obscured by the presence of these thermal vacancies. However, subsequent annealing of these alloys at 400°C revealed larger differences in hardness behaviour between the ternary alloys and FeAl. For those alloys with the generic composition Fe,Al,,X, considerable softening occurred during
412
P. R. Munroe, C. H. Kong
this heat treatment, except for Fed9Al,,Ni, and Feb9Al,,Ti,. That is, the hardness of these alloys was comparable to, but marginally harder than, binary FeAl. However, it is worth noting that the marginally higher hardness values obtained after this heat treatment increased with increasing difference in atomic number between iron and the ternary addition. This hardening behaviour is presumably related to solid solution hardening associated with the ternary additions. The increasing hardness with increasing difference in atomic number between iron and the ternary addition is presumably associated with atomic size misfit as suggested by Schneibel et ~1.~~The exception to this behaviour was Fe,Al,Ni,, which exhibited a much higher hardness than either the copper and cobalt containing alloys, suggesting that a more complex hardening behaviour was in operation in this alloy. This behaviour may be related to the defect structure evolved during heat treatment at 400°C. Voids, in addition to dislocations, were observed in Fe,,Al,,Ni, following this heat treatment, and it is possible that these voids induce additional hardening in this alloy.‘*,13 In addition, it has been suggested that the nickel atoms in this alloy induce core changes around glide dislocations which inhibit their motion and so harden the alloy. 23A very high hardness was also observed in the Fe49Al,0Ti, following the 400°C anneal, although some of this hardening may be attributable to the atomic size misfit effect, this alloy contained a large number of fine second phase particles which may also have contributed a hardening effect. In contrast, the Fe,,Al,,X, alloys softened to a much lower degree during annealing at 400°C compared with FeAl. That is, their equilibrium hardness values at this temperature were much higher than that of the binary composition. It is possible that the higher concentration of ternary additions in these alloys leads to greater vacancysolute interactions, which thus inhibit the removal of vacancies from the matrix during low temperature annealing. That is, the high hardnesses of these alloys following annealing at 400°C are still attributable to the presence of thermal vacancies trapped in the lattice. Alternatively, the higher hardness of these alloys may be associated with solid solution hardening effects. Consistent with this was the observation that the hardness levels increased with increasing atomic number difference between iron and the ternary addition, with the exception of the Nicontaining alloy, which was again much harder
than any of the other single-phase alloys. The high hardness of the Cr-containing alloy would presumably be due in part to the high density of fine particles which precipitated following the low temperature heat treatment. Schneibel et al. found similar trends in the yield strengths of Fe45Al-5X alloys (where X = Cu, Ni, Co, Fe, Mn, Cr, V, Ti) with small B and Zr additions.22,24 In that case, a linear relationship between the yield strength and the atomic size misfit for all ternary alloys was established with the exception of nickel-containing alloy, which again showed a much stronger strengthening effect. However, these workers noted that the manganese alloy was marginally softer than the binary alloy. The results of ALCHEMI studies revealed that Cu, Ni, Co, Mn and Ti atoms preferentially occupied the iron sublattice, whereas V atoms preferentially occupied the aluminium sublattice and the Cr atoms occupied both sublattice sites in approximately equal proportions. Recently Kao et al.” predicted the atom site preference of ternary additions of transition metals in a range of B2-structured intermetallic compounds, including FeAl, based on the calculation of standard enthalpies for binary compound formation. For FeAl it was predicted that ternary elements with atomic numbers greater than Fe, such as Co, Ni and Cu, would prefer to occupy the iron sublattice, and ternary elements with atomic numbers less than Fe, such as Cr, Ti, V and Mn would prefer to occupy the iron sublattice. However, these predictions for Cu, Mn, Cr and Ti were limited by being subject to significant statistical error. Schneibel and co-workers used ALCHEMI to determine the lattice site occupancy of ternary additions in a range of Fee45Al-5X alloys. 22 Their results were broadly in agreement with those obtained here, except that both titanium and chromium were observed to exhibit a strong preference for the aluminium sublattice. Earlier work by Munroe and Baker also showed that chromium occupied the aluminium sublattice in alloys based on Fe-40Al.26 The differences in behaviour for these results with the observations made here may arise because the alloys previously examined were iron-rich, and thus in these alloys the location of the ternary additions on the aluminium sublattice would be expected to reduce the number of iron antisite defects. In the present study, however, the location of the titanium atoms on the iron sublattice would minimise the number of anti-site defects. The even distribution of chromium atoms between both iron and aluminium sublattices in these alloys, would lead to
Vacancy hardening in near stoichiometric FeAl no iron anti-site defects, but the presence of the chromium atoms on the aluminium sublattice would produce constitutional vacancies. Presumably, for Fe,,Al,,X, it is energetically more favourable to contain a small number of constitutional vacancies than to occupy the iron sublattice in significant numbers. The microstructures of all the alloys following air-cooling from 950°C were observed to contain very low density of
413
It has been shown that many of these ~1 1l> dislocations were in the form of dipoles. ~1112 dislocation dipoles have been observed elsewhere in extruded iron-rich FeA1.27,32 In these cases, it was assumed that their formation was associated with dislocation interactions during high temperature deformation. Recently, Baker and Horton observed in stoichiometric FeAl following room temperature deformation dislocations with < Ill> Burgers vectors which were in the form of either dipoles or exhibited C-shape morphologies, similar to that observed in this study.33 It was suggested that these dipoles and C-shape dislocations arose not through mutual dislocation interaction but through interaction of gliding dislocations with quenched-in vacancies. It is likely that the dipoles formed here are also associated with presence of vacancies. The M-shape dislocation in Figure 5 appears to be formed by the pinning action of a ~11 l> dipole with a
414
P. R. Munroe,
was observed. Neither these precipitates nor the APB-type contrast were observed in the samples air-cooled from high temperature. This indicates that they were formed during subsequent low temperature annealing. This alloy is significantly harder than near-stoichiometric binary FeAl, following the same heat treatment, suggesting that in addition to solid solution effects these precipitates may provide some contribution to hardness. It is also clear that the solubility of chromium in nearstoichiometric FeAl at 400°C is greater than lo/o but less than 5%, but at higher temperatures the solubility of chromium in FeAl is greater than 5%. Examination of the Fe-Al-Cr alloy phase diagram 873 K isotherm (the lowest temperature available) shows that Fe,,Al,,Cr, lies in the (FeAl) B2 plus disordered (a-Fe, Cr) phase field.34 This is somewhat contrary to the observed microstructure, where FeAl, particles were also observed. The APB-like contrast is therefore related to the formation of the disordered (a-Fe, Cr) phase. It is, however, interesting to note that these disordered regions are associated with the FeAl, particles. This suggests that localised chemical variations occur in these regions, and the local decrease in aluminium content in the disordered regions, leads to adjacent regions locally rich in aluminium, and so the formation of FeAl,. However, these areas were sufficiently fine to preclude detailed microanalytical studies. Further studies, using high resolution microanalytical techniques are underway to investigate the development of the microstructure in these FeAl+Cr alloys.
5 CONCLUSIONS
(1) The
variation in hardness with heat treatment for alloys with the composition Fe49Al,,X, is little affected by ternary additions. The exceptions to this behaviour are those alloys where either titanium or nickel are the ternary additions. Significant hardening in Fe,,Al,,Ti, is associated with presence of second phase particles. In contrast, the hardening in FeJyA1,,Ni, is believed to be associated with the interaction of nickel additions with the cores of glide dislocations and the significant effects which nickel has on the defect microstructure which evolves in FeAl-based alloys during low temperature heat treatment. (2) Ternary additions are highly effective in hardening alloys with the composition Fe,,Al,,X,. This is believed to be associated with the inter-
C. H. Kong
action of vacancies with the solute atoms as well as solid solution hardening effects. Ternary additions affect the dislocation structures (3) which evolve in FeAl during low temperature vacancy relieving heat treatment. Vanadium and cobalt additions lead to the formation of < 11 1> dislocations in addition to
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