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Observation of 〈001〉 dislocations and a mechanism for transgranular fracture on {001} in FeAl
Acta memll, mater. Vol. 39, No. 5, pp. 1011-1017, 1991
0956-7151/91 $3.00 + 0,00 Copyright © 1991 Pergamon Press plc
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OBSERVATION OF (001) DISLOCATIONS AND A MECHANISM FOR TRANSGRANULAR FRACTURE ON {001} IN FeA1 P. R. MUNROE and I. BAKER Thayer School of Engineering, Dartmouth College, Hanover, NH 03755, U.S.A.
(Received 13 February 1990; receivedfor publication 12 November 1990) Alam'aet--Transmission electron microscopy of iron-rich, B2-structured, FeA1 alloys which had been compressed at room temperature revealed the presence of (001) dislocations, in addition to APB-coupled <111) dislocations which are normally observed at room temperature. It appears that the (001) dislocations form through the interaction of the (11 I) dislocations. This reaction to produce <001) dislocations is outlined and shown to be energetically favorable. The production of these (001) dislocations on {100} planes is suggested to be a cause of transgranular cleavage on {100} in B2 alloys deforming by < I 11) slip. Rtmmt---La microscopie 61ectronique en transmission d'alliages FeA1 (de structure B2) riches en fer, comprim~s a la teml~rature ambiante, met en 6vidence la presence de dislocations (001), en plus des dislocations (I 11), coupl6es aux parois d'antiphase, que l'on observe en principe :i la temp6rature ambiante. II semble que ces dislocations se forment ~i partir de rinteraction des dislocations (I 11). On pr~sente rapidement cette r~action et l'on montre qu'eUe est 6nerg6tiquement favorable. On sugg~re que la creation de ces dislocations (00 ! ) darts les plans {100} est une cause du clivage transgranutaire suivant les plans {I00} dam les alliages B2 qui se d6forment par le glissement (111). Zusammenfassung~In eisenreichen FeA1-Legierungen mit B2-Struktur, die bei Raumtemperatur druckvefformt worden find, werden elektronenmikroskopisch (001)-Versetzungen beobachtet. Diese treten zus/itzlich zu den normalerweise bei Raumtemperatur beobachteten, mit Antiphasengrenzen behafteten (111 )-Versetzungen auf. Es scheint, dab die (001)-Versetzungen aus einer Reaktion yon (111 )-Versetzungen entstehen. Diese Reaktion wird dargestellt; sie ist energetisch bevorzugt. Die Entstehung dieser (001)-Versetzungen auf {111}-Ebenen wird als eine Ursache der transgranularen Spaltung auf {I 1l}-Ebenen in B2-Legierungen, die sich mit (11 l)-Gleitung verformen, angesehen.
1. INTRODUCTION FeAI is a B2-structured compound which exists over a wide range of composition (34-51 at.% AI). Previous studies have shown that FeA1 deforms by (111 ) slip at all compositions at room temperature [1-14], and by (001) slip at elevated temperatures [5-7, 12-14], the transition temperature from (111) to (001) slip increasing with increasing iron-rich deviations from stoichiometry [7, 14]. A significant difference between the two slip vectors is that whereas 111 ) slip provides sufficient slip systems to satisfy Von Mises criterion for plastic flow in polycrystals <001) slip does not. The ductility of FeAI varies with composition: the ductility can also be affected by several extrinsic factors such as the testing environment [15] and the concentration of vacancies retained after heat treatment [16]. Stoichiometric FeAI is brittle at room temperature, generally fracturing before yielding in tensile tests [17], but the ductility increases as FeAI becomes increasingly iron-rich, for example, Fe-36.5 at.% A1 can exhibit as much as 12% tensile elongation [15]. The fracture mode of FeAI at room
temperature is affected by both grain size and composition [18]. Stoichiometric FeA1 exhibits predominantly intergranular fracture, but the incidence of transgranular fracture increases with both increasing iron-rich deviation from stoichiometry and increasing grain size [18]. This paper reports upon the observation of (001) dislocations in FeA1 after room temperature compression. These (001) dislocations are formed through the interaction of gliding < 111 ) dislocations. A mechanism for transgranular fracture based upon the occurrence of these <001) dislocations is outlined. 2. EXPERIMENTAL Three FeAl alloys, with compositions of F¢-34 A1, F¢-40 Al and F¢--45 A1 (compositions in at.%), were studied. The processing routes of these alloys have been described elsewhere [14]. Cylinders ~ 10ram long by 3 mm diameter were annealed at 1123 K for one hour (this removed the residual dislocations resulting from prior processing) and water-quenched. The Fe--40Al and Fe--45Al samples were then
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1012
MUNROE and BAKER: OBSERVATION OF (001) DISLOCATIONS
annealed at 673 K for 120h to reduce the high vacancy concentrations which are retained after annealing at temperatures above ~ I 0 0 0 K [16]. Fe-34 A1 can exist in both the B2 and DO3 states, the DO3 state being the equilibrium phase below ~ 700 K [19]. Thus, the Fe-34 AI sample was not annealed to reduce the vacancy concentration. So that the B2 microstructure would be retained, it was compressed after quenching alone (vacancy retention appears to be less pronounced in Fe-34 AI [16]). The cylinders were then strained ~ 1% in compression at a strain rate of ~ 1 0 - ~ s -~, at room temperature. Thin foils were prepared from all the alloys as described elsewhere [20], and examined in a JEOL 2000FX transmission electron microscope (TEM) operating at 200 kV. Dislocation analyses were performed using the g.b. = 0 invisibility criterion [21]. Samples of Fc-34AI, Fc-40Al and Fe-45Al were fractured in bending at both room temperatures and liquid nitogen temperatures and the fracture surfaces were examined using selected area electron channeling in the J E O L 2000FX operated in S E M mode. 3. RESULTS The microstructures of all the alloys were examined prior to compression and consisted of equiaxed, dislocation-free grains ~ 50 # m in diameter. Selected area difffraction patterns from Fe-34 AI indicated that water quenching from 1123K had retained the B2 structure (Fig. 1), (111) superlattice reflections from the DO3 structure were extremely diffuse, indicating that the DO3 phase had been largely suppressed by water quenching. The dislocation substructure of Fe--34 AI after compression is shown in Fig. 2. The dislocations marked a show either invisibility or residual contrastt when g = 110 and g = 10I, indicating that b = l l I , whilst the dislocations marked b show either invisibility or weak contrast when g = 110, g = 01T and g = 10I, indicating that b = 111. Trace analysis was used to show that these (111 ) dislocations are screw in character. In addition, dislocations marked c were observed, which are invisible when g = 200 and g = 10I, indicating a [010] Burgers' vector. The Burgers' vector must be a [010], rather than a/2[010], since no fault contrast was observed between the dislocations. This was also confirmed by imaging the dislocations with g = 020 and s -- 0, when a double image was observed (Fig. 3), consistent with g. b = 2 [21].Trace analysis showed that the predominant general directionof the
tFeA1 is strongly anisotropic. The Zoner anisotropy factor, A, of Fv--34AI is 4.4, and of F¢-40 AI is 3.8 [22], and there is often residual contrast even when a screw dislocation is viewed under conditions which satisfy the g" b = 0 invisibility criterion.
C)©?
Fig. 1. Selected area diffraction pattern from (110) zone axis in F¢-34 A1, heat treated at 1123 K for I h and waterquenched. Note the absence of DO3(III } superlattice reflections, which suggests that the DO 3 structure has been suppressed.
dislocation line was [100]. Tilting experiments which examined how the dislocation separation changed with the beam normal showed that the [010] dislocations lay on the (001) plane. A (100) line direction has been shown to be a stable orientation for (001) { 100} dislocations in anisotropic B2 compounds [23]. To investigate whether the [010] dislocations were part of a loop or a dipole, ((111 ) dislocation dipoles are commonly observed in FeAI alloys [14]), they were imaged using both + g and - g two-beam conditions. If they were dislocation loops or dipoles, a change in the spacing between the two dislocations would be expected on reversing g [21]. No change in spacing was observed, indicating that the Burgers' vectors of the two dislocations are the same. In addition, dynamical fringe contrast from each dislocation, when s -- 0, lies on the same side of the defect also indicating that each dislocation has the same sign (Fig. 4). Figure 5 shows images of a relatively isolated dislocation network similar to that described as c in Fig. 2, within the same grain as the dislocations described above. At each of the junctions of the two [010] dislocations ( I 11 ) dislocations can be observed. At the junction marked X there is both a [I11] and a [ I l l ] dislocation present, at the junction marked Y only a [111] dislocation can be observed. It was assumed that the [ I l l ] dislocation at the junction Y was not present in the foil. These can be most clearly seen when g = 200 and the [010] dislocations are out of contrast [Fig. 5(a)]. One of the [010] dislocations in Fig. 5, marked a, is relatively straight with a [100] line direction, whereas the part marked b initially curves away from the two junctions but then straightens to take up a line direction close to [100] away from the junctions, giving the defect a "coathanger" shape. Jogs can also be seen along the length of the [010] dislocations, marked J [Fig. 5(b)],
MUNROE and BAKER: OBSERVATION OF (001) DISLOCATIONS
1013
e
i
/
Fig. 2. Bright field transmission electron micrograph of dislocation structure in Fe--34AI after 1% compression at room temperature. The dislocations marked a and b have [I 1I] and [111] Burgers' vectors respectively, those marked c have a [010] Burgers' vectors. Diffracting vectors as shown, beam directions are (a) and (b) [111], (c) [113], (d) and (e) [001].
presumably resulting from the intersection of these dislocations. Similar dislocations configurations were also observed in Fe--40 A1 and Fe--45 AI, for example see Fig. 6. In addition to the "coathanger" networks commonly observed in Fe-34A1, discrete (001) dislocations were observed in both Fe--40 AI and Fe--45 A1, often with jogs apparent along their length. The incidence of these discrete (001) dislocations
was observed to increase with increasing aluminum content. The ( I 1 1 ) dislocations at the junctions of the (001) dislocations may not have been observed, because they were not present in the thin foil section. Selected area channeling patterns (SACP) have been used to determine the cleavage plane in transgranular fracture in intermetallic compounds, for example in Al3Ti-based alloys [24]. Attempts were
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MUNROE and BAKER: OBSERVATION OF (001> DISLOCATIONS a
III
X
D
Y III
Fig. 3. Bright field transmission electron micrograph of [0I0] dislocations in Fe--34 AI after 1% compression at room temperature, imaged with g = 020 and s ~ 0. Note the double image consistent with g'b = 2. B ~ [001]. made to investigate the plane or planes on which transgranular fracture occurred in Fe--34 AI, Fe--40 AI and Fe--45 AI by SACP using the JEOL 2000FX operating at 80 kV. N o channeling patterns could be obtained (although "high quality" patterns could be obtained from a silicon standard), indicating that sufficient localized plastic flow had occurred prior to fracture to blur any channeling pattern information. 4. DISCUSSION The (001) dislocations observed in these alloys appear to have arisen through the interaction of gliding (111) dislocations. Thus these (001) dislocations are different from (001) dislocations observed after high temperature deformation. As described above the temperature at which (001) slip
t
I' 1,2,:¸!
J
Fig. 5. (a) Bright field transmission electron micrograph of [010] dislocations, marked a and b, in Fe-34 AI after I% compression at room temperature, imaged with g = 500 so that g'b = 0. The [010] dislocations can still be observed due to residual contrast. At the junctions, marked X and Y, of the two [010] dislocations, (11 I) dislocations can be observed. B ~ [001]. (b) Weak-beam electron micrograph of [010] dislocations, marked a and b, in Fe-34AI after 1% compression at room temperature, g = 110, B~[001] s > 0.2 nm- 1. Note the jog, marked J on a. is normally observed in FeAI alloys increases with increasing iron content, however, (001) slip would not be expected even in the most aluminum-rich alloy, Fe--45 AI, at deformation temperatures less than ~ 8 0 0 K [7]. The production of (001) dislocations through the interaction of (111) dislocations is reminiscent of similar observations in b.c.c, metals. In a b.c.c, metal the dislocation reaction between two screw a / 2 ( l 11) dislocations to form an a ( 0 0 1 ) dislocation e.g.
Fig. 4. Bright field transmission electron micrograph of [010] dislocations in Fe--34 AI after 1% compression at room temperature, imaged with g = 0I 1 and s ~ 0. Note that the dynamical fringe contrast from each [010] dislocation lies on the same side of the dislocation, indicating that they each have the same sign. B ~ [I 11].
a/2 [111] + a/2 [TlI]--,a
[010]
on (OT1)
(001)
(011)
is energetically favorable. The reaction has been inferred previously in b.c.c, iron and iron-based alloys, after tensile straining over a range of temperatures, from the networks of (001) and a / 2 ( l l l > dislocations observed [25-29]. The resulting (001) dislocations have been shown to be either pure edge
MUNROE and BAKER: OBSERVATION OF (001) DISLOCATIONS
1015
which produces an equilibrium spacing between the two partials of ~ 5 nm [30]. The pairs of (001) dislocations appear to have formed through the interaction of two < 111) APBcoupled dislocation pairs. In the above example for Fe--34 AI, two a (111) dislocations are observed, specifically [111] and [TIT]. The interaction of a a / 2 ( l l l ) dislocation from each unit dislocation could produce an a(001) dislocation. Specifically
a/2 [I11] + a/2 [TlI]--,a [010]. A possible mechanism for the formation of the (001) dislocations through the intersection o f two APB-eoupled (111) dislocations is illustrated in Fig. 7. The leading a/2(l 11) dislocations react to produce an a(001) dislocation as indicated above [Fig. 7(b)], This is then pushed away from the intersection plane by the two unreaeted a/2(111) dislocations which continue to glide down their slip planes. The a/2(l 11) are pushed toward the junction by other gliding (111) dislocations on the slip plane and react to form a second (001> dislocation [Fig. 7(c)]. The two (001> dislocations would then try and repel, but both are pinned at each end at the junction of the interacting (111) dislocations. The second (001) dislocation appears to maintain an energetically favorable <100) line direction, but the first (001) dislocation is repelled. It experiences
Fig. 6. (a) Bright field transmission electron micrograph of Fe--40A1, after 1% compression at room temperature, A typical (111) dislocation is marked a, and a(001) dislocation is marked b. Diffraction vector as shown, B ~ [001]. (b) Bright field transmission electron micrograpli of dislocation structure in Fe--45AI, after 1% compressive strain at room temperature, typical (111) dislocations are marked a, and (001) dislocations marked b and c. (001) dislocations similar to those observed in Fe-34AI are labeled b. Discrete dislocations are labeled c. Diffraction vector as shown, B ~ [001].
7m) (ol~Ozt)
or pure screw in character, a<110) dislocations have also been observed, suggesting that a further reaction between two a (001) dislocations to produce a (01 I) dislocations had also occurred [25, 28] e.g. a [001] + a [010]--*a [011]. This indicates that the (001) dislocations are glissile in b.c.c, metals. Glissile (001) dislocations were also observed during in situ straining of a HSLA steel [29]. In B2 compounds (111) unit dislocations are split into two a/2(111) dislocations separted by a antiphase boundary (APB) [Fig. 7(a). The separation of these two partials increases with decreasing APB energy and, thus, for FeA! alloys with increasing iron-rich deviations from stoichiometry [19]. For example, for a ( l l l ) screw dislocation in Fe-34 AI, in the B2 state, the APB energy is ~ - l l 0 m J m -2, AM 39~5 ~ S
Fig. 7. Schematic illustration of the formation of an a<001) dislocation. (a) Two (111 ) dislocations, each consisting of two a / 2 < l l l ) dislocations separated by an APB glide on
intersecting {011} planes. (b) The leading a / 2 ( l i l ) dislocations interact to form an a(001) dislocation, it is still connected to the trailing a/2
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MUNROE and BAKER: OBSERVATION OF (001) DISLOCATIONS
competing forces, the repulsion from the other (001) dislocation, but two straightening forces, one making the dislocation take up a low energy (001) line direction and another minimizing the line length. The above reactions may also be considered energetically. For simplicity only isotropic elasticity will be considered (i.e. the dislocation self energy is Gb2/4n for a screw and Gb2/4n(1 - v ) for an edge, where G is the shear modulus, b is the Burgers' vector and v is Poisson's ratio) and the interaction energy between the dislocations will be ignored. Before collision, [Fig. 7(a)], the energy is from two pairs of APB-coupled a / 2 ( l l l ) screw dislocations. If the APB energy per unit length of dislocation is/~, then the energy per unit length of dislocation is 3Ga 2
~
4n
+2//.
In the intermediate condition [Fig. 7(b)], when two APBs, two a / 2 ( l 11) dislocations and one edge (001 dislocation are present, the energy per unit length of dislocation is 3 Ga 2
4n
Ga 2
-~ - - + 4~z(l - - v)
28.
Taking v = 1/3, the energy per unit length of dislocation is 3Ga 2 - -
4n
+28.
That is, there is no net gain or reduction in energy. The final product of two a ( 0 0 1 ) edge dislocations has an energy of Ga 2
3Ga 2
2"4n(1 - v-----~= 4n which is a net reduction in energy. The difference between the reactants and the products being twice the APB energy 8. That is, this reaction is more energetically favorable as the composition of the FeAI becomes less iron-rich, since the APB energy between the I/2(111) partials increasest [30]. This is consistent with the observation of more (001) dislocations as the aluminum concentration increases. It was suggested by Cottrell that the interaction between two a/2(111) dislocation to produce a ( 0 0 D edge dislocation may act as a crack nucleus in b.c.c, metals [31]. It is possible, therefore, that the interaction of APB-coupled (111) dislocations to produce (001) edge dislocations in FeA1 lead to transgranular cleavage on {001} planes. Although it was not possible to obtain the fracture plane in FeAI
tThe APB energy has only been measured for FeAI for aluminum contents between 25 and 36% [30]. Here APB energy increases linearly with increasing aluminum content. The APB energy has not been measured for Fe--40AI and Fe--45 AI, but by linear extrapolation of the data of Crawford and Ray it is deemed to increase with increasing aluminum content.
using SACP's (further studies are in progress), it is pertinent to note that {001} cleavage has been observed in other B2 compounds, notably//-CuZnbased alloys, which exhibit (111 ) slip, when tested in either water or liquid metals [32, 33]. There are no other reports, to the authors' knowledge, of the fracture plane in B2 compounds. The lower number of <001) dislocations in the most iron-rich composition is consistent with the observation of increasing ductility as the compositions become more ironrich. It should be noted that stoichiometric FeAI fractures intergranularly, although the reasons for intergranular fracture in this alloy are not understood. It is, therefore, possible that through alloying additions which decrease the APB energy of FeAI, the incidence of (001) dislocations would be reduced with a resulting improvement in the ductility of iron-rich FeAI. 5. CONCLUSIONS Transmission electron microscopy of iron-rich FeAI alloys deformed in compression at room temperature reveal the presence of <001) dislocations, in addition to (111) dislocations. It appears that the (001 ) dislocations form through an energetically favorable reaction between two (111) dislocations. Transgranular fracture in FeAI is accomplished by significant local plastic strain. A mechanism is outlined which may produce transgranuiar cleavage on {001} planes. Acknowledgements--The authors gratefully acknowledge Dr F. D. Lemkey of the United Technologies Research Center for providing the ingots and Mr D. J. Gaydosh and Dr J. D. Whittenberger of the NASA Lewis Research Center for extruding them. F. Liu and P. Nagpal are acknowledged for their assistance with specimen preparation. This work is supported by the Office of Basic Energy Sciences of the U.S. Department of Energy through grant DE-FG02-87ER45311. The use of the Dartmouth College Electron Microscope Center is also acknowledged. REFERENCES
I. I. L. F. Ray, R. C. Crawford and D. J. H. Cockayne, Phil. Mag. 21, 1027 (1970). 2. R. C. Crawford, Phil. Mag. 33, 529 (1976). 3. T. Yamagata and H. Yoshida, Mater. Sci. Engng 12, 95 (1973). 4. T. Yamagata, Trans. Japan Inst. metall. 18, 715 (1977). 5. Y. Umakoshi and M. Yamaguehi, Phil. Mag. A 41, 573 (1980). 6. Y. Umakoshi and M. Yamaguchi, Phil. Mag. A 44, 711 (1981). 7. M.G. Mendiratta, H-M. Kim and H. A. Lipsitt, Metall. Trans. A 15A, 395 (1984). 8. C. T. Chou and P. B. Hirsch, Phil. Mag. A 44, 1415 (1981). 9. C. T. Chou and P. B. Hirsch, Inst. Phys. Conf. Ser. 61, 459 (1982). 10. C. T. Chou and P. B. Hirsch, Proe. R. Soc. A387, 91 (1983). 1I. P. M. Hazzeldine and P. B. Hirsch, Proc. M R S 81, 75 (1987).
MUNROE and BAKER:
OBSERVATION OF (001) DISLOCATIONS
12. I. Baker and D. J. Gaydosb, Proc. MRS 81, 315 (1987). 13. I. Baker and D. J. Gaydosh, Mater. Sci. Engng 96, 147 (1987). 14. P. R. Munroe and I. Baker, J. Mater. Sci. 24, 4246 (1989). 15. C. T. Liu, E. H. Lee and C. G. McKamey, Scripta metall. 23, 875 (1989). 16. P. Nagpal and I. Baker, Metall. Trans..4. 21A, 2281 (1990). 17. M. G. Mendiratta, S. K. Ehlers and D. K. Chaterjee, in Rapid Solidfication Processing--Materials and Technologies llI (edited by R. Mehrabian), p. 240. NBS, Gaithersburg, Md. 18. F. Liu and I. Baker, unpublished results (1989). 19. Binary Phase Alloy diagrams (edited by T. B. Massalski), p. 112. Am. Soc. Metals, Metals Park, Ohio. 20. I. Baker and D. J. Gaydosh, Physica status solidi (a) 96, 185 (1986). 21. P. B. Hirsch, A. Howie, R. B. Nicholson, D. W. Pashley and M. J. Whelan, Electron Microscopy of Thin Crystals. R. E. Kreiger, Malabar FA (1965).
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22. H. J. Leamy, E. D. Gibson and F. X. Kayser, ,4cta metall. 15, 1827 (1967). 23. A. K. Head, Physiea status solidi 19, 185 (1967). 24. E. P. George, W. D. Porter, H. M. Henson, W. C. Oliver and B. F. Oliver, J. Mater. Res. 4, 78 (1989). 25. W. Carrington, K. F. Hale and D. McClean, Proc. R. Soc. A259, 203 (1960). 26. S, M. Ohr and D. N. Beshers, Phil. Mag. 10, 219 (1964). 27. A. S. Keh, Phil. Mag. 12, 9 (1965). 28. D. J. Dingley and K. F. Hale, Proc. R. Soc. A295, 55 (1966). 29. A. W. Sleeswyk and S. Mandziej, Scripta metall, mater. 24, 7 (1990). 30. R. C. Crawford and I. L. F. Ray, Phil. Mag. 35, 549 (1977). 31. A. H. Cottrell, T.M.S.-,4.LM.E. 212, 192 (1958). 32. W. K. Blancbard and L. A. Heldt, Metall. Trans..4 20A, 1439 (1989). 33. M. M. Shea and N. S. Stoloff, Mater. Sci. Engng 12, 245 (1973).
0956-7151/91 $3.00 + 0,00 Copyright © 1991 Pergamon Press plc
Printed in Great Britain. All rights reserved
OBSERVATION OF (001) DISLOCATIONS AND A MECHANISM FOR TRANSGRANULAR FRACTURE ON {001} IN FeA1 P. R. MUNROE and I. BAKER Thayer School of Engineering, Dartmouth College, Hanover, NH 03755, U.S.A.
(Received 13 February 1990; receivedfor publication 12 November 1990) Alam'aet--Transmission electron microscopy of iron-rich, B2-structured, FeA1 alloys which had been compressed at room temperature revealed the presence of (001) dislocations, in addition to APB-coupled <111) dislocations which are normally observed at room temperature. It appears that the (001) dislocations form through the interaction of the (11 I) dislocations. This reaction to produce <001) dislocations is outlined and shown to be energetically favorable. The production of these (001) dislocations on {100} planes is suggested to be a cause of transgranular cleavage on {100} in B2 alloys deforming by < I 11) slip. Rtmmt---La microscopie 61ectronique en transmission d'alliages FeA1 (de structure B2) riches en fer, comprim~s a la teml~rature ambiante, met en 6vidence la presence de dislocations (001), en plus des dislocations (I 11), coupl6es aux parois d'antiphase, que l'on observe en principe :i la temp6rature ambiante. II semble que ces dislocations se forment ~i partir de rinteraction des dislocations (I 11). On pr~sente rapidement cette r~action et l'on montre qu'eUe est 6nerg6tiquement favorable. On sugg~re que la creation de ces dislocations (00 ! ) darts les plans {100} est une cause du clivage transgranutaire suivant les plans {I00} dam les alliages B2 qui se d6forment par le glissement (111). Zusammenfassung~In eisenreichen FeA1-Legierungen mit B2-Struktur, die bei Raumtemperatur druckvefformt worden find, werden elektronenmikroskopisch (001)-Versetzungen beobachtet. Diese treten zus/itzlich zu den normalerweise bei Raumtemperatur beobachteten, mit Antiphasengrenzen behafteten (111 )-Versetzungen auf. Es scheint, dab die (001)-Versetzungen aus einer Reaktion yon (111 )-Versetzungen entstehen. Diese Reaktion wird dargestellt; sie ist energetisch bevorzugt. Die Entstehung dieser (001)-Versetzungen auf {111}-Ebenen wird als eine Ursache der transgranularen Spaltung auf {I 1l}-Ebenen in B2-Legierungen, die sich mit (11 l)-Gleitung verformen, angesehen.
1. INTRODUCTION FeAI is a B2-structured compound which exists over a wide range of composition (34-51 at.% AI). Previous studies have shown that FeA1 deforms by (111 ) slip at all compositions at room temperature [1-14], and by (001) slip at elevated temperatures [5-7, 12-14], the transition temperature from (111) to (001) slip increasing with increasing iron-rich deviations from stoichiometry [7, 14]. A significant difference between the two slip vectors is that whereas 111 ) slip provides sufficient slip systems to satisfy Von Mises criterion for plastic flow in polycrystals <001) slip does not. The ductility of FeAI varies with composition: the ductility can also be affected by several extrinsic factors such as the testing environment [15] and the concentration of vacancies retained after heat treatment [16]. Stoichiometric FeAI is brittle at room temperature, generally fracturing before yielding in tensile tests [17], but the ductility increases as FeAI becomes increasingly iron-rich, for example, Fe-36.5 at.% A1 can exhibit as much as 12% tensile elongation [15]. The fracture mode of FeAI at room
temperature is affected by both grain size and composition [18]. Stoichiometric FeA1 exhibits predominantly intergranular fracture, but the incidence of transgranular fracture increases with both increasing iron-rich deviation from stoichiometry and increasing grain size [18]. This paper reports upon the observation of (001) dislocations in FeA1 after room temperature compression. These (001) dislocations are formed through the interaction of gliding < 111 ) dislocations. A mechanism for transgranular fracture based upon the occurrence of these <001) dislocations is outlined. 2. EXPERIMENTAL Three FeAl alloys, with compositions of F¢-34 A1, F¢-40 Al and F¢--45 A1 (compositions in at.%), were studied. The processing routes of these alloys have been described elsewhere [14]. Cylinders ~ 10ram long by 3 mm diameter were annealed at 1123 K for one hour (this removed the residual dislocations resulting from prior processing) and water-quenched. The Fe--40Al and Fe--45Al samples were then
1011
1012
MUNROE and BAKER: OBSERVATION OF (001) DISLOCATIONS
annealed at 673 K for 120h to reduce the high vacancy concentrations which are retained after annealing at temperatures above ~ I 0 0 0 K [16]. Fe-34 A1 can exist in both the B2 and DO3 states, the DO3 state being the equilibrium phase below ~ 700 K [19]. Thus, the Fe-34 AI sample was not annealed to reduce the vacancy concentration. So that the B2 microstructure would be retained, it was compressed after quenching alone (vacancy retention appears to be less pronounced in Fe-34 AI [16]). The cylinders were then strained ~ 1% in compression at a strain rate of ~ 1 0 - ~ s -~, at room temperature. Thin foils were prepared from all the alloys as described elsewhere [20], and examined in a JEOL 2000FX transmission electron microscope (TEM) operating at 200 kV. Dislocation analyses were performed using the g.b. = 0 invisibility criterion [21]. Samples of Fc-34AI, Fc-40Al and Fe-45Al were fractured in bending at both room temperatures and liquid nitogen temperatures and the fracture surfaces were examined using selected area electron channeling in the J E O L 2000FX operated in S E M mode. 3. RESULTS The microstructures of all the alloys were examined prior to compression and consisted of equiaxed, dislocation-free grains ~ 50 # m in diameter. Selected area difffraction patterns from Fe-34 AI indicated that water quenching from 1123K had retained the B2 structure (Fig. 1), (111) superlattice reflections from the DO3 structure were extremely diffuse, indicating that the DO3 phase had been largely suppressed by water quenching. The dislocation substructure of Fe--34 AI after compression is shown in Fig. 2. The dislocations marked a show either invisibility or residual contrastt when g = 110 and g = 10I, indicating that b = l l I , whilst the dislocations marked b show either invisibility or weak contrast when g = 110, g = 01T and g = 10I, indicating that b = 111. Trace analysis was used to show that these (111 ) dislocations are screw in character. In addition, dislocations marked c were observed, which are invisible when g = 200 and g = 10I, indicating a [010] Burgers' vector. The Burgers' vector must be a [010], rather than a/2[010], since no fault contrast was observed between the dislocations. This was also confirmed by imaging the dislocations with g = 020 and s -- 0, when a double image was observed (Fig. 3), consistent with g. b = 2 [21].Trace analysis showed that the predominant general directionof the
tFeA1 is strongly anisotropic. The Zoner anisotropy factor, A, of Fv--34AI is 4.4, and of F¢-40 AI is 3.8 [22], and there is often residual contrast even when a screw dislocation is viewed under conditions which satisfy the g" b = 0 invisibility criterion.
C)©?
Fig. 1. Selected area diffraction pattern from (110) zone axis in F¢-34 A1, heat treated at 1123 K for I h and waterquenched. Note the absence of DO3(III } superlattice reflections, which suggests that the DO 3 structure has been suppressed.
dislocation line was [100]. Tilting experiments which examined how the dislocation separation changed with the beam normal showed that the [010] dislocations lay on the (001) plane. A (100) line direction has been shown to be a stable orientation for (001) { 100} dislocations in anisotropic B2 compounds [23]. To investigate whether the [010] dislocations were part of a loop or a dipole, ((111 ) dislocation dipoles are commonly observed in FeAI alloys [14]), they were imaged using both + g and - g two-beam conditions. If they were dislocation loops or dipoles, a change in the spacing between the two dislocations would be expected on reversing g [21]. No change in spacing was observed, indicating that the Burgers' vectors of the two dislocations are the same. In addition, dynamical fringe contrast from each dislocation, when s -- 0, lies on the same side of the defect also indicating that each dislocation has the same sign (Fig. 4). Figure 5 shows images of a relatively isolated dislocation network similar to that described as c in Fig. 2, within the same grain as the dislocations described above. At each of the junctions of the two [010] dislocations ( I 11 ) dislocations can be observed. At the junction marked X there is both a [I11] and a [ I l l ] dislocation present, at the junction marked Y only a [111] dislocation can be observed. It was assumed that the [ I l l ] dislocation at the junction Y was not present in the foil. These can be most clearly seen when g = 200 and the [010] dislocations are out of contrast [Fig. 5(a)]. One of the [010] dislocations in Fig. 5, marked a, is relatively straight with a [100] line direction, whereas the part marked b initially curves away from the two junctions but then straightens to take up a line direction close to [100] away from the junctions, giving the defect a "coathanger" shape. Jogs can also be seen along the length of the [010] dislocations, marked J [Fig. 5(b)],
MUNROE and BAKER: OBSERVATION OF (001) DISLOCATIONS
1013
e
i
/
Fig. 2. Bright field transmission electron micrograph of dislocation structure in Fe--34AI after 1% compression at room temperature. The dislocations marked a and b have [I 1I] and [111] Burgers' vectors respectively, those marked c have a [010] Burgers' vectors. Diffracting vectors as shown, beam directions are (a) and (b) [111], (c) [113], (d) and (e) [001].
presumably resulting from the intersection of these dislocations. Similar dislocations configurations were also observed in Fe--40 A1 and Fe--45 AI, for example see Fig. 6. In addition to the "coathanger" networks commonly observed in Fe-34A1, discrete (001) dislocations were observed in both Fe--40 AI and Fe--45 A1, often with jogs apparent along their length. The incidence of these discrete (001) dislocations
was observed to increase with increasing aluminum content. The ( I 1 1 ) dislocations at the junctions of the (001) dislocations may not have been observed, because they were not present in the thin foil section. Selected area channeling patterns (SACP) have been used to determine the cleavage plane in transgranular fracture in intermetallic compounds, for example in Al3Ti-based alloys [24]. Attempts were
1014
MUNROE and BAKER: OBSERVATION OF (001> DISLOCATIONS a
III
X
D
Y III
Fig. 3. Bright field transmission electron micrograph of [0I0] dislocations in Fe--34 AI after 1% compression at room temperature, imaged with g = 020 and s ~ 0. Note the double image consistent with g'b = 2. B ~ [001]. made to investigate the plane or planes on which transgranular fracture occurred in Fe--34 AI, Fe--40 AI and Fe--45 AI by SACP using the JEOL 2000FX operating at 80 kV. N o channeling patterns could be obtained (although "high quality" patterns could be obtained from a silicon standard), indicating that sufficient localized plastic flow had occurred prior to fracture to blur any channeling pattern information. 4. DISCUSSION The (001) dislocations observed in these alloys appear to have arisen through the interaction of gliding (111) dislocations. Thus these (001) dislocations are different from (001) dislocations observed after high temperature deformation. As described above the temperature at which (001) slip
t
I' 1,2,:¸!
J
Fig. 5. (a) Bright field transmission electron micrograph of [010] dislocations, marked a and b, in Fe-34 AI after I% compression at room temperature, imaged with g = 500 so that g'b = 0. The [010] dislocations can still be observed due to residual contrast. At the junctions, marked X and Y, of the two [010] dislocations, (11 I) dislocations can be observed. B ~ [001]. (b) Weak-beam electron micrograph of [010] dislocations, marked a and b, in Fe-34AI after 1% compression at room temperature, g = 110, B~[001] s > 0.2 nm- 1. Note the jog, marked J on a. is normally observed in FeAI alloys increases with increasing iron content, however, (001) slip would not be expected even in the most aluminum-rich alloy, Fe--45 AI, at deformation temperatures less than ~ 8 0 0 K [7]. The production of (001) dislocations through the interaction of (111) dislocations is reminiscent of similar observations in b.c.c, metals. In a b.c.c, metal the dislocation reaction between two screw a / 2 ( l 11) dislocations to form an a ( 0 0 1 ) dislocation e.g.
Fig. 4. Bright field transmission electron micrograph of [010] dislocations in Fe--34 AI after 1% compression at room temperature, imaged with g = 0I 1 and s ~ 0. Note that the dynamical fringe contrast from each [010] dislocation lies on the same side of the dislocation, indicating that they each have the same sign. B ~ [I 11].
a/2 [111] + a/2 [TlI]--,a
[010]
on (OT1)
(001)
(011)
is energetically favorable. The reaction has been inferred previously in b.c.c, iron and iron-based alloys, after tensile straining over a range of temperatures, from the networks of (001) and a / 2 ( l l l > dislocations observed [25-29]. The resulting (001) dislocations have been shown to be either pure edge
MUNROE and BAKER: OBSERVATION OF (001) DISLOCATIONS
1015
which produces an equilibrium spacing between the two partials of ~ 5 nm [30]. The pairs of (001) dislocations appear to have formed through the interaction of two < 111) APBcoupled dislocation pairs. In the above example for Fe--34 AI, two a (111) dislocations are observed, specifically [111] and [TIT]. The interaction of a a / 2 ( l l l ) dislocation from each unit dislocation could produce an a(001) dislocation. Specifically
a/2 [I11] + a/2 [TlI]--,a [010]. A possible mechanism for the formation of the (001) dislocations through the intersection o f two APB-eoupled (111) dislocations is illustrated in Fig. 7. The leading a/2(l 11) dislocations react to produce an a(001) dislocation as indicated above [Fig. 7(b)], This is then pushed away from the intersection plane by the two unreaeted a/2(111) dislocations which continue to glide down their slip planes. The a/2(l 11) are pushed toward the junction by other gliding (111) dislocations on the slip plane and react to form a second (001> dislocation [Fig. 7(c)]. The two (001> dislocations would then try and repel, but both are pinned at each end at the junction of the interacting (111) dislocations. The second (001) dislocation appears to maintain an energetically favorable <100) line direction, but the first (001) dislocation is repelled. It experiences
Fig. 6. (a) Bright field transmission electron micrograph of Fe--40A1, after 1% compression at room temperature, A typical (111) dislocation is marked a, and a(001) dislocation is marked b. Diffraction vector as shown, B ~ [001]. (b) Bright field transmission electron micrograpli of dislocation structure in Fe--45AI, after 1% compressive strain at room temperature, typical (111) dislocations are marked a, and (001) dislocations marked b and c. (001) dislocations similar to those observed in Fe-34AI are labeled b. Discrete dislocations are labeled c. Diffraction vector as shown, B ~ [001].
7m) (ol~Ozt)
or pure screw in character, a<110) dislocations have also been observed, suggesting that a further reaction between two a (001) dislocations to produce a (01 I) dislocations had also occurred [25, 28] e.g. a [001] + a [010]--*a [011]. This indicates that the (001) dislocations are glissile in b.c.c, metals. Glissile (001) dislocations were also observed during in situ straining of a HSLA steel [29]. In B2 compounds (111) unit dislocations are split into two a/2(111) dislocations separted by a antiphase boundary (APB) [Fig. 7(a). The separation of these two partials increases with decreasing APB energy and, thus, for FeA! alloys with increasing iron-rich deviations from stoichiometry [19]. For example, for a ( l l l ) screw dislocation in Fe-34 AI, in the B2 state, the APB energy is ~ - l l 0 m J m -2, AM 39~5 ~ S
Fig. 7. Schematic illustration of the formation of an a<001) dislocation. (a) Two (111 ) dislocations, each consisting of two a / 2 < l l l ) dislocations separated by an APB glide on
intersecting {011} planes. (b) The leading a / 2 ( l i l ) dislocations interact to form an a(001) dislocation, it is still connected to the trailing a/2
1016
MUNROE and BAKER: OBSERVATION OF (001) DISLOCATIONS
competing forces, the repulsion from the other (001) dislocation, but two straightening forces, one making the dislocation take up a low energy (001) line direction and another minimizing the line length. The above reactions may also be considered energetically. For simplicity only isotropic elasticity will be considered (i.e. the dislocation self energy is Gb2/4n for a screw and Gb2/4n(1 - v ) for an edge, where G is the shear modulus, b is the Burgers' vector and v is Poisson's ratio) and the interaction energy between the dislocations will be ignored. Before collision, [Fig. 7(a)], the energy is from two pairs of APB-coupled a / 2 ( l l l ) screw dislocations. If the APB energy per unit length of dislocation is/~, then the energy per unit length of dislocation is 3Ga 2
~
4n
+2//.
In the intermediate condition [Fig. 7(b)], when two APBs, two a / 2 ( l 11) dislocations and one edge (001 dislocation are present, the energy per unit length of dislocation is 3 Ga 2
4n
Ga 2
-~ - - + 4~z(l - - v)
28.
Taking v = 1/3, the energy per unit length of dislocation is 3Ga 2 - -
4n
+28.
That is, there is no net gain or reduction in energy. The final product of two a ( 0 0 1 ) edge dislocations has an energy of Ga 2
3Ga 2
2"4n(1 - v-----~= 4n which is a net reduction in energy. The difference between the reactants and the products being twice the APB energy 8. That is, this reaction is more energetically favorable as the composition of the FeAI becomes less iron-rich, since the APB energy between the I/2(111) partials increasest [30]. This is consistent with the observation of more (001) dislocations as the aluminum concentration increases. It was suggested by Cottrell that the interaction between two a/2(111) dislocation to produce a ( 0 0 D edge dislocation may act as a crack nucleus in b.c.c, metals [31]. It is possible, therefore, that the interaction of APB-coupled (111) dislocations to produce (001) edge dislocations in FeA1 lead to transgranular cleavage on {001} planes. Although it was not possible to obtain the fracture plane in FeAI
tThe APB energy has only been measured for FeAI for aluminum contents between 25 and 36% [30]. Here APB energy increases linearly with increasing aluminum content. The APB energy has not been measured for Fe--40AI and Fe--45 AI, but by linear extrapolation of the data of Crawford and Ray it is deemed to increase with increasing aluminum content.
using SACP's (further studies are in progress), it is pertinent to note that {001} cleavage has been observed in other B2 compounds, notably//-CuZnbased alloys, which exhibit (111 ) slip, when tested in either water or liquid metals [32, 33]. There are no other reports, to the authors' knowledge, of the fracture plane in B2 compounds. The lower number of <001) dislocations in the most iron-rich composition is consistent with the observation of increasing ductility as the compositions become more ironrich. It should be noted that stoichiometric FeAI fractures intergranularly, although the reasons for intergranular fracture in this alloy are not understood. It is, therefore, possible that through alloying additions which decrease the APB energy of FeAI, the incidence of (001) dislocations would be reduced with a resulting improvement in the ductility of iron-rich FeAI. 5. CONCLUSIONS Transmission electron microscopy of iron-rich FeAI alloys deformed in compression at room temperature reveal the presence of <001) dislocations, in addition to (111) dislocations. It appears that the (001 ) dislocations form through an energetically favorable reaction between two (111) dislocations. Transgranular fracture in FeAI is accomplished by significant local plastic strain. A mechanism is outlined which may produce transgranuiar cleavage on {001} planes. Acknowledgements--The authors gratefully acknowledge Dr F. D. Lemkey of the United Technologies Research Center for providing the ingots and Mr D. J. Gaydosh and Dr J. D. Whittenberger of the NASA Lewis Research Center for extruding them. F. Liu and P. Nagpal are acknowledged for their assistance with specimen preparation. This work is supported by the Office of Basic Energy Sciences of the U.S. Department of Energy through grant DE-FG02-87ER45311. The use of the Dartmouth College Electron Microscope Center is also acknowledged. REFERENCES
I. I. L. F. Ray, R. C. Crawford and D. J. H. Cockayne, Phil. Mag. 21, 1027 (1970). 2. R. C. Crawford, Phil. Mag. 33, 529 (1976). 3. T. Yamagata and H. Yoshida, Mater. Sci. Engng 12, 95 (1973). 4. T. Yamagata, Trans. Japan Inst. metall. 18, 715 (1977). 5. Y. Umakoshi and M. Yamaguehi, Phil. Mag. A 41, 573 (1980). 6. Y. Umakoshi and M. Yamaguchi, Phil. Mag. A 44, 711 (1981). 7. M.G. Mendiratta, H-M. Kim and H. A. Lipsitt, Metall. Trans. A 15A, 395 (1984). 8. C. T. Chou and P. B. Hirsch, Phil. Mag. A 44, 1415 (1981). 9. C. T. Chou and P. B. Hirsch, Inst. Phys. Conf. Ser. 61, 459 (1982). 10. C. T. Chou and P. B. Hirsch, Proe. R. Soc. A387, 91 (1983). 1I. P. M. Hazzeldine and P. B. Hirsch, Proc. M R S 81, 75 (1987).
MUNROE and BAKER:
OBSERVATION OF (001) DISLOCATIONS
12. I. Baker and D. J. Gaydosb, Proc. MRS 81, 315 (1987). 13. I. Baker and D. J. Gaydosh, Mater. Sci. Engng 96, 147 (1987). 14. P. R. Munroe and I. Baker, J. Mater. Sci. 24, 4246 (1989). 15. C. T. Liu, E. H. Lee and C. G. McKamey, Scripta metall. 23, 875 (1989). 16. P. Nagpal and I. Baker, Metall. Trans..4. 21A, 2281 (1990). 17. M. G. Mendiratta, S. K. Ehlers and D. K. Chaterjee, in Rapid Solidfication Processing--Materials and Technologies llI (edited by R. Mehrabian), p. 240. NBS, Gaithersburg, Md. 18. F. Liu and I. Baker, unpublished results (1989). 19. Binary Phase Alloy diagrams (edited by T. B. Massalski), p. 112. Am. Soc. Metals, Metals Park, Ohio. 20. I. Baker and D. J. Gaydosh, Physica status solidi (a) 96, 185 (1986). 21. P. B. Hirsch, A. Howie, R. B. Nicholson, D. W. Pashley and M. J. Whelan, Electron Microscopy of Thin Crystals. R. E. Kreiger, Malabar FA (1965).
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22. H. J. Leamy, E. D. Gibson and F. X. Kayser, ,4cta metall. 15, 1827 (1967). 23. A. K. Head, Physiea status solidi 19, 185 (1967). 24. E. P. George, W. D. Porter, H. M. Henson, W. C. Oliver and B. F. Oliver, J. Mater. Res. 4, 78 (1989). 25. W. Carrington, K. F. Hale and D. McClean, Proc. R. Soc. A259, 203 (1960). 26. S, M. Ohr and D. N. Beshers, Phil. Mag. 10, 219 (1964). 27. A. S. Keh, Phil. Mag. 12, 9 (1965). 28. D. J. Dingley and K. F. Hale, Proc. R. Soc. A295, 55 (1966). 29. A. W. Sleeswyk and S. Mandziej, Scripta metall, mater. 24, 7 (1990). 30. R. C. Crawford and I. L. F. Ray, Phil. Mag. 35, 549 (1977). 31. A. H. Cottrell, T.M.S.-,4.LM.E. 212, 192 (1958). 32. W. K. Blancbard and L. A. Heldt, Metall. Trans..4 20A, 1439 (1989). 33. M. M. Shea and N. S. Stoloff, Mater. Sci. Engng 12, 245 (1973).