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Structure and development of slip lines during plastic deformation of the intermetallic phases Fe3A1 and CuZn
Materials Science and Engineering A239 – 240 (1997) 180 – 187
Structure and development of slip lines during plastic deformation of the intermetallic phases Fe3A1 and CuZn A. Brinck *, C. Engelke, W. Kopmann, H. Neuha¨user Institut fu¨r Metallphysik und Nukleare Festko¨rperphysik, TU Braunschweig, Mendelssohnstr. 3, D-38106 Braunschweig, Germany
Abstract The development and fine structure of slip lines on Fe3Al and b-CuZn single crystals has been investigated by video recording through a far distance optical microscope during compressive deformation and by scanning force microscopy (SFM) after deformation at room temperature up to temperatures beyond the yield stress peak. Different deformation modes, depending on temperature, are identified in Fe3A1 crystals and similarly found in b-CuZn. Homogeneous slip occurs at room temperature and is superposed on heterogeneous localized slip connected with stress serrations in a range of intermediate temperatures up to the yield stress peak, where the type of activated slip systems changes. These variations appear to be connected with the generation and motion of imperfect superdislocations at low and of perfect ones at intermediate temperatures, which, however, are gradually uncoupled with increasing temperature by an antiphase relaxation (local re-ordering) process, until beyond the peak dislocations of another slip system become more favourable. © 1997 Elsevier Science S.A. Keywords: Slip lines; Plastic deformation; Intermetallics; Fe3Al; CuZn; AFM
1. Introduction
2. Experimental details
The analysis of slip lines produced by the movement of dislocations emerging at the crystal surface, provides an easy way to investigate the distribution and activity of dislocation sources, as well as the type of activated slip systems, e.g. for DO3-ordered Fe3Al [1–4] and B2-ordered CuZn [5 – 8]. In the following, we present a systematic investigation on slip line formation and fine structure in these materials at various temperatures, from records taken through an optical microscope during in-situ deformation. The fine structure of the rather faint lines with low step heights is resolved and measured by means of scanning force microscopy. The observed changes of the characteristics of slip line development and structure are compared for the two types of intermetallic phases and are correlated with the simultaneously measured temperature dependence of the flow stress in view of the basic dislocation mechanisms.
Single crystals were produced after induction melting of the pure metals by the Bridgman method in resistance furnaces; Fe–24.8at.%Al crystals in Al2O3 powder crucibles, to avoid any carbon, and Cu–47.5at.%Zn crystals in pure graphite crucibles. Crystals mostly orientated for single slip were selected and prepared for compression tests (Fe3A1 8× 3× 3 mm3, b-CuZn diameter 3.9 mm, gauge length 8 mm) and for tensile tests (b-CuZn diameter 3.9 mm, length 80 mm). After homogenization by annealing for 24 h at 1275 K (Fe3Al) and 775 K (CuZn), respectively, the crystals were cooled down slowly to achieve DO3 order for Fe3Al (cooling rate 0.4 K min − 1) and B2 order for CuZn (cooling rate 1.7 K min − 1). A smooth surface for slip line observations was produced by electrolytic polishing. During compression within a vacuum vessel at about 10 − 2 Pa, with a deformation rate of e; = 8× 10 − 5 s − 1, the specimens were heated from room temperature to 875 K by focusing infrared lamps. During straining, one surface of the crystals was imaged through a quartz window by means of an optical microscope with a working distance of 150 mm providing 2 mm resolution
* Corresponding author. 0921-5093/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved. PII S 0 9 2 1 - 5 0 9 3 ( 9 7 ) 0 0 5 7 9 - 0
A. Brinck et al. / Materials Science and Engineering A239–240 (1997) 180–187
(QUESTAR QM-100), using either bright or dark field illumination. The development of slip lines was recorded on magnetic tape with an attached CCD camera. The tensile tests on b-CuZn were performed in an Instron 1185 machine equipped with a heating chamber under a protective atmosphere of Ar. After deformation the crystal surface was examined by SFM (RASTERSCOPE 4000, DME) using Si cantilevers with an Si tip (NANOSCOPE, cone angle 35°). This method enabled fine slip lines to be resolved in these bcc materials, contrary to EM replica. All investigations were restricted to the first stages of deformation, extending up to a few percent of resolved strain.
3. Slip line formation and structure
3.1. Dependence of the CRSS of temperature Fig. 1a and b compare the dependencies of the critical resolved shear stress (crss) t0 and the strain
181
rate sensitivity S= (1/kT) (t/( ln e; on temperature for the two intermetallics Fe3Al and b-CuZn, in the range from room temperature up to temperatures beyond the so-called yield stress anomaly, where t0 increases with T. Three ranges can be differentiated in Fig. 1a for Fe3Al: A range I of high yield stress around room temperature with smooth load-elongation curves, a range II with relatively low yield stresses where the load-elongation curve exhibits serrations (for small deformation) increasing with T, in particular in the ‘anomalous’ part, before the maximum of the yield stress and a range III beyond this maximum, where the yield stress, decreases again and the load-elongation curves are smooth again, showing however, pronounced yield points at first loading. For b-CuZn, range I is ending already around 215 K [26], therefore Fig. 1b only shows ranges II and III with the same characteristics as for Fe3Al. The results shown in Fig. 1a agree reasonably with measurements in literature, considering the distinct dependencies on the crystal orientations [3] and on the exact alloy composition [9]. For example, for Fe3Al it is known that small amounts of carbon cause formation of carbides [10,11] which increase the flow stress (cf. [3]). The rather sharp drop of t0, seen in Fig. 1a for Fe3Al slightly above room temperature, is attributed to a transition from nucleation and glide of imperfect to perfect superdislocations (cf. [12] and Section 4). The crss values for b-CuZn in Fig. 1b are taken from tensile tests on long crystals, which exhibit a considerable concentration gradient, i.e. they correspond to the softest part of the crystals where, according to ICO-OES analysis, the composition is Cu–47(90.5)at.%Zn. The measured values are in reasonable agreement with literature [13,14]. For short compression specimens cut from parts of the long crystals we observed hardening effects by a- or g-precipitates, when the Zn concentration reached the limits of the b-phase.
3.2. Structure and de6elopment of slip lines
Fig. 1. Dependence of the crss t0 and of the strain rate sensitivity S =(1/kT) (t/( ln e; on deformation temperature T for (a) Fe3Al and (b) CuZn single crystals, oriented for single slip (characteristic ranges are indicated by I, II, III).
a) For Fe3Al, five regimes with different modes of slip line structure could be identified within the temperature ranges shown in Fig. 1a ([12]): (1) In crystals oriented for maximum shear stress on the {110} planes homogeneous slip occurs in faint propagating deformation bands at temperatures up to 325 K on the {110}111 system. The bands can be observed in LM bright field phase contrast (slight defocusing) and appear as a propagating ‘curtain’ on the video, but cannot be resolved properly on single video prints. Fig. 2 shows an example of such a deformation band front taken by SFM. The propagation rate 6prop is around several 10 mm s − 1,
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Fig. 2. Front region of a homogeneous deformation band propagating in range I (Fig. 1a) on a Fe3Al crystal oriented for single slip deformed in compression at room temperature (the wavy surface is the result of electropolishing), showing slip step heights, clustered with less than 2b (b: Burgers vector of a perfect superdislocation) per 103 neighbouring crystallographic slip planes.
corresponding to a local strain rate e; loc :5e; ext, where e; ext =8× 10 − 5 s − 1 is the externally imposed strain rate. (2) In crystals oriented for a maximum shear stress on the {211} planes, localized slip bands with rather high step heights (up to 500 nm) appear during loading and disappear during unloading (‘pseudoelastic effect’ [15]), at high stresses around room temperature and up to temperatures depending on predeformation. Fig. 3a shows LM micrographs (dark field illumination) taken during loading and unloading, Fig. 3b shows an SFM micrograph taken after several loading cycles on one of the few remaining slip band clusters. For more details cf. [16]. Here, the growth rate of slip band step height is S: :5 ×10 − 3 mm s − 1 (at the same e; ext as before), corresponding to a local strain rate e; loc of some 10e; ext. (3) In range II for both orientations at low stresses, inhomogeneous slip occurs on {110}111 in distinct and localized deformation bands, which formed rapidly (Fig. 4a,b). They are, together with static and dynamic strain ageing phenomena [17], connected with serrations in the load-elongation curve for T\ 475 K. Often the homogeneous slip mode (1) continues at both sides of a previously formed inhomogeneous deformation band (Fig. 4c). The rapid growth of step height is in the order of S: :30 mm s − 1 at T = 475 K, corresponding to a local strain rate of e; loc :750e; ext. This value increases with temperature in range II, as shown in Table 1 up to e; loc :4500e; ext. (4) While all slip lines shown before, (with the exception of the pseudoelastic deformation [16]), occur on the {110} slip planes, near the temperature of the
flow stress peak, also slip planes of type {112} and {321} are activated (Fig. 5). Here slip is stable again with a local strain rate of e; loc : 20e; ext. (5) For T exceeding the flow stress peak temperature, slip starts to occur in a very homogeneous manner on the {112} system, with step heights as faint as in mode (1), but no longer in propagating but in randomly appearing deformation bands. Owing to oxidation effects of the surface, the fine structure can no longer be resolved by SFM. Here the local strain rate is e; loc : e; ext. b) For b-CuZn, the slip line characteristics show close similarities to those mentioned above. (1) Around room temperature homogeneous slip occurs in faint deformation band fronts (Fig. 6), which propagate during the micro-yield region on the slip system {110}111, similar to a)(1) in Fe3Al. The front with single slip step heights of S= 0.5–5 nm (near the limit of resolution due to problems with electrolytic polishing) propagates in occasional bursts correlated with slight load drops (Fig. 7). The propagation rate 6prop is about 40–100 mm s − 1 in the average, but reaches several mm s − 1 for short time intervals. (2) After micro-yield at room temperature, faint slip line bundles develop in random distribution on the crystal surface, for crystals oriented for single glide always on the {011}111 system. (3) With increasing temperature, the slip lines get distinctly clustered with step heights of slip lines reaching S= 50–100 nm and undeformed regions between them (Fig. 8). They appear rapidly and are connected with serrations observed in CuZn, too, in
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Fig. 3. Slip bands during pseudoelastic deformation of a Fe3Al crystal oriented for maximum shear stress on {112} planes. (a) Series of video frames (optical micrographs) with slip bands in dark field illumination during loading (1. row up to 2. frame in 2. row) and unloading (succeeding frames). Time inset gives h:min:s:s/100; (b) SFM micrograph (top view and 3D-view) a few slip bands remaining after several unloading/loading cycles. The clustering amounts to about 720 partial dislocation Burgers vectors bp ( =b/4 with b: Burgers vector of perfect superdislocation) per 103 slip planes.
the temperature range corresponding to range II in Fig. 1. (4) From the literature [6,7,18 – 20], we take the result that, for temperatures exceeding that of the yield stress maximum in Fig. 1b, the mode of slip changes from {011}111 to the {011}112 slip system. Here the deformation is homogeneous and slip lines are not resolved.
4. Discussion Comparing the results for Fe3A1 and b-CuZn intermetallic phases, we find the following common features: 1. Homogeneous glide at low T, with slowly propagating fronts. 2. Clustered glide at intermediate T, rapidly forming slip bands connected with serrations in the yield
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Fig. 4. Localized clustered slip bands in Fe3Al appearing rapidly during deformation in range II (Fig. 1a) at T= 475 K in Fe3Al. (a) Successive video frames (optical micrographs) with indicated time scale. (b) SFM micrograph of a typical section. (c) SFM micrograph of a localized slip band with homogeneous continuation at both sides. (d) Step profile of the slip band shown in c. The clustering of the steep part is about 25 perfect superdislocation Burgers vectors per 103 slip planes.
curve; flow stress increasing with temperature up to a stress maximum at the peak temperature Tp. 3. Change of the active slip system around Tp, homogeneous glide for higher T \Tp. Point (1) indicates that, starting from few first activated dislocation sources, the dislocations easily crossslip to neighbouring slip planes [21]. The high stress acting in this range of temperature for Fe3A1 suggests that they are imperfect superdislocations [23], i.e. they are trailing an APB behind them, as confirmed in TEM [3,22]. The high stress, on the other hand, helps in cross-slip of the split screw dislocations [24,25]. With some increase of temperature, when the stress drops to distinctly lower values, the production of perfect superdislocations with narrow APBs between them seems possible [23,12] (for Fe3Al oriented for single glide at T \325 K, for CuZn already below room temperature [26]). The low external stress is able to activate only few favourably situated and configurated sources where some overstress is acting; then each source is able to generate a bunch of dislocations resulting in clustered slip. This clustering and the rapid dislocation generation appear to be en-
hanced with increasing T by local re-ordering processes in the APBs [27–30,19], whose effectivity increases with decreasing dislocation velocity and which cause pinning, breakaway and obstacle destruction effects, similar to short range order hardening [31,29]. As dislocation movement requires lower stress than dislocation production in this region, the slip bands develop so rapidly that some load relaxation of the machine-specimen system occurs and serrations on the load-elongation curve appear. The re-ordering effect (‘APB relaxation’ [32], considered in some detail for b-CuZn in [30]) increases with raising temperature, resulting in an increase of the measured flow stress because the dislocations get more and more uncoupled and move as imperfect dislocations trailing again, an APB behind them, as observed in TEM (CuZn: [20], Fe3Al [22]). This increase would continue if the crystal could not activate dislocations on another slip system which move at lower stress, such as more strongly coupled superdislocations, i.e. {112}111 in Fe3Al, {011}112 in b-CuZn (cf. point (3)).
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Table 1 Variation with temperature of slip modes, of activated types of slip planes, of clustering of slipa and of local strain rates in the active slip bandsb T [K]
Slip mode
Slip plane
Clustering [b/103 d]
e; loc [10−5 s−1]
300 300
Reversible Stable
{112} {110}
5180 B2
\100 40
475 575 675 775
Instable Instable Instable Instable
{110} {110} {110} {110}
25 25 25 25
6000 7500 10 000 38 000
825
Stable
{110}, {211}, {312}
60
160
875
Stable
{211}
a b
None
8
Given in perfect superdislocation Burgers vectors per 103 slip plane distances d. For an external strain rate of e; ext = 8×10−5 s−1 for Fe3Al single crystals.
As the cross-slip activity does not change markedly with T [22], the above re-ordering processes seem to be more probable for explaining the flow stress anomaly in these systems than the cross-slip locking process ([20,25]). The re-ordering mechanism also provides a straight-forward explanation for the static and dynamic strain-ageing phenomena observed in these two alloys [32,17,14,12] and recently modeled phenomenologically by [33]. Another common feature observed in all the alloys showing the flow stress maximum at elevated temperatures, i.e. the long-range ordered intermetallics Fe3Al, b-CuZn, but also short-range ordered Cu-based solid solutions, is the pronounced yield point on first loading of a single crystal at T \ Tp [34,35]. This may be discussed in terms of the yield point theories [34–37] for systems with low initial dislocation density and a weak dependence of dislocation velocity on stress (i.e. high S, cf. Fig. 1). The number of dislocations with appropriate character appears to be very low at the transition to the new slip system, but they multiply
Fig. 5. Optical micrograph taken during deformation of Fe3Al at 775 K (single glide orientation), showing simultaneous slip activity on {110} (mainly activated), {211} and {321} slip planes.
rapidly during the first stages of deformation and move rather easily at stresses as low as the lower yield stress. This view is supported by our observations of extremely fine slip, suggesting homogeneous viscous slip of dislocations which also has been confirmed directly in TEM in-situ observations (Fe3Al: [22], b-CuZn: [20], and quite similar in Cu–Ge solid solutions: [38]).
5. Conclusions From the correlation of the temperature dependencies of the crss and the strain rate sensitivity with the development and fine structure of slip lines in Fe3Al and b-CuZn intermetallic phases, the following possible dislocation mechanisms have been inferred: In the low temperature region (around room temperature) in DO3ordered Fe3Al, homogeneous slip prevails, probably caused by single imperfect superdislocations, trailing APBs behind them, which move at high stress with cross-slip processes to propagate slip in continuously moving fronts. When the temperature is raised (\325– 475 K for Fe3Al, \ 215 K for CuZn [26]) perfect
Fig. 6. SFM micrograph of slip lines on b-CuZn at room temperature (homogeneous deformation range)
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We have emphasized here the effects of solute mobility on the flow stress and its temperature dependence to explain, in particular, the observed static and dynamic ageing phenomena and to point at the similarities to short range ordering alloys. The ageing effects are supposed to occur in parallel with the possible specific dislocation-structure based mechanisms in long range ordered alloys (e.g. [39–41]) suggested to account for the ‘anomalous’ t0(T) behaviour.
Acknowledgements Fig. 7. Propagation of a homogeneous deformation band front in b-CuZn at room temperature, taken from the positions (s) on the video frames (symbols +), compared with the simultaneously recorded stress (symbols ), showing slight serrations correlated with propagation bursts (note the limited field view of 1 mm of the video frames).
superdislocations may be produced which move at low stress, again multiplying by cross-slip. With increasing temperature, local re-ordering processes in the APB between the partials of the superdislocation pair occur with increasing probability. This may lead to dynamic strain aging effects (‘short range ordering in the long range ordered structure’ [29]), implying clustered slip and stress serrations as well in Fe3Al and also in b-CuZn. The partials of the superdislocation pair get gradually uncoupled, the flow stress has to increase (‘flow stress anomaly’). This increase of stress is interrupted when, around the yield stress peak, another slip system is energetically more favourable, in Fe3Al another slip plane, in CuZn another Burgers vector. This change is accompanied by a transition to smooth and viscous glide, which causes a pronounced yield point at first loading, due to dislocation multiplication and a rapidly decreasing yield stress with increasing temperature.
Fig. 8. SFM micrograph of clustered slip lines on b-CuZn deformed at 425 K. Note the difference to Fig. 6 in the step height scale.
Our thanks are due to Prof. D. Zachmann, Institute for Geosciences, for carrying out the ICP-OES analyses of our specimens. We acknowledge gratefully the financial support by the Deutsche Forschungsgemeinschaft (SPP Verformung und Bruch geordneter Mischkristalle).
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S. Hanada, O. Izumi, Phys. Status Solidi A 40 (1977) 589. N. Brown, Philos. Mag. 4 (1959) 693; Acta Metall. 7 (1959) 210. K. Sumino, Sci. Rep. Res. Inst. Tohoku Univ. 10 (A) (1958) 283. R.S. Rudman, Acta Metall. 10 (1962) 253. A. Nohara, T. Imura, Jpn. Appl. Phys. 25 (1986) 283. J.C. Fisher, Acta Metall. 9 (1954) 9. D.G. Morris, Scr. Metall. Mater. 26 (1992) 733. Y. Brechet, Y. Estrin, Scr. Mater. 35 (1996) 217.
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Structure and development of slip lines during plastic deformation of the intermetallic phases Fe3A1 and CuZn A. Brinck *, C. Engelke, W. Kopmann, H. Neuha¨user Institut fu¨r Metallphysik und Nukleare Festko¨rperphysik, TU Braunschweig, Mendelssohnstr. 3, D-38106 Braunschweig, Germany
Abstract The development and fine structure of slip lines on Fe3Al and b-CuZn single crystals has been investigated by video recording through a far distance optical microscope during compressive deformation and by scanning force microscopy (SFM) after deformation at room temperature up to temperatures beyond the yield stress peak. Different deformation modes, depending on temperature, are identified in Fe3A1 crystals and similarly found in b-CuZn. Homogeneous slip occurs at room temperature and is superposed on heterogeneous localized slip connected with stress serrations in a range of intermediate temperatures up to the yield stress peak, where the type of activated slip systems changes. These variations appear to be connected with the generation and motion of imperfect superdislocations at low and of perfect ones at intermediate temperatures, which, however, are gradually uncoupled with increasing temperature by an antiphase relaxation (local re-ordering) process, until beyond the peak dislocations of another slip system become more favourable. © 1997 Elsevier Science S.A. Keywords: Slip lines; Plastic deformation; Intermetallics; Fe3Al; CuZn; AFM
1. Introduction
2. Experimental details
The analysis of slip lines produced by the movement of dislocations emerging at the crystal surface, provides an easy way to investigate the distribution and activity of dislocation sources, as well as the type of activated slip systems, e.g. for DO3-ordered Fe3Al [1–4] and B2-ordered CuZn [5 – 8]. In the following, we present a systematic investigation on slip line formation and fine structure in these materials at various temperatures, from records taken through an optical microscope during in-situ deformation. The fine structure of the rather faint lines with low step heights is resolved and measured by means of scanning force microscopy. The observed changes of the characteristics of slip line development and structure are compared for the two types of intermetallic phases and are correlated with the simultaneously measured temperature dependence of the flow stress in view of the basic dislocation mechanisms.
Single crystals were produced after induction melting of the pure metals by the Bridgman method in resistance furnaces; Fe–24.8at.%Al crystals in Al2O3 powder crucibles, to avoid any carbon, and Cu–47.5at.%Zn crystals in pure graphite crucibles. Crystals mostly orientated for single slip were selected and prepared for compression tests (Fe3A1 8× 3× 3 mm3, b-CuZn diameter 3.9 mm, gauge length 8 mm) and for tensile tests (b-CuZn diameter 3.9 mm, length 80 mm). After homogenization by annealing for 24 h at 1275 K (Fe3Al) and 775 K (CuZn), respectively, the crystals were cooled down slowly to achieve DO3 order for Fe3Al (cooling rate 0.4 K min − 1) and B2 order for CuZn (cooling rate 1.7 K min − 1). A smooth surface for slip line observations was produced by electrolytic polishing. During compression within a vacuum vessel at about 10 − 2 Pa, with a deformation rate of e; = 8× 10 − 5 s − 1, the specimens were heated from room temperature to 875 K by focusing infrared lamps. During straining, one surface of the crystals was imaged through a quartz window by means of an optical microscope with a working distance of 150 mm providing 2 mm resolution
* Corresponding author. 0921-5093/97/$17.00 © 1997 Elsevier Science S.A. All rights reserved. PII S 0 9 2 1 - 5 0 9 3 ( 9 7 ) 0 0 5 7 9 - 0
A. Brinck et al. / Materials Science and Engineering A239–240 (1997) 180–187
(QUESTAR QM-100), using either bright or dark field illumination. The development of slip lines was recorded on magnetic tape with an attached CCD camera. The tensile tests on b-CuZn were performed in an Instron 1185 machine equipped with a heating chamber under a protective atmosphere of Ar. After deformation the crystal surface was examined by SFM (RASTERSCOPE 4000, DME) using Si cantilevers with an Si tip (NANOSCOPE, cone angle 35°). This method enabled fine slip lines to be resolved in these bcc materials, contrary to EM replica. All investigations were restricted to the first stages of deformation, extending up to a few percent of resolved strain.
3. Slip line formation and structure
3.1. Dependence of the CRSS of temperature Fig. 1a and b compare the dependencies of the critical resolved shear stress (crss) t0 and the strain
181
rate sensitivity S= (1/kT) (t/( ln e; on temperature for the two intermetallics Fe3Al and b-CuZn, in the range from room temperature up to temperatures beyond the so-called yield stress anomaly, where t0 increases with T. Three ranges can be differentiated in Fig. 1a for Fe3Al: A range I of high yield stress around room temperature with smooth load-elongation curves, a range II with relatively low yield stresses where the load-elongation curve exhibits serrations (for small deformation) increasing with T, in particular in the ‘anomalous’ part, before the maximum of the yield stress and a range III beyond this maximum, where the yield stress, decreases again and the load-elongation curves are smooth again, showing however, pronounced yield points at first loading. For b-CuZn, range I is ending already around 215 K [26], therefore Fig. 1b only shows ranges II and III with the same characteristics as for Fe3Al. The results shown in Fig. 1a agree reasonably with measurements in literature, considering the distinct dependencies on the crystal orientations [3] and on the exact alloy composition [9]. For example, for Fe3Al it is known that small amounts of carbon cause formation of carbides [10,11] which increase the flow stress (cf. [3]). The rather sharp drop of t0, seen in Fig. 1a for Fe3Al slightly above room temperature, is attributed to a transition from nucleation and glide of imperfect to perfect superdislocations (cf. [12] and Section 4). The crss values for b-CuZn in Fig. 1b are taken from tensile tests on long crystals, which exhibit a considerable concentration gradient, i.e. they correspond to the softest part of the crystals where, according to ICO-OES analysis, the composition is Cu–47(90.5)at.%Zn. The measured values are in reasonable agreement with literature [13,14]. For short compression specimens cut from parts of the long crystals we observed hardening effects by a- or g-precipitates, when the Zn concentration reached the limits of the b-phase.
3.2. Structure and de6elopment of slip lines
Fig. 1. Dependence of the crss t0 and of the strain rate sensitivity S =(1/kT) (t/( ln e; on deformation temperature T for (a) Fe3Al and (b) CuZn single crystals, oriented for single slip (characteristic ranges are indicated by I, II, III).
a) For Fe3Al, five regimes with different modes of slip line structure could be identified within the temperature ranges shown in Fig. 1a ([12]): (1) In crystals oriented for maximum shear stress on the {110} planes homogeneous slip occurs in faint propagating deformation bands at temperatures up to 325 K on the {110}111 system. The bands can be observed in LM bright field phase contrast (slight defocusing) and appear as a propagating ‘curtain’ on the video, but cannot be resolved properly on single video prints. Fig. 2 shows an example of such a deformation band front taken by SFM. The propagation rate 6prop is around several 10 mm s − 1,
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Fig. 2. Front region of a homogeneous deformation band propagating in range I (Fig. 1a) on a Fe3Al crystal oriented for single slip deformed in compression at room temperature (the wavy surface is the result of electropolishing), showing slip step heights, clustered with less than 2b (b: Burgers vector of a perfect superdislocation) per 103 neighbouring crystallographic slip planes.
corresponding to a local strain rate e; loc :5e; ext, where e; ext =8× 10 − 5 s − 1 is the externally imposed strain rate. (2) In crystals oriented for a maximum shear stress on the {211} planes, localized slip bands with rather high step heights (up to 500 nm) appear during loading and disappear during unloading (‘pseudoelastic effect’ [15]), at high stresses around room temperature and up to temperatures depending on predeformation. Fig. 3a shows LM micrographs (dark field illumination) taken during loading and unloading, Fig. 3b shows an SFM micrograph taken after several loading cycles on one of the few remaining slip band clusters. For more details cf. [16]. Here, the growth rate of slip band step height is S: :5 ×10 − 3 mm s − 1 (at the same e; ext as before), corresponding to a local strain rate e; loc of some 10e; ext. (3) In range II for both orientations at low stresses, inhomogeneous slip occurs on {110}111 in distinct and localized deformation bands, which formed rapidly (Fig. 4a,b). They are, together with static and dynamic strain ageing phenomena [17], connected with serrations in the load-elongation curve for T\ 475 K. Often the homogeneous slip mode (1) continues at both sides of a previously formed inhomogeneous deformation band (Fig. 4c). The rapid growth of step height is in the order of S: :30 mm s − 1 at T = 475 K, corresponding to a local strain rate of e; loc :750e; ext. This value increases with temperature in range II, as shown in Table 1 up to e; loc :4500e; ext. (4) While all slip lines shown before, (with the exception of the pseudoelastic deformation [16]), occur on the {110} slip planes, near the temperature of the
flow stress peak, also slip planes of type {112} and {321} are activated (Fig. 5). Here slip is stable again with a local strain rate of e; loc : 20e; ext. (5) For T exceeding the flow stress peak temperature, slip starts to occur in a very homogeneous manner on the {112} system, with step heights as faint as in mode (1), but no longer in propagating but in randomly appearing deformation bands. Owing to oxidation effects of the surface, the fine structure can no longer be resolved by SFM. Here the local strain rate is e; loc : e; ext. b) For b-CuZn, the slip line characteristics show close similarities to those mentioned above. (1) Around room temperature homogeneous slip occurs in faint deformation band fronts (Fig. 6), which propagate during the micro-yield region on the slip system {110}111, similar to a)(1) in Fe3Al. The front with single slip step heights of S= 0.5–5 nm (near the limit of resolution due to problems with electrolytic polishing) propagates in occasional bursts correlated with slight load drops (Fig. 7). The propagation rate 6prop is about 40–100 mm s − 1 in the average, but reaches several mm s − 1 for short time intervals. (2) After micro-yield at room temperature, faint slip line bundles develop in random distribution on the crystal surface, for crystals oriented for single glide always on the {011}111 system. (3) With increasing temperature, the slip lines get distinctly clustered with step heights of slip lines reaching S= 50–100 nm and undeformed regions between them (Fig. 8). They appear rapidly and are connected with serrations observed in CuZn, too, in
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Fig. 3. Slip bands during pseudoelastic deformation of a Fe3Al crystal oriented for maximum shear stress on {112} planes. (a) Series of video frames (optical micrographs) with slip bands in dark field illumination during loading (1. row up to 2. frame in 2. row) and unloading (succeeding frames). Time inset gives h:min:s:s/100; (b) SFM micrograph (top view and 3D-view) a few slip bands remaining after several unloading/loading cycles. The clustering amounts to about 720 partial dislocation Burgers vectors bp ( =b/4 with b: Burgers vector of perfect superdislocation) per 103 slip planes.
the temperature range corresponding to range II in Fig. 1. (4) From the literature [6,7,18 – 20], we take the result that, for temperatures exceeding that of the yield stress maximum in Fig. 1b, the mode of slip changes from {011}111 to the {011}112 slip system. Here the deformation is homogeneous and slip lines are not resolved.
4. Discussion Comparing the results for Fe3A1 and b-CuZn intermetallic phases, we find the following common features: 1. Homogeneous glide at low T, with slowly propagating fronts. 2. Clustered glide at intermediate T, rapidly forming slip bands connected with serrations in the yield
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Fig. 4. Localized clustered slip bands in Fe3Al appearing rapidly during deformation in range II (Fig. 1a) at T= 475 K in Fe3Al. (a) Successive video frames (optical micrographs) with indicated time scale. (b) SFM micrograph of a typical section. (c) SFM micrograph of a localized slip band with homogeneous continuation at both sides. (d) Step profile of the slip band shown in c. The clustering of the steep part is about 25 perfect superdislocation Burgers vectors per 103 slip planes.
curve; flow stress increasing with temperature up to a stress maximum at the peak temperature Tp. 3. Change of the active slip system around Tp, homogeneous glide for higher T \Tp. Point (1) indicates that, starting from few first activated dislocation sources, the dislocations easily crossslip to neighbouring slip planes [21]. The high stress acting in this range of temperature for Fe3A1 suggests that they are imperfect superdislocations [23], i.e. they are trailing an APB behind them, as confirmed in TEM [3,22]. The high stress, on the other hand, helps in cross-slip of the split screw dislocations [24,25]. With some increase of temperature, when the stress drops to distinctly lower values, the production of perfect superdislocations with narrow APBs between them seems possible [23,12] (for Fe3Al oriented for single glide at T \325 K, for CuZn already below room temperature [26]). The low external stress is able to activate only few favourably situated and configurated sources where some overstress is acting; then each source is able to generate a bunch of dislocations resulting in clustered slip. This clustering and the rapid dislocation generation appear to be en-
hanced with increasing T by local re-ordering processes in the APBs [27–30,19], whose effectivity increases with decreasing dislocation velocity and which cause pinning, breakaway and obstacle destruction effects, similar to short range order hardening [31,29]. As dislocation movement requires lower stress than dislocation production in this region, the slip bands develop so rapidly that some load relaxation of the machine-specimen system occurs and serrations on the load-elongation curve appear. The re-ordering effect (‘APB relaxation’ [32], considered in some detail for b-CuZn in [30]) increases with raising temperature, resulting in an increase of the measured flow stress because the dislocations get more and more uncoupled and move as imperfect dislocations trailing again, an APB behind them, as observed in TEM (CuZn: [20], Fe3Al [22]). This increase would continue if the crystal could not activate dislocations on another slip system which move at lower stress, such as more strongly coupled superdislocations, i.e. {112}111 in Fe3Al, {011}112 in b-CuZn (cf. point (3)).
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Table 1 Variation with temperature of slip modes, of activated types of slip planes, of clustering of slipa and of local strain rates in the active slip bandsb T [K]
Slip mode
Slip plane
Clustering [b/103 d]
e; loc [10−5 s−1]
300 300
Reversible Stable
{112} {110}
5180 B2
\100 40
475 575 675 775
Instable Instable Instable Instable
{110} {110} {110} {110}
25 25 25 25
6000 7500 10 000 38 000
825
Stable
{110}, {211}, {312}
60
160
875
Stable
{211}
a b
None
8
Given in perfect superdislocation Burgers vectors per 103 slip plane distances d. For an external strain rate of e; ext = 8×10−5 s−1 for Fe3Al single crystals.
As the cross-slip activity does not change markedly with T [22], the above re-ordering processes seem to be more probable for explaining the flow stress anomaly in these systems than the cross-slip locking process ([20,25]). The re-ordering mechanism also provides a straight-forward explanation for the static and dynamic strain-ageing phenomena observed in these two alloys [32,17,14,12] and recently modeled phenomenologically by [33]. Another common feature observed in all the alloys showing the flow stress maximum at elevated temperatures, i.e. the long-range ordered intermetallics Fe3Al, b-CuZn, but also short-range ordered Cu-based solid solutions, is the pronounced yield point on first loading of a single crystal at T \ Tp [34,35]. This may be discussed in terms of the yield point theories [34–37] for systems with low initial dislocation density and a weak dependence of dislocation velocity on stress (i.e. high S, cf. Fig. 1). The number of dislocations with appropriate character appears to be very low at the transition to the new slip system, but they multiply
Fig. 5. Optical micrograph taken during deformation of Fe3Al at 775 K (single glide orientation), showing simultaneous slip activity on {110} (mainly activated), {211} and {321} slip planes.
rapidly during the first stages of deformation and move rather easily at stresses as low as the lower yield stress. This view is supported by our observations of extremely fine slip, suggesting homogeneous viscous slip of dislocations which also has been confirmed directly in TEM in-situ observations (Fe3Al: [22], b-CuZn: [20], and quite similar in Cu–Ge solid solutions: [38]).
5. Conclusions From the correlation of the temperature dependencies of the crss and the strain rate sensitivity with the development and fine structure of slip lines in Fe3Al and b-CuZn intermetallic phases, the following possible dislocation mechanisms have been inferred: In the low temperature region (around room temperature) in DO3ordered Fe3Al, homogeneous slip prevails, probably caused by single imperfect superdislocations, trailing APBs behind them, which move at high stress with cross-slip processes to propagate slip in continuously moving fronts. When the temperature is raised (\325– 475 K for Fe3Al, \ 215 K for CuZn [26]) perfect
Fig. 6. SFM micrograph of slip lines on b-CuZn at room temperature (homogeneous deformation range)
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We have emphasized here the effects of solute mobility on the flow stress and its temperature dependence to explain, in particular, the observed static and dynamic ageing phenomena and to point at the similarities to short range ordering alloys. The ageing effects are supposed to occur in parallel with the possible specific dislocation-structure based mechanisms in long range ordered alloys (e.g. [39–41]) suggested to account for the ‘anomalous’ t0(T) behaviour.
Acknowledgements Fig. 7. Propagation of a homogeneous deformation band front in b-CuZn at room temperature, taken from the positions (s) on the video frames (symbols +), compared with the simultaneously recorded stress (symbols ), showing slight serrations correlated with propagation bursts (note the limited field view of 1 mm of the video frames).
superdislocations may be produced which move at low stress, again multiplying by cross-slip. With increasing temperature, local re-ordering processes in the APB between the partials of the superdislocation pair occur with increasing probability. This may lead to dynamic strain aging effects (‘short range ordering in the long range ordered structure’ [29]), implying clustered slip and stress serrations as well in Fe3Al and also in b-CuZn. The partials of the superdislocation pair get gradually uncoupled, the flow stress has to increase (‘flow stress anomaly’). This increase of stress is interrupted when, around the yield stress peak, another slip system is energetically more favourable, in Fe3Al another slip plane, in CuZn another Burgers vector. This change is accompanied by a transition to smooth and viscous glide, which causes a pronounced yield point at first loading, due to dislocation multiplication and a rapidly decreasing yield stress with increasing temperature.
Fig. 8. SFM micrograph of clustered slip lines on b-CuZn deformed at 425 K. Note the difference to Fig. 6 in the step height scale.
Our thanks are due to Prof. D. Zachmann, Institute for Geosciences, for carrying out the ICP-OES analyses of our specimens. We acknowledge gratefully the financial support by the Deutsche Forschungsgemeinschaft (SPP Verformung und Bruch geordneter Mischkristalle).
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