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The influence of microstructure on the ductility of iron aluminides
Intermetallics 7 (1999) 1121±1129
The in¯uence of microstructure on the ductility of iron aluminides D.G. Morris*, M.A. Morris-MunÄoz Department of Physical Metallurgy, CENIM, CSIC, Avenida Gregorio del Amo 8, 28040 Madrid, Spain Dedicated to Dr. Gerhard Sautho on the occasion of his 60th birthday Received 1 March 1999; accepted 11 March 1999
Abstract Considerable eort has been devoted over the last decade to the development of iron aluminides as materials for high temperature applications, where their good oxidation and corrosion resistance, combined with reasonable strength, may be utilised. Poor formability and ductility, however, particularly at room temperature, has hampered the exploitation of these materials. The present review examines the present state of understanding of the factors which in¯uence the ductility. Recent research has made clear the important in¯uence of testing environment, the role of Al content and minor additions of B, as well as the eect of quenched-in vacancies. The extent to which other factors, such as alloying additions and microstructural features, aect the ductility has not received the same attention, and is examined in the present study. Alloy strengthening, by almost any mechanism, is seen to lead to a dramatic loss of ductility. The only parameter allowing both strength increase and ductility improvement for a given set of Al/B/ vacancy/environment conditions is the grain size. The best ductility for a given alloy, which should have as low an Al content as compatible with other requirements, is obtained by re®ning the grain size and by maintaining the alloy in the softest possible state. For the most part these conclusions are drawn from analysis of the behaviour of B2 ordered FeAl alloys, although similar trends seem also to apply to alloys of slightly lower Al content where DO3 ordering can occur. The observations drawn can be understood in terms of the mechanisms leading to the nucleation and propagation of brittle fracture, either as transgranular cleavage cracks or as grain boundary cracks. The possible role of additional factors, such as the texture, or grain and grain boundary distribution, surface layers producing protective stress eects, and strain homogenising or crack arresting dispersions, has not been suciently evaluated to determine whether any further improvements of ductility are possible. # 1999 Elsevier Science Ltd. All rights reserved. Keywords: A. Iron aluminides; D. Microstructure
1. Introduction Iron aluminides fail during tensile testing after limited plastic deformation associated with uniform gauge length elongation, and considerable work hardening, before ®nal fairly brittle failure with little local necking. As the Al content of iron aluminides increases from the Fe3Al composition to the FeAl composition there is a gradual decrease in ductility and a trend away from the ductile failure typical of Fe base alloys to cleavage failure and eventually to intergranular failure near stoichiometric FeAl. Fig. 1 and Table 1 [1±8] show the gradual fall in ductility as the Al content increases. For a long time the scatter in experimental data on fracture ductility of such iron aluminides has been so large that it has been dicult to deduce trends and explanations of the cause of ductility variations. It is only relatively recently that the important role played by such factors as environmental * Corresponding author.
embrittlement [9±14], and hence Al content and strain rate [15], B segregating to grain boundaries and modifying grain boundary behaviour [5,15±17], and hardening by quenched-in vacancies [5,7], have become clear. Since many of the mechanical testing experiments carried out earlier lent scant attention to some or all of these factors, the literature shows very wide scatter of fracture data both as a function of speci®c test parameters as well as between investigators and reports. Fig. 1 and Table 1 make clear the dramatic loss of ductility as the Al content increases from near 40 to 50% in FeAl alloys, when testing is carried out in air instead of under vacuum or in dry oxygen, and also in the absence of B from the alloy. Note that throughout this text all compositions are given as atomic percentages. The dramatic eect due to the choice of environment where testing is carried out [9±14] has been explained by the reaction of freshly-exposed intermetallic at a propagating crack tip with water vapour, whereby atomic hydrogen is injected into the material in
0966-9795/99/$ - see front matter # 1999 Elsevier Science Ltd. All rights reserved. PII: S0966-9795(99)00038-2
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Fig 1. Tensile ductility of well-annealed FeAl in air and in vacuum as a function of Al content. This ®gure is taken from the review of Table 1 In¯uence of Al content , B additions and testing atmosphere on tensile ductility of FaAla Aluminium content (%) Atmosphere/B addition?
35
40
45
50
Air No B
TG 2%
Mixed 1%/3%
IG 3%/1%
IG 0%
Vac/oxygen No B
IG 5.5/18%
IG 5%
±
±
Air B Oxygen B
±
TG 4% IG 17%
TG 2% ±
±
±
±
a TG, IG and mixed refer to transgranular, intergranular and mixed (transgranular and intergranular) failure modes. All alloys examined are expected to have B2 order as it is unlikely that any DO3 order appears for Al contents above 35%. The data shown are approximate average behaviours reported by a number of workers, including [1,4,5,10,11,16].
front of the tip leading to easier failure [10]. Since such embrittlement aects cleavage planes more than grain boundaries, an additional eect of the environmental attack is the tendency to cleavage failure rather than intergranular failure [10] (Table 1). It may also be noted that since the environmental attack takes place especially on cleavage planes it is a phenomenon found also for single crystals [13]. In Table 1 a gradual reduction of ductility is evident as the Al content increases. Increases in the Al content, as well as testing in vacuum/oxygen, as well as the absence of B, all tend to lead to intergranular failure [1,4,5,10,11,16,17]. These trends are due to the intrinsic weakness of grain boundaries at high Al levels, the preference for environmental attack on cleavage planes, and the eect of B on improving grain boundary adhesion. It should be noted that over the range of alloys and test conditions discussed here the general mechanical
properties of FeAl alloys are only slightly modi®ed (for example by Al content over the range of say 35 to above 45% and even as high as 48%) and hence the various test parameters truly aect fracture behaviour, not general plasticity. For example, the hardness of wellannealed alloys over the range 35% to above 45% Al [18], the yield stress of alloys over the range 40±48% Al [1] and also the work hardening rate over the same 40± 48% Al range [1], as well as the Hall±Petch slope over the range from below 35% Al to above 45% Al [19], are essentially constant, such that any change of ductility recorded (Table 1, Fig. 1) is indeed due to changes of fracture behaviour. Deformation mechanisms in iron aluminides at temperatures near ambient temperatures involve the motion of superdislocations of <111> Burgers vector generally gliding on {110} planes with perhaps some cross slip onto {112} planes. This aspect of deformation has been examined on many occasions, and further detailed references may be found in the review by Baker and Munroe [8]. The factors discussed above concern in¯uences on the propagation of an already existing crack, for example the eects of environmental embrittlement. The nucleation of such a crack presumably occurs by the interaction of various dislocations or slip systems. A possible nucleation mechanism has been proposed by Munroe and Baker [20] involving the reaction of <111> superdislocations of dierent Burgers vectors on intersecting glide planes to produce what are eectively small pile-ups of edge <100> dislocations. If the edge dislocations produced are suciently sessile they could represent the initial stage of formation of a cleavage crack on a {100} plane. This plane has indeed been reported to be the cleavage plane for iron aluminides [7,21]. Any alloying or microstructural factor which in¯uences the tendency for <111> superdislocations to come together to produce such sessile <100> edge dislocation pile-ups will in¯uence the tendency to crack initiation. Coarser microstructures, such as larger grain sizes, would allow longer slip bands of the initial <111> superdislocations to be created with correspondingly higher stresses available to create larger numbers of <100> edge dislocations. It has also been suggested that the presence of hydrogen may change the mobility of <111> and <100> dislocations [22] or the elastic energy of <111> or <100> dislocations [23]; less mobile <100> dislocations of lower elastic energy should be easier to create, allowing easier crack initiation. As a ®nal comment, it is emphasized that the intention of the present short review is to demonstrate in¯uences of alloying and microstructural parameters on fracture behaviour and ductility, and as such the present discussion of the general fracture mechanisms, and in¯uences of testing conditions (including Al and B content) on behaviour is deliberately restricted to being a very brief overview. Excellent, more detailed, reviews
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have been presented elsewhere, e.g. [8,9]. Fundamental understanding of the fracture processes, however, both crack initiation and crack propagation, remain poorly or incompletely understood, even though there is now a much better empirical description of the in¯uences of the various test and alloy parameters. 2. Vacancy hardening, and the in¯uence of annealing treatments on ductility It is now well established that large numbers of vacancies will be present in iron aluminides cooled from high temperatures, for example after casting, recrystallization or a solutionising anneal, and these will only be removed by long anneals at relatively low temperatures [18]. Such vacancies are known to lead to considerable hardening of the materials [18,24], presumably by vacancies acting as solute for pinning dislocations, while vacancy aggregates such as dislocation loops [25] and voids [26] may also cause similar hardening. The role of such point defect hardening on ductility is also now clear, as illustrated in Fig. 2, for a Fe±45Al alloy, both with and without B additions [5,8]. Furnace cooling, which leaves a fairly high vacancy concentration, is seen to lead to fairly low ductility for these alloys, while removing vacancies by a slow cool from high temperature and a long anneal at low temperature leads to progressively improved ductility. Such improvements of ductility by vacancy removal have now been seen on a number of occasions [7,11,27]. The relationship between hardening and embrittlement by vacancies has been evaluated in more detail by Pang and Kumar [27], as summarised in Fig. 3, showing the yield stress after dierent high temperature anneals, quenches, and low temperature ageing, and the related tensile ductility. It should be noted that these results were obtained on a Fe±40Al alloy containing in addition 0.6%
Fig. 2. In¯uence of dierent heat treatments to reduce the retained vacancy concentration on the ductility of Fe±45Al alloys, with or without 400 ppm B [5].
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carbon, and this material contained some carbon in the form of fairly large perovskite (Fe3AlC) carbides, which may themselves play some role in aecting fracture [27], as discussed later. During high temperature annealing and subsequent ageing, some part of the carbon may dissolve and subsequently precipitate [27] and hence annealing eects reported may be compound eects due to both vacancy and to carbon eects. At the same time, the interaction of lattice vacancies and interstitial C atoms is likely to modify the mobility and tendency to agglomeration of vacancy defects [27]. The ®lled symbol data points of Fig. 3 show the eect of ageing at about 400 C, starting from a high temperature extruded state, on ductility and yield stress: the eect of such ageing is a steady reduction of stress (from about 500 MPa to near 300 MPa) and a concomitant increase of ductility (from near 2% to above 4%). The open symbol data points in Fig. 3 show the results of several experiments on quenched material which is subsequently aged at dierent temperatures, with mechanical testing carried out at various times during ageing. These data show, interestingly, that there is some slight age hardening, probably associated with the carbon precipitation, and at the same time a notable embrittlement. One such age hardening-embrittlement cycle is indicated by the sequence of arrows in Fig. 3. Subsequent annealing to soften leads to an increase in ductility again. Despite some scatter in the ductility-stress data of Fig. 3 about the best ®t line, it is clear for this material that ductility and yield stress are essentially reciprocally related by a single master curve. During ageing of such quenched materials it is well known that the vacancies
Fig. 3. In¯uence of heat treatments on yield stress and tensile ductility of Fe±40Al containing 0.6% carbon [27]. The ®lled symbol data points show the softening and ductilising eect of ageing material extruded (at 1000 C) at low temperature (400±500 C). The open symbol data points show ageing data for material well quenched from high temperature (1000 C), when the grain size has increased somewhat. The sequence of arrows shows an example of the variation of yield stress and ductility on progressively ageing a quenched material at one ageing temperature (500 C). Data for other ageing temperatures (400 C for strong±brittle materials and 650 C for softer±ductile materials) show similar stress±ductility trends.
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initially present agglomerate, generally as dislocation loops of <111> Burgers vector, which subsequently grow, interact, and ®nish as a low density of <100> dislocations, [25]. The implication of the master curve representation of Fig. 3 is that the precise mode of strengthening, by a solution of vacancies or by a population of <111> or <100> dislocations, is immaterial. The only relevant parameter determining ductility is the yield stress, independent of the precise mechanism causing hardening. The earlier section described failure in terms of mechanisms of crack formation and growth. It seems clear from the present data that it is the absolute level of stress alone that determines the ease of crack nucleation and propagation, with little to no in¯uence of the precise mechanism that causes the hardening. 3. In¯uence of grain size on ductility In general terms it is well established that grain size re®nement is a potent way to improve ductility or to bring about a brittle to ductile transition [28]. These eects are deduced from the grain size dependence of the yield stress, for the initiation of plastic ¯ow, as described by the Hall±Petch relationship (y i ky dÿ1=2 ) and from the similar dependence of the fracture stress, characterised by a smaller (or zero) value of internal stress (i ) and a higher value of the grain size dependency slope kf . While some researchers have reported an increase in ductility as the grain size is reduced [7] other times it is reported that grain size re®nement has only a minor eect on ductility [8,17]. For example, Cohron et al. [17] compared the ductilities of Fe±45Al with 300 ppm B with grain sizes of 22 mm and 120 mm to ®nd almost identical ductility in the two cases. Baker and Munroe [8] also concluded that the slight improvements in ductility reported by grain re®nement were generally overshadowed or confused with other factors aecting ductility, for example due to particles introduced to ensure grain re®nement or due to environmental embrittlement eect, and it is perhaps only in those cases where other factors aecting ductility are very well controlled, or their eects are avoided, that grain size in¯uences can be detected. Fig. 4 shows variations of tensile ductility of Fe-40Al as the grain size varies over a very wide range Ð from single crystal down to grain sizes of about 5 mm Ð and shows a clear increase in ductility, particularly for the very ®ne grain sizes [7,21,29,30]. It should be emphasized, however, that the ®nest grain sizes are obtained using powder metallurgy/rapid solidi®cation techniques, and that dispersoid particles are often present for grain boundary pinning, but these may also serve to disperse slip and delay crack nucleation. Over a narrow grain size range, however, and particularly for the fairly large
Fig. 4. Eect of grain size reduction, from single crystal to about 5 mm, on tensile ductility of Fe±40Al [7,21,29,30] (®lled symbol data points) and Fe±40Al±0.6C [27] (open symbol data points).
grain sizes obtained by conventional casting/extrusion methods, the variation of ductility with grain size is very small, as illustrated by the data on Fe-40Al with carbon [27]. An interesting point to be made with regards the data of Fig. 4 is that the re®nement of grain size will be associated with a signi®cant strengthening of the material [7], as indeed expected from the Hall±Petch relationship. As discussed earlier, and seen in Fig. 3, an increase in strength is expected to be associated with a reduction in ductility, at least for a constant grain size. The observed increase in ductility associated with grain size reduction and hardening is indeed proof of the powerful eect of grain size reduction on delaying cracking. Fig. 4 also suggests that the addition of carbon to Fe± 40Al serves to increase the ductility signi®cantly. In this sense the C addition appears to produce similar gains in ductility to that achieved by B additions, Fig. 2 for example. This interpretation should be considered with great care, however, because the carbon-containing alloy achieved the good ductility shown in Fig. 4 only after careful annealing at low temperature to remove vacancy hardening, and as seen in Fig. 2 such annealing can considerably improve the ductility levels. 4. In¯uence of substitutional alloying on ductility The in¯uence of substitutional alloying on ductility has been examined in great detail for alloys containing about 28% Al [31±34], where DO3 order may be found, with less work reported on the B2 alloys containing 40% Al or more [35]. Eects of such alloying on Fe28Al-based alloys are summarised in Fig. 5, where yield stress and tensile ductility are shown. The present discussion will consider changes of ductility and stress in terms of such alloying, considering that all other parameters are constant, e.g. grain size, test environment,
D.G. Morris, M.A. Morris-MunÄoz / Intermetallics 7 (1999) 1121±1129
Fig. 5. Changes of yield stress and tensile ductility with additions of various alloying elements (Mo, Nb, W) to the Fe±28Al or to the Fe± 28Al±5Cr base [31±34]. The Cr addition to the Fe±28Al base leads to increased strength and ductility. Additions of Mo, Nb, or W to the respective alloy bases lead to increased strength with only small changes of ductility for small alloying additions, and typically further strengthening with some loss of ductility for larger alloying additions.
strain rate, etc. As dierent amounts of various alloying additions were used for the experiments of Fig. 5 the grain size varied from around 50 to 100 mm for the binary alloy, or with Cr addition, to near 30 mm with large additions of Mo, Nb or W. As seen from Fig. 4 such reductions of grain size do not lead to major changes of ductility. The most dramatic eect shown in Fig. 5 is the increase in ductility on adding Cr to the binary Fe±28Al alloy. This increase was ®rst ascribed to changes of plasticity of the material [36], but has subsequently been associated with changes in the nature of surface ®lms and the extent of environmental embrittlement [37,38]. Indeed it is possible that all the changes of ductility in Fig. 5 are related to the environmental embrittlement eect and some quality of the surface ®lms produced before or during testing. It should also be remembered that all the alloys discussed here show DO3 order after annealing at moderate temperatures, and the eects seen may be somewhat aected by slight changes of the degree of this order. Fig. 5 shows that, with the exception of the ductilising eect of Cr, alloying by any addition (of Nb, Mo or W) leads to a progressive loss of ductility. Hardening and embrittlement occur when the alloying additions are dissolved as solute (e.g. 0.5±1% Mo) or as second phase intermetallics or carbide particles (large additions of all the solutes examined). Nevertheless, for the Fe±28Al± 5Cr intermetallic base it is possible to achieve notable strengthening (for example to a yield stress near 600MPa) by Mo and particularly by Nb additions without dramatic loss of ductility. Whilst the ternary (or quaternary) additions to Fe± 28Al(+Cr) can strengthen and lead to only a gradual
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loss of ductility, there is no evidence that alloying with any substitutional element (including Cr) to Fe±45Al produces any bene®cial eects [35]. In this study, even the addition of Cr leads to a fall in ductility, and this for any testing environment. Some other studies have examined alloying additions to FeAl alloys, for example [6], and have typically considered alloying additions that lead to second phase precipitation. Gaydosh et al. [6] considered additions of B and C alone to the Fe± 40Al base, as well as small additions of substitutional elements such as Zr and Hf, and also combined additions of the light and the heavy elements. The only material showing improved room temperature ductility over the binary base alloy was the one containing both Zr and B, with neither Zr alone, C alone, nor C+B alone improving ductility. The improvement of ductility observed in this case may be due to the eect of B in increasing grain boundary strength, although there may also be an eect of particles of Zr intermetallics or borides on increasing ductility Ð by tending to homogenise slip Ð or on decreasing ductility Ð if the coarse particles lead to stress concentrations and hence nucleate cracks. 5. Yield stress and grain size in¯uences on ductility The previous discussions have demonstrated the important role that grain size of a given material, and extent of hardening for a material of given grain size, will play in determining the tensile ductility (Figs. 3 and 4), and this for material tested under reasonably similar conditions of test environment and strain rate, and possessing similar chemical compositions in terms of Al content, B additions, and so on. Fig. 6 attempts to correlate the tensile ductility (in air, typical laboratory strain rates, etc.) of thermally-relaxed Fe40Al (generally containing small amounts of B or C, but no signi®cant amounts of substitutional additions) with the two parameters Ð grain size and yield stress [6,7,27,30,35,39±45]. The causes of strengthening and ductility changes for a given material and grain size are multiple, ranging from the extent to which vacancies are retained or completely annealed out with very low dislocation density, some retention of work hardening from material preparation, small alloying and impurity additions, etc. It should also be remarked that, while the data are collected with approximate grain size groups in Fig. 6, the grain sizes reported in the literature vary steadily from very small to very large without having the precise values suggested in Fig. 6. The lines connecting data of approximately the same grain size in Fig. 6 show that to a reasonable approximation it is indeed possible to correlate ductility with yield stress. Despite the extensive work carried out on iron aluminides over the past decades the available data are still limited and only 5±10 data points are available
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Fig. 6. Variation of tensile ductility in Fe±40Al alloys as a function of the grain size and the yield stress. Data taken from the literature sources [6,7,27,30,35,39±45]. Data are grouped corresponding to approximately constant grain sizes (microns), with strength±ductility data correlated for each approximate grain size. The data points refer generally to single phase or dispersion-strengthened materials prepared by typical casting, thermomechanical processing, or powder metallurgy methods. The open data points refer to the Fe±40Al±0.6C material [27] discussed earlier, where strength and ductility variations are obtained in the same material by heat treatments. Some of the ®lled data points lie o their grain-size trend lines, and are shown attached to this by dotted lines: these data refer to materials of lower Al content (e.g. 36%Al, [45]), when the ductility-strength is above the trend line, or of higher Al content (e.g. 45%Al, [35]), when the ductility±strength is below the trend line, taking account of the in¯uence of Al content on the ductility.
for representation at each grain size. The scatter of the data is generally not excessive bearing in mind the number of experimental variables which change slightly from one literature case to another. The ®lled data points of Fig. 6 refer to samples prepared under similar conditions by the same workers during a single study or by a number of workers during several studies. The open data points, linked by a dotted line, show data on the same Fe±40Al±0.6C alloy [27] as in Fig. 3 and refer to the in¯uence on yield stress and ductility of dierent quench and ageing conditions to modify the number of retained vacancies and vacancy aggregates. The smaller slope of the line connecting these data suggests that vacancy hardening is a powerful way of hardening while retaining relatively high ductility (for the grain size and hardness), or otherwise that the presence of C in solution in the quenched and lightly aged material may be responsible for this high hardening. Grain size reduction by a large amount is con®rmed to be a powerful way of increasing the ductility Ð yield stress combination, with reductions of grain size by a factor of 4±5 leading to the equivalent change over the entire grain size range. For example, grain size reduction from 200±50 mm to 10±2 mm allows increases of yield stress from 300±450 to 600±750 MPa for 5% tensile
ductility or, otherwise, at a yield stress of 500 MPa the ductility increases from 0±4 to 10±20%. Fig. 6 shows also several individual data points connected to their grain-size lines by dotted connecting lines (for example there are two such data points near 12±15% ductility and 650 MPa). These data refer to materials which do not have 40% Al but instead 36% Al or 45% Al [35,45]. The several data points displaced to the right of their grain size lines (for example the two data points described before: ductility 12±15%, stress 650 MPa) correspond to an alloy containing 36% Al [45]. This material is somewhat more ductile (or stronger for a given ductility and grain size) than the corresponding Fe±40Al alloy, but not to an extent that suggests that dierent failure mechanisms operate. Similarly, the one data point displaced to the left of its grain size line (near a ductility of 8% and yield stress of 300 MPa) corresponds to an alloy containing 45% Al [35]. Again this material is more brittle than the Fe-40Al alloy of same grain size and yield stress, but not to an extent incompatible with the present yield stressductility-grain size relationships. 6. Other microstructural parameters which may aect ductility There have been some reports that the presence of dispersoid particles within the iron aluminide matrix can serve to distribute plastic ¯ow, avoiding or at least delaying the formation of intense slip bands and stress concentrations and thereby delaying crack nucleation and failure. Data in Fig. 6 of Maziasz et al. [45] (grain size line of 1±3 mm and ductility of 12±15%), and data of Strothers and Vedula [30] (also grain size line of 1±3 mm and ductility of about 7%) correspond to such cases of dispersed particles throughout the material. The question of importance is whether these materials show their good ductility because of their ®ne grain size and relatively low yield stress Ð i.e. they sit on the ductility± yield stress line that corresponds to the 1±3 mm grain size Ð or whether they show exceptionally good ductility Ð i.e. they sit above the true ductility±yield stress line for their grain size. The data presently available are insucient to provide a clear answer. In another study of Fe±40Al alloys, however, Gaydosh et al. [6] decided that rather large dispersed particles, of the order of several microns in size, could nucleate cracking either at the particle-matrix interface or by particle cracking, and such coarse particles should clearly be avoided. Another interesting question concerns the special role of carbide particles in iron aluminides where, again, the evidence is con¯icting. In a study of the eect of carbon in Fe3Al, Kerr [46] concluded that carbon additions could lead to embrittlement, although it was not completely clear whether carbon in solution, some in¯uence
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at grain boundaries, or Fe3AlC perovskite particles were responsible. On the other hand, Pang and Kumar [27] decided that C additions to Fe±40Al could enhance ductility, as implied in Fig. 4, where the C-containing alloy shows greater ductility than the carbon-free Fe40Al of similar grain size. As commented earlier, however, these results should be treated with caution, since the results on the Fe±Al±C alloy correspond to material in a well-relaxed state, and Figs. 1 and 2 suggest that better ductility may be found for similarly well-relaxed Fe±40Al than suggested in Fig. 4. Fig. 6 also shows that Fe±40Al alloys and the Fe±40Al±0.6C alloy of grain size about 100-200 mm have similar ductility, about 3±4%, when relaxed to a yield stress of about 300±350 MPa. Continuing the debate on the eect of C additions on ductility, Baligidad et al. [47,48] suggest that perovskite carbides present in Fe3Al may increase tensile ductility by trapping atomic hydrogen and reducing the sensitivity to environmental embrittlement. These arguments are supported by analogous evidence in steels [49,50] which suggests that carbide particles trap hydrogen, lower its diusivity in the lattice, and hence increase ductility by avoiding environmental embrittlement. The data con®rming such improved ductility is limited, however, and more research is needed. There have also been several reports that some special dislocation arrangement or recovered dislocation/subgrain/®ne-grain state, as found after annealing to near the beginning of recrystallization, may lead to improvements of ductility [51±54]. For example [51], a warm rolled and stress relieved Fe3Al base alloy, subsequently annealed at various temperatures to recover-recrystallize (plus a long anneal at 500 C to give a constant state of DO3 order) showed a notable increase in ductility for annealing temperatures of 600±700 C, where the rolled structure relaxed somewhat but did not yet show signs of recrystallization. During these anneals the yield stress did not change noticeably, con®rming that the improvements in ductility were due to subtle changes of dislocation structure, not major recovery or recrystallization. In the same way [52], hot rolling a similar Fe3Al at a range of temperatures, followed by a recovery anneal at 700 C, showed the best ductility when rolling was carried out at 500±700 C, where there were no signs of recrystallization, neither during hot rolling nor during the subsequent anneal, and an elongated grain structure was set up by rolling. Further studies [53] showed that it was neither the elongated grain structure itself, nor changes of environmental sensitivity, that was responsible for the improvements in ductility, but rather the relaxation of dislocations towards a subgrain substructure [53,54]. It was suggested that this relaxed dislocation substructure might ensure that subsequent slip was homogenised throughout the material, and the tendency to concentrated strain and stress was reduced. These results bear some resemblance to studies that
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show that DO3 order is characterised by lower ductility than is B2 order [55,56]. The origin of this eect is again possibly found in the greater tendency of dislocations in DO3 alloys to dissociate into partial dislocations (e.g. 2coupled instead of 4-coupled <111> dislocations as needed to constitute perfect dislocations) and the tendency to remain on the same slip planes, thereby increasing the local strain and stress concentrations. While both the eects described here have some relevance to the behaviour of Fe3Al base alloys under laboratory conditions, they have little importance when considering alloys of high Al contents, and where long times at moderate-to-high temperatures are considered. Such higher-Al alloys do not show DO3 order, of course and the special, relaxed dislocation substructures are likely to be unstable under these long time, moderate temperature conditions. There are, in fact, no reports of an analogous increase in ductility of FeAl alloys by processing or recovering at some intermediate temperature range, which may either suggest this as a topic for further research, or may be a con®rmation that the ductility enhancement eect is indeed restricted to Fe3Al base alloys showing DO3 order. A variety of other factures may play some role in modifying the ductility of iron aluminides, and some of these are brie¯y described here. A recent study of textures formed in Fe3Al base alloys [57] has considered the dierent textures produced at the surface of rolled sheet and in the centre during rolling and subsequent annealing [58]. The elastic anisotropy of iron aluminides, combined with varying textures from one place to another, may be a factor in¯uencing mechanical behaviour. Deliberate heat treatments to create surface ®lms on iron aluminide samples, or on analogous Al-containing steels [59], can lead to protective ceramic layers that serve not only to provide chemical protection (for example from subsequent environmental/water vapour embrittlement) but also create compressive stresses that can delay plastic deformation and crack formation at the surface. Such factors remain interesting topics for further investigation. 7. Summary The present review has examined the factors known to modify the ductility of iron aluminides. This topic is of great importance since the low ductility and toughness of these materials are seen as major reasons for their limited use today. The following ``rules'' should be followed to obtain high strength iron aluminides with good ductility. . The Al content should be reduced to the lowest possible level compatible with needs of, for example, oxidation resistance or density. Al levels certainly
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below 40% and if possible below 35% should be considered. Environmental (i.e. water vapour) embrittlement should be avoided or limited by atmosphere control, if this is possible. Additions of B (perhaps also C) help to improve the ductility by increasing the grain boundary cohesion or strength. The eect tends also to produce a change from intergranular failure to cleavage failure, and may be related to a reduction in the extent of environmental attack. Hardening by vacancies should be avoided. In fact any hardening mechanism, by solute, by precipitate or by vacancies, will lower ductility. There is perhaps some indication that vacancy hardening is less deleterious than other hardening mechanisms. Grain size re®nement is the most attractive way of improving both ductility and strength at the same time. For very ®ne grain sizes, only obtained in the presence of dispersoid particles needed to retain the ®ne grain size, there is some debate about whether such dispersoids may lead to additional improvements in ductility. Alloying, whether leading to solution hardening or to precipitate hardening, will generally also lower ductility. The only exception to this statement has been the case of Cr addition to Fe±28Al. This ductility improvement is not obtained for Fe± 45Al, and it would be interesting to know whether Fe±35/40Al alloys show such ductility improvement. Several other features of the microstructure have been discussed which may lead to interesting changes in ductility. The in¯uences of these parameters on ductility may be understood in terms of concentrated slip bands leading to crack nucleation, followed by crack propagation with only local plasticity taking place. Crack propagation, and possibly crack nucleation, is sensitive to hydrogen injection below the free surface as a result of reaction of the aluminide with water vapour.
References [1] Baker I, Xiao H, Klein O, Nelson C, Whittenberger JD. Acta Metall Mater 1995;43:1723. [2] Klein O, Baker I. Scripta Metall Mater 1994;30:1413. [3] Xiao H, Baker I. Scripta Metall Mater 1993;28:1411. [4] Nagpal P, Baker I. Mater Charact 1991;27:167. [5] Klein O, Baker I. Scripta Metall Mater 1994;30:627. [6] Gaydosh DJ, Draper SL, Nathal MV. Metall Trans 1989;20A:1701. [7] Gaydosh DJ, Draper SL, Noebe RD, Nathal MV. Mater Sci Eng 1992;A150:7. [8] Baker I, Munroe PR. Int Mater Rev 1997;42:181.
[9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43]
[44] [45] [46] [47] [48] [49] [50] [51] [52]
Stolo NS, Liu CT. Intermetallics 1994;2:75. Liu CT, Lee EH, McKamey CG. Scripta Metall 1989;23:875. Gaydosh DJ, Nathal MV. Scripta Metall Mater 1990;24:1281. Lynch RJ, Heldt LA, Milligan WW. Scripta Metall Mater 1991;25:2147. Lynch RJ, Gee KA, Heldt LA. Scripta Metall Mater 1994;30:945. Lynch RJ, Harburn M, Maucione L, Heldt LA. Scripta Metall Mater 1994;30:1157. Baker I, Klein O, Nelson C, George EP. Scripta Metall Mater 1994;30:863. Liu CT, George EP. Scripta Metall Mater 1990;24:1285. Cohron JW, Lin Y, Zee RH, George EP. Acta Mater 1998;46:6245. Nagpal P, Baker I. Metall Trans 1990;21A:2281. Baker I, Nagpal P, Liu F, Munroe PR. Acta Metall Mater 1991;39:1637. Munroe PR, Baker I. Acta Metall Mater 1991;39:1011. Nathal MV, Liu CT. Intermetallics 1995;3:77. Schultz PA, Davenport JW. Scripta Metall Mater 1992;27:629. Li JCM, Liu CT. Scripta Metall Mater 1992;27:1701. Chang YA, Pike LM, Liu CT, Bilbrey AR, Stone DS. Intermetallics 1993;1:107. Morris MA, George O, Morris DG. Mater Sci Eng 1998;A258:99. Morris DG, Joye JC, Leboeuf M. Phil Mag 1994;A69:961. Pang L, Kumar KS. Acta Mater 1998;46:4017. Cottrell AH. The mechanical properties of matter. NY: J.H. Wiley, 1963. Gaydosh DJ, Crimp MA. Koch CC, Liu CT, Stolo NS, editors. Pittsburgh, PA: Materials Research Society, 1985. vol. 39, p. 429. In: High-temperature ordered intermetallic alloys. Strothers SD, Vedula K. Prog Powder Metall 1987;43:597. McKamey CG, Maziasz PJ. Intermetallics 1998;6:303. Sun Y, Xue F, Mei J, Yu X, Zhang L. Scripta Metall Mater 1995;32:1181. Zhang Z, Sun Y, Guo J. Scripta Metall Mater 1995;33:2013. Zhang Z, Sun Y, Shen G. Scripta Mater 1998;38:21. Schneibel JH, George EP, Anderson IM. Intermetallics 1997;5:185. McKamey CG, Horton JA, Liu CT. J Mater Res 1989;4:1156. McKamey CG, Liu CT. Scripta Metall Mater 1990;24:2119. Balasubramaniam R. Scripta Mater 1996;34:127. Tassa O, Testani C, Ricci Bitti R. Scripta Metall Mater 1992;26:1813. Baccino R, San Filippo D, Moret F, Lefort A, Webb G. In: Powder metallurgy world congress, PM94, Paris, France, 6±9 June 1994, p.II.1239. Technical Information Sheet on FeAl 40 Grade 3, CEREM, CEA, Grenoble, France, c/o R. Baccino. Morris DG, Gunther S. Mater Sci Eng 1996;A208:7. Morris, DG. In Deevi SC, Maziasz PJ, Sikka VK, Cahn RW, editors. Int. Symp. on Nickel and Iron Aluminides, Processing, Properties and Applications, Materials Park, Ohio: ASM, 1997. p. 73. Carleton R, George EP, Zee RH. Intermetallics 1995;3:433. Maziasz PJ, Alexander DJ, Wright JL. Intermetallics 1997;5:547. Kerr WR. Metall Trans 1986;17A:2298. Baligidad RG, Prakash U, Radhakrishna A, Ramakrishna Rao V, Rao PK, Ballal NB. Scripta Mater 1997;36:667. Baligidad RG, Prakash U, Krishna AR. Mater Sci Eng 1998;A249:97. Hirth JP. Metall Trans 1980;11A:861. Gehrmann F, Grabke HJ, Riecke E. In: Turnbull A, editor. Hydrogen transport and cracking in metals. London: Inst. of Mater., London, 1995. p. 216. McKamey CG, Pierce DH. Scripta Metall Mater 1993;28:1173. Viswanathan S, Shelton BR, Wright JK, Sikka VK. Scripta Metall Mater 1993;29:589.
D.G. Morris, M.A. Morris-MunÄoz / Intermetallics 7 (1999) 1121±1129 [53] Lynch RJ, Gee KA, Heldt LA. Scripta Metall Mater 1994;30:945. [54] Morris DG, Leboeuf M. Meall Mater 1994;42:1817. [55] McKaney CG, Devan JH, Torturelli PF, Sikka VK. J Mater Res 1991;6:1779. [56] Sikka VK, McKamey CG, Howell CR, Baldwin RH. Fabrication and mechanical properties of fe3al-based aluminides,
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ORNL/TM-11465, Oak Ridge National Laboratory, Oak Ridge, TN (March 1990). [57] Morris DG, Gunther S. Scripta Mater 1996;35:1211. [58] Mao W, Sun Z. Scripta Metall Mater 1993;29:217. [59] Chao J, Gonzalez-Carrasco JL, IbanÄez J, Escudero ML, Gonzalez-Doncel G. Metall Mater Trans 1996;27A:3809.
The in¯uence of microstructure on the ductility of iron aluminides D.G. Morris*, M.A. Morris-MunÄoz Department of Physical Metallurgy, CENIM, CSIC, Avenida Gregorio del Amo 8, 28040 Madrid, Spain Dedicated to Dr. Gerhard Sautho on the occasion of his 60th birthday Received 1 March 1999; accepted 11 March 1999
Abstract Considerable eort has been devoted over the last decade to the development of iron aluminides as materials for high temperature applications, where their good oxidation and corrosion resistance, combined with reasonable strength, may be utilised. Poor formability and ductility, however, particularly at room temperature, has hampered the exploitation of these materials. The present review examines the present state of understanding of the factors which in¯uence the ductility. Recent research has made clear the important in¯uence of testing environment, the role of Al content and minor additions of B, as well as the eect of quenched-in vacancies. The extent to which other factors, such as alloying additions and microstructural features, aect the ductility has not received the same attention, and is examined in the present study. Alloy strengthening, by almost any mechanism, is seen to lead to a dramatic loss of ductility. The only parameter allowing both strength increase and ductility improvement for a given set of Al/B/ vacancy/environment conditions is the grain size. The best ductility for a given alloy, which should have as low an Al content as compatible with other requirements, is obtained by re®ning the grain size and by maintaining the alloy in the softest possible state. For the most part these conclusions are drawn from analysis of the behaviour of B2 ordered FeAl alloys, although similar trends seem also to apply to alloys of slightly lower Al content where DO3 ordering can occur. The observations drawn can be understood in terms of the mechanisms leading to the nucleation and propagation of brittle fracture, either as transgranular cleavage cracks or as grain boundary cracks. The possible role of additional factors, such as the texture, or grain and grain boundary distribution, surface layers producing protective stress eects, and strain homogenising or crack arresting dispersions, has not been suciently evaluated to determine whether any further improvements of ductility are possible. # 1999 Elsevier Science Ltd. All rights reserved. Keywords: A. Iron aluminides; D. Microstructure
1. Introduction Iron aluminides fail during tensile testing after limited plastic deformation associated with uniform gauge length elongation, and considerable work hardening, before ®nal fairly brittle failure with little local necking. As the Al content of iron aluminides increases from the Fe3Al composition to the FeAl composition there is a gradual decrease in ductility and a trend away from the ductile failure typical of Fe base alloys to cleavage failure and eventually to intergranular failure near stoichiometric FeAl. Fig. 1 and Table 1 [1±8] show the gradual fall in ductility as the Al content increases. For a long time the scatter in experimental data on fracture ductility of such iron aluminides has been so large that it has been dicult to deduce trends and explanations of the cause of ductility variations. It is only relatively recently that the important role played by such factors as environmental * Corresponding author.
embrittlement [9±14], and hence Al content and strain rate [15], B segregating to grain boundaries and modifying grain boundary behaviour [5,15±17], and hardening by quenched-in vacancies [5,7], have become clear. Since many of the mechanical testing experiments carried out earlier lent scant attention to some or all of these factors, the literature shows very wide scatter of fracture data both as a function of speci®c test parameters as well as between investigators and reports. Fig. 1 and Table 1 make clear the dramatic loss of ductility as the Al content increases from near 40 to 50% in FeAl alloys, when testing is carried out in air instead of under vacuum or in dry oxygen, and also in the absence of B from the alloy. Note that throughout this text all compositions are given as atomic percentages. The dramatic eect due to the choice of environment where testing is carried out [9±14] has been explained by the reaction of freshly-exposed intermetallic at a propagating crack tip with water vapour, whereby atomic hydrogen is injected into the material in
0966-9795/99/$ - see front matter # 1999 Elsevier Science Ltd. All rights reserved. PII: S0966-9795(99)00038-2
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Fig 1. Tensile ductility of well-annealed FeAl in air and in vacuum as a function of Al content. This ®gure is taken from the review of Table 1 In¯uence of Al content , B additions and testing atmosphere on tensile ductility of FaAla Aluminium content (%) Atmosphere/B addition?
35
40
45
50
Air No B
TG 2%
Mixed 1%/3%
IG 3%/1%
IG 0%
Vac/oxygen No B
IG 5.5/18%
IG 5%
±
±
Air B Oxygen B
±
TG 4% IG 17%
TG 2% ±
±
±
±
a TG, IG and mixed refer to transgranular, intergranular and mixed (transgranular and intergranular) failure modes. All alloys examined are expected to have B2 order as it is unlikely that any DO3 order appears for Al contents above 35%. The data shown are approximate average behaviours reported by a number of workers, including [1,4,5,10,11,16].
front of the tip leading to easier failure [10]. Since such embrittlement aects cleavage planes more than grain boundaries, an additional eect of the environmental attack is the tendency to cleavage failure rather than intergranular failure [10] (Table 1). It may also be noted that since the environmental attack takes place especially on cleavage planes it is a phenomenon found also for single crystals [13]. In Table 1 a gradual reduction of ductility is evident as the Al content increases. Increases in the Al content, as well as testing in vacuum/oxygen, as well as the absence of B, all tend to lead to intergranular failure [1,4,5,10,11,16,17]. These trends are due to the intrinsic weakness of grain boundaries at high Al levels, the preference for environmental attack on cleavage planes, and the eect of B on improving grain boundary adhesion. It should be noted that over the range of alloys and test conditions discussed here the general mechanical
properties of FeAl alloys are only slightly modi®ed (for example by Al content over the range of say 35 to above 45% and even as high as 48%) and hence the various test parameters truly aect fracture behaviour, not general plasticity. For example, the hardness of wellannealed alloys over the range 35% to above 45% Al [18], the yield stress of alloys over the range 40±48% Al [1] and also the work hardening rate over the same 40± 48% Al range [1], as well as the Hall±Petch slope over the range from below 35% Al to above 45% Al [19], are essentially constant, such that any change of ductility recorded (Table 1, Fig. 1) is indeed due to changes of fracture behaviour. Deformation mechanisms in iron aluminides at temperatures near ambient temperatures involve the motion of superdislocations of <111> Burgers vector generally gliding on {110} planes with perhaps some cross slip onto {112} planes. This aspect of deformation has been examined on many occasions, and further detailed references may be found in the review by Baker and Munroe [8]. The factors discussed above concern in¯uences on the propagation of an already existing crack, for example the eects of environmental embrittlement. The nucleation of such a crack presumably occurs by the interaction of various dislocations or slip systems. A possible nucleation mechanism has been proposed by Munroe and Baker [20] involving the reaction of <111> superdislocations of dierent Burgers vectors on intersecting glide planes to produce what are eectively small pile-ups of edge <100> dislocations. If the edge dislocations produced are suciently sessile they could represent the initial stage of formation of a cleavage crack on a {100} plane. This plane has indeed been reported to be the cleavage plane for iron aluminides [7,21]. Any alloying or microstructural factor which in¯uences the tendency for <111> superdislocations to come together to produce such sessile <100> edge dislocation pile-ups will in¯uence the tendency to crack initiation. Coarser microstructures, such as larger grain sizes, would allow longer slip bands of the initial <111> superdislocations to be created with correspondingly higher stresses available to create larger numbers of <100> edge dislocations. It has also been suggested that the presence of hydrogen may change the mobility of <111> and <100> dislocations [22] or the elastic energy of <111> or <100> dislocations [23]; less mobile <100> dislocations of lower elastic energy should be easier to create, allowing easier crack initiation. As a ®nal comment, it is emphasized that the intention of the present short review is to demonstrate in¯uences of alloying and microstructural parameters on fracture behaviour and ductility, and as such the present discussion of the general fracture mechanisms, and in¯uences of testing conditions (including Al and B content) on behaviour is deliberately restricted to being a very brief overview. Excellent, more detailed, reviews
D.G. Morris, M.A. Morris-MunÄoz / Intermetallics 7 (1999) 1121±1129
have been presented elsewhere, e.g. [8,9]. Fundamental understanding of the fracture processes, however, both crack initiation and crack propagation, remain poorly or incompletely understood, even though there is now a much better empirical description of the in¯uences of the various test and alloy parameters. 2. Vacancy hardening, and the in¯uence of annealing treatments on ductility It is now well established that large numbers of vacancies will be present in iron aluminides cooled from high temperatures, for example after casting, recrystallization or a solutionising anneal, and these will only be removed by long anneals at relatively low temperatures [18]. Such vacancies are known to lead to considerable hardening of the materials [18,24], presumably by vacancies acting as solute for pinning dislocations, while vacancy aggregates such as dislocation loops [25] and voids [26] may also cause similar hardening. The role of such point defect hardening on ductility is also now clear, as illustrated in Fig. 2, for a Fe±45Al alloy, both with and without B additions [5,8]. Furnace cooling, which leaves a fairly high vacancy concentration, is seen to lead to fairly low ductility for these alloys, while removing vacancies by a slow cool from high temperature and a long anneal at low temperature leads to progressively improved ductility. Such improvements of ductility by vacancy removal have now been seen on a number of occasions [7,11,27]. The relationship between hardening and embrittlement by vacancies has been evaluated in more detail by Pang and Kumar [27], as summarised in Fig. 3, showing the yield stress after dierent high temperature anneals, quenches, and low temperature ageing, and the related tensile ductility. It should be noted that these results were obtained on a Fe±40Al alloy containing in addition 0.6%
Fig. 2. In¯uence of dierent heat treatments to reduce the retained vacancy concentration on the ductility of Fe±45Al alloys, with or without 400 ppm B [5].
1123
carbon, and this material contained some carbon in the form of fairly large perovskite (Fe3AlC) carbides, which may themselves play some role in aecting fracture [27], as discussed later. During high temperature annealing and subsequent ageing, some part of the carbon may dissolve and subsequently precipitate [27] and hence annealing eects reported may be compound eects due to both vacancy and to carbon eects. At the same time, the interaction of lattice vacancies and interstitial C atoms is likely to modify the mobility and tendency to agglomeration of vacancy defects [27]. The ®lled symbol data points of Fig. 3 show the eect of ageing at about 400 C, starting from a high temperature extruded state, on ductility and yield stress: the eect of such ageing is a steady reduction of stress (from about 500 MPa to near 300 MPa) and a concomitant increase of ductility (from near 2% to above 4%). The open symbol data points in Fig. 3 show the results of several experiments on quenched material which is subsequently aged at dierent temperatures, with mechanical testing carried out at various times during ageing. These data show, interestingly, that there is some slight age hardening, probably associated with the carbon precipitation, and at the same time a notable embrittlement. One such age hardening-embrittlement cycle is indicated by the sequence of arrows in Fig. 3. Subsequent annealing to soften leads to an increase in ductility again. Despite some scatter in the ductility-stress data of Fig. 3 about the best ®t line, it is clear for this material that ductility and yield stress are essentially reciprocally related by a single master curve. During ageing of such quenched materials it is well known that the vacancies
Fig. 3. In¯uence of heat treatments on yield stress and tensile ductility of Fe±40Al containing 0.6% carbon [27]. The ®lled symbol data points show the softening and ductilising eect of ageing material extruded (at 1000 C) at low temperature (400±500 C). The open symbol data points show ageing data for material well quenched from high temperature (1000 C), when the grain size has increased somewhat. The sequence of arrows shows an example of the variation of yield stress and ductility on progressively ageing a quenched material at one ageing temperature (500 C). Data for other ageing temperatures (400 C for strong±brittle materials and 650 C for softer±ductile materials) show similar stress±ductility trends.
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initially present agglomerate, generally as dislocation loops of <111> Burgers vector, which subsequently grow, interact, and ®nish as a low density of <100> dislocations, [25]. The implication of the master curve representation of Fig. 3 is that the precise mode of strengthening, by a solution of vacancies or by a population of <111> or <100> dislocations, is immaterial. The only relevant parameter determining ductility is the yield stress, independent of the precise mechanism causing hardening. The earlier section described failure in terms of mechanisms of crack formation and growth. It seems clear from the present data that it is the absolute level of stress alone that determines the ease of crack nucleation and propagation, with little to no in¯uence of the precise mechanism that causes the hardening. 3. In¯uence of grain size on ductility In general terms it is well established that grain size re®nement is a potent way to improve ductility or to bring about a brittle to ductile transition [28]. These eects are deduced from the grain size dependence of the yield stress, for the initiation of plastic ¯ow, as described by the Hall±Petch relationship (y i ky dÿ1=2 ) and from the similar dependence of the fracture stress, characterised by a smaller (or zero) value of internal stress (i ) and a higher value of the grain size dependency slope kf . While some researchers have reported an increase in ductility as the grain size is reduced [7] other times it is reported that grain size re®nement has only a minor eect on ductility [8,17]. For example, Cohron et al. [17] compared the ductilities of Fe±45Al with 300 ppm B with grain sizes of 22 mm and 120 mm to ®nd almost identical ductility in the two cases. Baker and Munroe [8] also concluded that the slight improvements in ductility reported by grain re®nement were generally overshadowed or confused with other factors aecting ductility, for example due to particles introduced to ensure grain re®nement or due to environmental embrittlement eect, and it is perhaps only in those cases where other factors aecting ductility are very well controlled, or their eects are avoided, that grain size in¯uences can be detected. Fig. 4 shows variations of tensile ductility of Fe-40Al as the grain size varies over a very wide range Ð from single crystal down to grain sizes of about 5 mm Ð and shows a clear increase in ductility, particularly for the very ®ne grain sizes [7,21,29,30]. It should be emphasized, however, that the ®nest grain sizes are obtained using powder metallurgy/rapid solidi®cation techniques, and that dispersoid particles are often present for grain boundary pinning, but these may also serve to disperse slip and delay crack nucleation. Over a narrow grain size range, however, and particularly for the fairly large
Fig. 4. Eect of grain size reduction, from single crystal to about 5 mm, on tensile ductility of Fe±40Al [7,21,29,30] (®lled symbol data points) and Fe±40Al±0.6C [27] (open symbol data points).
grain sizes obtained by conventional casting/extrusion methods, the variation of ductility with grain size is very small, as illustrated by the data on Fe-40Al with carbon [27]. An interesting point to be made with regards the data of Fig. 4 is that the re®nement of grain size will be associated with a signi®cant strengthening of the material [7], as indeed expected from the Hall±Petch relationship. As discussed earlier, and seen in Fig. 3, an increase in strength is expected to be associated with a reduction in ductility, at least for a constant grain size. The observed increase in ductility associated with grain size reduction and hardening is indeed proof of the powerful eect of grain size reduction on delaying cracking. Fig. 4 also suggests that the addition of carbon to Fe± 40Al serves to increase the ductility signi®cantly. In this sense the C addition appears to produce similar gains in ductility to that achieved by B additions, Fig. 2 for example. This interpretation should be considered with great care, however, because the carbon-containing alloy achieved the good ductility shown in Fig. 4 only after careful annealing at low temperature to remove vacancy hardening, and as seen in Fig. 2 such annealing can considerably improve the ductility levels. 4. In¯uence of substitutional alloying on ductility The in¯uence of substitutional alloying on ductility has been examined in great detail for alloys containing about 28% Al [31±34], where DO3 order may be found, with less work reported on the B2 alloys containing 40% Al or more [35]. Eects of such alloying on Fe28Al-based alloys are summarised in Fig. 5, where yield stress and tensile ductility are shown. The present discussion will consider changes of ductility and stress in terms of such alloying, considering that all other parameters are constant, e.g. grain size, test environment,
D.G. Morris, M.A. Morris-MunÄoz / Intermetallics 7 (1999) 1121±1129
Fig. 5. Changes of yield stress and tensile ductility with additions of various alloying elements (Mo, Nb, W) to the Fe±28Al or to the Fe± 28Al±5Cr base [31±34]. The Cr addition to the Fe±28Al base leads to increased strength and ductility. Additions of Mo, Nb, or W to the respective alloy bases lead to increased strength with only small changes of ductility for small alloying additions, and typically further strengthening with some loss of ductility for larger alloying additions.
strain rate, etc. As dierent amounts of various alloying additions were used for the experiments of Fig. 5 the grain size varied from around 50 to 100 mm for the binary alloy, or with Cr addition, to near 30 mm with large additions of Mo, Nb or W. As seen from Fig. 4 such reductions of grain size do not lead to major changes of ductility. The most dramatic eect shown in Fig. 5 is the increase in ductility on adding Cr to the binary Fe±28Al alloy. This increase was ®rst ascribed to changes of plasticity of the material [36], but has subsequently been associated with changes in the nature of surface ®lms and the extent of environmental embrittlement [37,38]. Indeed it is possible that all the changes of ductility in Fig. 5 are related to the environmental embrittlement eect and some quality of the surface ®lms produced before or during testing. It should also be remembered that all the alloys discussed here show DO3 order after annealing at moderate temperatures, and the eects seen may be somewhat aected by slight changes of the degree of this order. Fig. 5 shows that, with the exception of the ductilising eect of Cr, alloying by any addition (of Nb, Mo or W) leads to a progressive loss of ductility. Hardening and embrittlement occur when the alloying additions are dissolved as solute (e.g. 0.5±1% Mo) or as second phase intermetallics or carbide particles (large additions of all the solutes examined). Nevertheless, for the Fe±28Al± 5Cr intermetallic base it is possible to achieve notable strengthening (for example to a yield stress near 600MPa) by Mo and particularly by Nb additions without dramatic loss of ductility. Whilst the ternary (or quaternary) additions to Fe± 28Al(+Cr) can strengthen and lead to only a gradual
1125
loss of ductility, there is no evidence that alloying with any substitutional element (including Cr) to Fe±45Al produces any bene®cial eects [35]. In this study, even the addition of Cr leads to a fall in ductility, and this for any testing environment. Some other studies have examined alloying additions to FeAl alloys, for example [6], and have typically considered alloying additions that lead to second phase precipitation. Gaydosh et al. [6] considered additions of B and C alone to the Fe± 40Al base, as well as small additions of substitutional elements such as Zr and Hf, and also combined additions of the light and the heavy elements. The only material showing improved room temperature ductility over the binary base alloy was the one containing both Zr and B, with neither Zr alone, C alone, nor C+B alone improving ductility. The improvement of ductility observed in this case may be due to the eect of B in increasing grain boundary strength, although there may also be an eect of particles of Zr intermetallics or borides on increasing ductility Ð by tending to homogenise slip Ð or on decreasing ductility Ð if the coarse particles lead to stress concentrations and hence nucleate cracks. 5. Yield stress and grain size in¯uences on ductility The previous discussions have demonstrated the important role that grain size of a given material, and extent of hardening for a material of given grain size, will play in determining the tensile ductility (Figs. 3 and 4), and this for material tested under reasonably similar conditions of test environment and strain rate, and possessing similar chemical compositions in terms of Al content, B additions, and so on. Fig. 6 attempts to correlate the tensile ductility (in air, typical laboratory strain rates, etc.) of thermally-relaxed Fe40Al (generally containing small amounts of B or C, but no signi®cant amounts of substitutional additions) with the two parameters Ð grain size and yield stress [6,7,27,30,35,39±45]. The causes of strengthening and ductility changes for a given material and grain size are multiple, ranging from the extent to which vacancies are retained or completely annealed out with very low dislocation density, some retention of work hardening from material preparation, small alloying and impurity additions, etc. It should also be remarked that, while the data are collected with approximate grain size groups in Fig. 6, the grain sizes reported in the literature vary steadily from very small to very large without having the precise values suggested in Fig. 6. The lines connecting data of approximately the same grain size in Fig. 6 show that to a reasonable approximation it is indeed possible to correlate ductility with yield stress. Despite the extensive work carried out on iron aluminides over the past decades the available data are still limited and only 5±10 data points are available
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Fig. 6. Variation of tensile ductility in Fe±40Al alloys as a function of the grain size and the yield stress. Data taken from the literature sources [6,7,27,30,35,39±45]. Data are grouped corresponding to approximately constant grain sizes (microns), with strength±ductility data correlated for each approximate grain size. The data points refer generally to single phase or dispersion-strengthened materials prepared by typical casting, thermomechanical processing, or powder metallurgy methods. The open data points refer to the Fe±40Al±0.6C material [27] discussed earlier, where strength and ductility variations are obtained in the same material by heat treatments. Some of the ®lled data points lie o their grain-size trend lines, and are shown attached to this by dotted lines: these data refer to materials of lower Al content (e.g. 36%Al, [45]), when the ductility-strength is above the trend line, or of higher Al content (e.g. 45%Al, [35]), when the ductility±strength is below the trend line, taking account of the in¯uence of Al content on the ductility.
for representation at each grain size. The scatter of the data is generally not excessive bearing in mind the number of experimental variables which change slightly from one literature case to another. The ®lled data points of Fig. 6 refer to samples prepared under similar conditions by the same workers during a single study or by a number of workers during several studies. The open data points, linked by a dotted line, show data on the same Fe±40Al±0.6C alloy [27] as in Fig. 3 and refer to the in¯uence on yield stress and ductility of dierent quench and ageing conditions to modify the number of retained vacancies and vacancy aggregates. The smaller slope of the line connecting these data suggests that vacancy hardening is a powerful way of hardening while retaining relatively high ductility (for the grain size and hardness), or otherwise that the presence of C in solution in the quenched and lightly aged material may be responsible for this high hardening. Grain size reduction by a large amount is con®rmed to be a powerful way of increasing the ductility Ð yield stress combination, with reductions of grain size by a factor of 4±5 leading to the equivalent change over the entire grain size range. For example, grain size reduction from 200±50 mm to 10±2 mm allows increases of yield stress from 300±450 to 600±750 MPa for 5% tensile
ductility or, otherwise, at a yield stress of 500 MPa the ductility increases from 0±4 to 10±20%. Fig. 6 shows also several individual data points connected to their grain-size lines by dotted connecting lines (for example there are two such data points near 12±15% ductility and 650 MPa). These data refer to materials which do not have 40% Al but instead 36% Al or 45% Al [35,45]. The several data points displaced to the right of their grain size lines (for example the two data points described before: ductility 12±15%, stress 650 MPa) correspond to an alloy containing 36% Al [45]. This material is somewhat more ductile (or stronger for a given ductility and grain size) than the corresponding Fe±40Al alloy, but not to an extent that suggests that dierent failure mechanisms operate. Similarly, the one data point displaced to the left of its grain size line (near a ductility of 8% and yield stress of 300 MPa) corresponds to an alloy containing 45% Al [35]. Again this material is more brittle than the Fe-40Al alloy of same grain size and yield stress, but not to an extent incompatible with the present yield stressductility-grain size relationships. 6. Other microstructural parameters which may aect ductility There have been some reports that the presence of dispersoid particles within the iron aluminide matrix can serve to distribute plastic ¯ow, avoiding or at least delaying the formation of intense slip bands and stress concentrations and thereby delaying crack nucleation and failure. Data in Fig. 6 of Maziasz et al. [45] (grain size line of 1±3 mm and ductility of 12±15%), and data of Strothers and Vedula [30] (also grain size line of 1±3 mm and ductility of about 7%) correspond to such cases of dispersed particles throughout the material. The question of importance is whether these materials show their good ductility because of their ®ne grain size and relatively low yield stress Ð i.e. they sit on the ductility± yield stress line that corresponds to the 1±3 mm grain size Ð or whether they show exceptionally good ductility Ð i.e. they sit above the true ductility±yield stress line for their grain size. The data presently available are insucient to provide a clear answer. In another study of Fe±40Al alloys, however, Gaydosh et al. [6] decided that rather large dispersed particles, of the order of several microns in size, could nucleate cracking either at the particle-matrix interface or by particle cracking, and such coarse particles should clearly be avoided. Another interesting question concerns the special role of carbide particles in iron aluminides where, again, the evidence is con¯icting. In a study of the eect of carbon in Fe3Al, Kerr [46] concluded that carbon additions could lead to embrittlement, although it was not completely clear whether carbon in solution, some in¯uence
D.G. Morris, M.A. Morris-MunÄoz / Intermetallics 7 (1999) 1121±1129
at grain boundaries, or Fe3AlC perovskite particles were responsible. On the other hand, Pang and Kumar [27] decided that C additions to Fe±40Al could enhance ductility, as implied in Fig. 4, where the C-containing alloy shows greater ductility than the carbon-free Fe40Al of similar grain size. As commented earlier, however, these results should be treated with caution, since the results on the Fe±Al±C alloy correspond to material in a well-relaxed state, and Figs. 1 and 2 suggest that better ductility may be found for similarly well-relaxed Fe±40Al than suggested in Fig. 4. Fig. 6 also shows that Fe±40Al alloys and the Fe±40Al±0.6C alloy of grain size about 100-200 mm have similar ductility, about 3±4%, when relaxed to a yield stress of about 300±350 MPa. Continuing the debate on the eect of C additions on ductility, Baligidad et al. [47,48] suggest that perovskite carbides present in Fe3Al may increase tensile ductility by trapping atomic hydrogen and reducing the sensitivity to environmental embrittlement. These arguments are supported by analogous evidence in steels [49,50] which suggests that carbide particles trap hydrogen, lower its diusivity in the lattice, and hence increase ductility by avoiding environmental embrittlement. The data con®rming such improved ductility is limited, however, and more research is needed. There have also been several reports that some special dislocation arrangement or recovered dislocation/subgrain/®ne-grain state, as found after annealing to near the beginning of recrystallization, may lead to improvements of ductility [51±54]. For example [51], a warm rolled and stress relieved Fe3Al base alloy, subsequently annealed at various temperatures to recover-recrystallize (plus a long anneal at 500 C to give a constant state of DO3 order) showed a notable increase in ductility for annealing temperatures of 600±700 C, where the rolled structure relaxed somewhat but did not yet show signs of recrystallization. During these anneals the yield stress did not change noticeably, con®rming that the improvements in ductility were due to subtle changes of dislocation structure, not major recovery or recrystallization. In the same way [52], hot rolling a similar Fe3Al at a range of temperatures, followed by a recovery anneal at 700 C, showed the best ductility when rolling was carried out at 500±700 C, where there were no signs of recrystallization, neither during hot rolling nor during the subsequent anneal, and an elongated grain structure was set up by rolling. Further studies [53] showed that it was neither the elongated grain structure itself, nor changes of environmental sensitivity, that was responsible for the improvements in ductility, but rather the relaxation of dislocations towards a subgrain substructure [53,54]. It was suggested that this relaxed dislocation substructure might ensure that subsequent slip was homogenised throughout the material, and the tendency to concentrated strain and stress was reduced. These results bear some resemblance to studies that
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show that DO3 order is characterised by lower ductility than is B2 order [55,56]. The origin of this eect is again possibly found in the greater tendency of dislocations in DO3 alloys to dissociate into partial dislocations (e.g. 2coupled instead of 4-coupled <111> dislocations as needed to constitute perfect dislocations) and the tendency to remain on the same slip planes, thereby increasing the local strain and stress concentrations. While both the eects described here have some relevance to the behaviour of Fe3Al base alloys under laboratory conditions, they have little importance when considering alloys of high Al contents, and where long times at moderate-to-high temperatures are considered. Such higher-Al alloys do not show DO3 order, of course and the special, relaxed dislocation substructures are likely to be unstable under these long time, moderate temperature conditions. There are, in fact, no reports of an analogous increase in ductility of FeAl alloys by processing or recovering at some intermediate temperature range, which may either suggest this as a topic for further research, or may be a con®rmation that the ductility enhancement eect is indeed restricted to Fe3Al base alloys showing DO3 order. A variety of other factures may play some role in modifying the ductility of iron aluminides, and some of these are brie¯y described here. A recent study of textures formed in Fe3Al base alloys [57] has considered the dierent textures produced at the surface of rolled sheet and in the centre during rolling and subsequent annealing [58]. The elastic anisotropy of iron aluminides, combined with varying textures from one place to another, may be a factor in¯uencing mechanical behaviour. Deliberate heat treatments to create surface ®lms on iron aluminide samples, or on analogous Al-containing steels [59], can lead to protective ceramic layers that serve not only to provide chemical protection (for example from subsequent environmental/water vapour embrittlement) but also create compressive stresses that can delay plastic deformation and crack formation at the surface. Such factors remain interesting topics for further investigation. 7. Summary The present review has examined the factors known to modify the ductility of iron aluminides. This topic is of great importance since the low ductility and toughness of these materials are seen as major reasons for their limited use today. The following ``rules'' should be followed to obtain high strength iron aluminides with good ductility. . The Al content should be reduced to the lowest possible level compatible with needs of, for example, oxidation resistance or density. Al levels certainly
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below 40% and if possible below 35% should be considered. Environmental (i.e. water vapour) embrittlement should be avoided or limited by atmosphere control, if this is possible. Additions of B (perhaps also C) help to improve the ductility by increasing the grain boundary cohesion or strength. The eect tends also to produce a change from intergranular failure to cleavage failure, and may be related to a reduction in the extent of environmental attack. Hardening by vacancies should be avoided. In fact any hardening mechanism, by solute, by precipitate or by vacancies, will lower ductility. There is perhaps some indication that vacancy hardening is less deleterious than other hardening mechanisms. Grain size re®nement is the most attractive way of improving both ductility and strength at the same time. For very ®ne grain sizes, only obtained in the presence of dispersoid particles needed to retain the ®ne grain size, there is some debate about whether such dispersoids may lead to additional improvements in ductility. Alloying, whether leading to solution hardening or to precipitate hardening, will generally also lower ductility. The only exception to this statement has been the case of Cr addition to Fe±28Al. This ductility improvement is not obtained for Fe± 45Al, and it would be interesting to know whether Fe±35/40Al alloys show such ductility improvement. Several other features of the microstructure have been discussed which may lead to interesting changes in ductility. The in¯uences of these parameters on ductility may be understood in terms of concentrated slip bands leading to crack nucleation, followed by crack propagation with only local plasticity taking place. Crack propagation, and possibly crack nucleation, is sensitive to hydrogen injection below the free surface as a result of reaction of the aluminide with water vapour.
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