Transition between planar and wavy slip in cyclically deformed short-range ordered alloys

Acta metall, mater. Vol. 42, No. 11, pp. 3759-3765, 1994

Pergamon

0956-7151(94)E0150-F

Copyright © 1994 Elsevier Science Ltd Printed in Great Britain. All rights reserved 0956-7151/94 $7.00 + 0.00

TRANSITION BETWEEN P L A N A R A N D WAVY SLIP IN CYCLICALLY D E F O R M E D SHORT-RANGE O R D E R E D ALLOYS K. W O L F , t H.-J. GUDLADT,:~ H. A. CALDERON§ and G. K O S T O R Z ETH Ziirich, lnstitut fiir Angewandte Physik, CH-8093 Ziirich, Switzerland (Received 28 February 1994) Abstract--Fatigue experiments in Ni-20 at.% Cr and Cu-17 at.%Mn single crystals are reported. Both alloys show short-range order but have different solid-solution hardening parameters. A transition from planar to wavy glide is observed in the late stages of fatigue life. In Cu-17 at.%Mn extrusions occur. Transmission electron microscopy shows that this is due to the development of three-dimensional dislocation arrays from the planar arrays observed in earlier stages. A similar tendency is found for Ni-20 at.% Cr, but the development of extrusions is only seen in the vicinity of cracks. The planarity of the glide and the transition from planar to wavy glide are discussed on the basis of short-range order and the parameters characterizing solid-solution hardening.

1. INTRODUCTION

which eventually give rise to extrusions and intrusions and to crack nucleation sites [6, 7]. In singlephase alloys with SRO, however, slip localization results from planar arrangements of dislocations on

Planar slip is frequently observed in concentrated solid solutions and alloys with coherent particles, Typically, extended planar dislocation pile-ups and

a few isolated slip planes, and the fatigued specimen can usually sustain higher cumulative plastic strains [8]. Additionally dislocation avalanches or strain bursts are typical of the fatigue behavior in these

multipole configurations are found in these materials after deformation as opposed to the threedimensional arrangements of dislocations in pure metals showing wavy slip [1]. Several factors are possible candidates to control the formation of planar slip in single-phase concentrated alloys during cyclic deformation under stress or strain control, Among them, a reduced stacking-fault energy (SFE) or an increased friction stress [2] in comparison with the pure solvent metals and short-range order (SRO) [3,4] have been proposed. Analysing published experimental data, Gerold and Karnthaler [3] suggested SRO to be the most plausible reason for the planarity of glide, based on the fact that glide planes soften when SRO is destroyed by shear. Planar glide and the corresponding softening are also observed in alloys containing ordered particles, such as Ni-A1 or Al-Li. In A1-Li alloys, for example, the degree of glide planarity can be influenced by altering the particle size [5]. Strain localization in metals with wavy glide takes place in the form of persistent slip bands (PSBs) fPresent address: Press- und Stanzwerk AG, FL-9492 Eschen, Liechtenstein. :~Present address: Institut ffir Werkstoffkunde, Fakult/it fiir Luft- und Raumfahrttechnik, Universit[it der Bundeswehr Miinchen, D-85577 Neubiberg, Germany. §Present address: Dept. Ingenieria Metalurgica, E.S.I.Q.I.E.-I.P.N., Apdo Postal 75-874, Mexico D.F. 07300. AMM 42/I I--L

alloys. However there is little information on whether or not strain localization, similar to PSBs in wavy-slip materials, may occur in the late stages of cyclic deformation. Thus it is of special interest to investigate whether or not a variation of the degree of SRO can bring about a transition from planar to wavy slip in single-phase concentrated solid solutions subject to cyclic straining. The amount of SRO can be varied either by comparing different SRO alloys, e.g. A1-Cu, C u - M n as well as Ni-Cr, or by comparing different compositions of a given alloy. In the present research, Ni-20 at.%Cr and Cu-17 at.%Mn were investigated. SRO in Ni-20 at.%Cr has recently been analyzed by diffuse neutron scattering techniques [9]. There is also a reduction of the SFE in this system with respect to pure Ni [10]. However, this dependence seems to be a very weak function of the Cr content [4]. The SFE in C u - M n alloys remains almost unaffected for Mn contents of up to 12at.%, but there is substantial SRO at this concentration [11]. The reduction of SFE by Mn in Cu is even smaller than by Cr in Ni while the degree of SRO is similar [12]. As a quantitative parameter, the diffuse antiphase boundary energy 7SRO can be used. Its calculation requires the knowledge of SRO parameters and pair interaction

3759

3760

WOLF

et al.:

CYCLICALLY DEFORMED SRO ALLOYS

potentials, and a generalized calculation scheme has recently been published [13]. For Ni-20 at.%Cr, 7SRO is 18 mJ/m 2 [13] and for Cu-17 at.%Mn, it is similar [12]. In both alloys planar dislocation arrays have been found. On the other hand, it is known that deformation proceeds by wavy slip in Ni-10 at.%Cr [4] where YsRois around 4 mJ/m 2 [13]. A good correlation between the degree of SRO and the planarity of slip is suggested by these results. In contrast, these two alloys (Cu-Mn, Ni-Cr) show considerable differences in the factors controlling solid-solution hardening, i.e. atomic size and elastic modulus misfits, The aim of the present paper is to report and discuss experimental findings related to the glide behavior during cyclic deformation in these two alloys and to specify the conditions under which a transition between planar and wavy glide may be expected,

2. EXPERIMENTAL DETAILS The alloys Ni-20 at.%Cr and Cu-17 at.%Mn were prepared in vacuum by induction melting. Cu-17 at.%Mn single crystals were grown with an orientation close to [I 2 3] using the Bridgman technique. These crystals were then polished to the final cylindrical specimen shape with a gauge length of 12 mm and a diameter of 3.9 mm. Single crystals of Ni-20 at.%Cr were produced by electron-beam zone melting. Cylindrical specimens 4.5 mm in diameter and 16 mm in length were cut by spark erosion from the original crystal. The orientation of the specimen axis was parallel to [i 2 3] and [[ 3 10]. For cyclic testing cylindrical grips of Nimonic 75A were welded to the specimens. Thereafter, the final gauge section withalengthofl4.8mmandadiameterof4mmwas spark-eroded. The Cu-17at.%Mn samples were homogenized for 48h at 1123K and the Ni-20 at.% Cr samples for 2 h at 1023 K. To establish a well defined thermodynamically stable condition (SRO equilibrium) the Cu-17 at.%Mn crystals were aged for 120 h at 483 K in oil and the Ni-20 at.%Cr specimens for 150 h at 743 K in evacuated quartz capsules. The fatigue tests were performed in a servohydraulic testing machine by applying a sinusoidal wave signal at 1 and 3 Hz. Normally, the plastic strain amplitude was kept constant within 2-5% of the desired amplitude and controlled by a computer, It was not possible to suppress the sudden appearance of strain instabilities. The development of a strain burst depends on the controlling technique but its origin is related directly to the deformation mechanisms in the tested material. Hysteresis loops were continuously monitored, and, if necessary, recorded using a computer. It was thus possible to observe these strain instabilities in detail during cyclic deformarion. Approximately 90 data points per cycle were recorded at a testing frequency of 3 Hz. Surface

observations of the specimens were carried out by scanning electron microscopy (SEM). Additionally, transmission electron microscopy (TEM) was performed in a JEOL 200CX microscope operating at an accelerating voltage of 200 kV. 3. RESULTS The cyclic hardening behavior was investigated under strain control by means of step tests. The corresponding cyclic hardening curve obtained from step tests of a Ni'20 at.%Cr single crystal is shown in Fig. l(a) for plastic strain amplitudes between 8.6 x 10 -5 and 3.2 x 10 -3. Figure l(b) shows the results of step tests o f a Cu-17 at.%Mn single crystal for plastic strain amplitudes ranging from 1.6 x 10 -4 to 1.2 x 10-3, The corresponding plastic strain amplitudes are also plotted in both figures. The cyclic hardening curves are interrupted by sudden stress drops, accompanied by a spontaneous strain exceeding the desired plastic strain amplitude, i.e. strain bursts. The strain burst activity decreases with increasing cycle number and plastic strain amplitude. In later stages of deformation, only a few bursts are seen. A detailed description of the strain burst behavior of Ni-20 at.% Cr single crystals is given elsewhere [14, 15]. The amount of the plastic strain amplitude caused by strain bursts in these step tests is more pronounced for Ni-20 at.%Cr than for Cu-17at.%Mn single crystals. Consequently, this lo0

50 (o)

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~ 60 . . . . . . . "~ ~ ~ ~2 4o " ' . , . | L . [ 20 j 0

. 0

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.

. 400

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. 600

800

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' ' 0 60 90 120 7~'¢~"=Y'4NTp~'" Fig. 1. Cyclic hardening curves (A)from multiple step tests, (a) for a Ni-20at.%Cr single crystal with 7~t (B) ranging from 8.65 x 10-5 to 3.2 x 10-3, (b) for a Cu-17at.%Mn single crystal with y¢ (B) ranging from 1.6 x 10-4 to 1.2 x 10-3.

WOLF et al.:

CYCLICALLY DEFORMED SRO ALLOYS

leads to higher stress drops in Ni-20 at.%Cr specimens than in Cu-17 at.%Mn. The cyclic hardening from one step of the plastic strain amplitude to the following one is more pronounced for C u - 1 7 a t . % M n specimens than for Ni-20at.% Cr single crystals. Generally the slope of the cyclic hardening curve is about four times higher for C u - 1 7 a t . % M n than for N i - 2 0 a t . % C r . In both cases, in the late stages of deformation all specimens are fully covered with slip traces less than 200 nm apart. At the beginning of this late stage a wellpronounced cyclic hardening can be seen in the cyclic stress response. The additional hardening results most likely from dislocation interactions between neighboring active slip planes which could promote for example local cross slip. The occurrence of strain bursts is connected with well-pronounced bands of slip marks on the specimen surface. As an example, Fig. 2 shows the slip marks and the typical hill-and-valley structure which results from the symmetrical cyclic deformation in a C u - 1 7 a t . % M n specimen after a step test. This topology can be easily followed by observing the displacements produced at the scratches marked by arrows. Figure 2 also shows that the smallest distance between slip marks is about 100nm. This surface structure is typical for both alloys. Later in the fatigue life this pattern is only changed by the occurrence of extrusions. As an example, Fig. 3(a) shows a SEM image of the specimen surface of a Ni-20 at.% Cr single crystal fatigued to a high cumulative plastic strain of 7cum= 290. In this case the deformation of extrusions is clearly visible besides planar slip. At higher magnification Fig. 3(b) shows crystallographic cracks within these extrusions. TEM has been performed to investigate the dislocation arrangement. A typical dislocation structure is shown in Fig. 4 for a Ni-20 at.% Cr single crystal with orientation [1 3 10] deformed at Vp~= 1.03 x 10 4 ( N = 106, Z = 44 MPa). The specimen was cut parallel to the primary slip plane. Typical planar slip can be seen from the pile-ups of edge, mixed (60 °) and screw

:

• ~

Fig. 2. Surface slip bands with typical hill-and-valley structure after a step test of Cu-17at.%Mn. The structural changes can be followed by observing the displacements produced at the scratches (see arrows),

3761

Fig. 3. (a) SEM image of a specimen surface of a Ni-20at.%Cr single crystal fatigued to 7p~. . . . = 290. Extrusions are visible. (b) Higher magnification shows a crystallographic crack within the extrusion. dislocations. The pile-ups of screw dislocations suggest planar cross-slip. The TEM images obtained from Cu-17 at.% Mn specimens (7~m = 100, Z = 55 MPa) show slip bands with the corresponding local cross slip events (Fig. 5). The dislocation arrays show not only planar structures, but also three-dimensional arrays of dislocations giving rise to the bands (see arrows) where

Fig. 4. Weak-beam image (zone axis parallel to [1 1 1], diffraction vector g = [~ 0 2]) of the dislocation microstructure of a Ni-20at.%Cr single crystal (Tp~= 1.03 × 10-4 N = 106). The primary slip plane is (1 1 1). The arrows indicate the cross-slip plane (1) and the critical (2) and conjugate (3) glide systems.

3762

WOLF et al.:

CYCLICALLY DEFORMED SRO ALLOYS cumulative strain Ycum of 1360. Bands with threedimensional arrays of dislocations can be seen together with typical planar arrangements of primary dislocations. Within these bands the dislocation density is very high and the activity of local cross-slip is clearly visible. Pile-ups of dislocations of all characters are also visible in this micrograph. The formation of the three-dimensional tangles of dislocations in combination with the development of extrusions [Fig. 3(a, b)] can be interpreted to be an early stage of PSB formation in this alloy. Similar to C u - 1 7 a t . % M n , a transition from planar to wavy glide can be observed in Ni-20 at.%Cr. However, this transition occurs at a somewhat later stage of the fatigue life.

Fig. 5. Weak-beam image (zone axis parallel to [1 ~ 1], diffraction vector g = [1 l l]) of a Cu-17at.%Mn single crystal showing a slip band with local cross-slip events. Band with dislocation tangles is marked by an arrow.

4. DISCUSSION

dislocation tangles are visible. Thus these dislocation bands, characteristic for late stages of the fatigue life, can be taken as a result of the transition from planar to wavy slip. Similar vein-like structures, suggesting wavy slip behavior, are shown in Fig. 6. Here, bright field is used to emphasize with its intrinsic lower resolution the differences between the threedimensional arrays in the bands and their surroundings where characteristics of planar glide still can be seen. The corresponding dislocation density in the bands is very high and it is assumed that these structures represent at early stage of PSB formation, Indeed, careful analysis of specimen surfaces (Fig. 7) reveals micro-PSBs which are normally associated with wavy slip. Figure 8 illustrates dislocation structures in Ni-20 at.%Cr in a foil parallel to (l ~ 1), i.e. perpendicular to the primary slip planes (1 1 1). In this orientation, the Burgers vector of the primary dislocations lies in the plane of observation. This Ni-20 at.%Cr specimen has been fatigued with a plastic strain amplitude of 7p~= 2.1 x 10 -3 to a

Planar slip is observed in both alloys during the early stages of cyclic deformation. This can be related most likely to SRO. In the past, the planarity of the glide has been related to a low SFE, an increase of the friction stress [2, 16], or to short-range order [3, 4]. In Cu-17 at.%Mn, the SFE decreases only by a relatively small amount with respect to the value for Cu. In the case of Ni-20 at.%Cr the SFE is lower than that of Ni but approximately independent of the Cr content. Experiments show that in Ni-10 at.%Cr deformation takes place by wavy glide [4]. Thus a change in SFE with respect to the pure metals cannot be responsible for the planarity of glide. Hong and Laird [2, 16] have based their interpretation of planar slip in Cu-A1 on an increase of the friction stress and the structure of the stacking-fault ribbon separating split dislocations. They developed a model that predicts planar glide for materials with a relatively low SFE, a high solute concentration, a large atomic misfit parameter 6 and a large value of the elastic shear modulus G. While all these features are certainly found in Cu-Al alloys where, in particular, the SFE depends strongly on the alloy

Fig. 6. Three dimensional dislocation arrays in the bands in Cu-17at.%Mn, resembling a vein-like structure (see arrow), can be seen beside characteristic planar slip in this bright-field TEM image (zone axis parallel to [1 ~ 1], diffraction vector g = [1 1 1]).

WOLF et al.: CYCLICALLY DEFORMED SRO ALLOYS

Fig. 7. SEM image of the surface ofa Cu-17at.%Mn single crystal showing a micro-PSB.

3763

composition, this model is clearly not corroborated

Fig. 8. Bright-field image of Ni-20at.%Cr single crystal (70_1= 2.14 × 10 -3, N = 1.6 × 105) with zone axis parallel to [1 ~ 1] and a diffraction vector g = [1 1 1]. Bands with threedimensional arrays of dislocations together with typical planar arrangements of primary dislocations can be seen.

by the present observations since the SFE of the materials used in the present study decreases only marginally with respect to the SFE of the pure materials. Furthermore, the solid solution parameters in Ni-base alloys are considerably lower than in Cu-base alloys (see below). Thus, the friction stress is not a likely candidate for controlling the planarity of the glide, There are strong differences in the parameters 6 and q describing solid solution hardening in Ni- and Cu-base alloys. Solid solution hardening has been extensively studied in terms of two factors characterizing the interaction between solute atoms and dislocations, namely the size misfit 6 [6 = l/a da/dc, where a is the lattice parameter and c the solute concentration] and the change of shear modulus G with alloying, r/(r/ = 1/G dG/dc). Several combinations of 6 and ~ have been suggested to appropriately describe solid solution hardening in Cu, Ag and Au alloys [17-20]. Table 1 shows the parameter 6 and q and some of their combinations for the two alloys of the present study and for the C u - Z n and Cu-A1 alloys studied by other authors [24, 25]. Comparison of the values in Table 1 shows that the solid solution hardening (in terms of 6 and q) expected from

mechanisms involving dislocation interactions with solute atoms in N i - 2 0 a t . % C r is considerably lower than that in Cu-base alloys. The most likely microstructural factor controlling planar glide is therefore the SRO present in the alloys under consideration. This cannot be recognized in systems such as C u - A I or C u - Z n where there is a strong solute concentration dependence of the S F E [2] and simultaneous SRO. In these two Cu alloys additional contributions from mechanisms such as the Suzuki effect have to be considered simultaneously to the contribution of S R O and therefore, the factor controlling planar slip cannot be unambiguously determined. In Ni-20 a t . % C r and Cu-17 a t . % M n , a strong decrease of the S F E or large values of 6 are avoided and the contributions from different microstructural parameters can thus be separated. In these alloys the Suzuki effect plays a minor role since the SFE remains essentially constant. The present observations strongly enforce the view that SRO is the decisive parameter controlling the formation of planar slip, as can be seen from a comparison with past experimental results. Table 2 shows the strong correlation between a given degree

Table 1. Parameters characterizing solid-solution hardening EFa

~=3 ~=16 Alloy 6 r/ Edge Screw E~b E2¢ Ni-20at.%Cr 0.033[22] -0.089[22] 0.184 0.613 0.53 0.040 Cu-17at.%Mn 0.103 111] -0.51 [11] 0.715 2.054 1.65 0.107 Cu-10at.%Al 0.068 [21] -0.48 [23l 0.63 1.48 1.09 0.138 Cu-22.7at.%Zn 0.061[21] -0.58[23] 0.59 1.43 1.00 0.157 The size misfit 6 is given by 1/a da/de where a = lattice constant and e= solute concentration. The modulus misfit ~/is given by I/G dG/dc where G is the shear modulus. G is calculated with G = [e,~(ct~- e~2)]~:2.Errors for 6 and ~ are about 10%. a~v = [ q ' -- 5 3 l,

bet

and r / '

(1

+~ i/2~[18];

(Iq'+ 1661+ I q ' - 166 2 I)[17];

~2 ~/(r/2+ 1662)[20]. 4

3764

WOLF et aL: CYCLICALLY DEFORMED SRO ALLOYS Table 2. Summary of experimentalobservationsrelating SRO to planar glide System 7/G [10-'3 m] Type of slip 7sRo[mJ/m 2] Strainbursts' Ni-10at.%Cr 11.2 [10] Wavy 4 -Ni-20at.%Cr 9.0 [10] Planar 10... 17 Yes Ni-25at.%Cr 9.1 [10] Planar 10...17 -Ni-30at.%Cr 9.5 [10] Planar 24 -Cu-3at.%Mn 10.3 [1l] Wavy Low No Cu-8at.%Mn 8.9 [1I] Planar ~ -Cu-12at.%Mn 10.4[11] Planar ~ -Cu-17at.%Mn Unknown Planar High Yes Cu-2at.%Al 9.1 [2] Wavy Low No Cu-6at.%Al 5.1 [2] Wavy/planar J, No Cu-1 lat.%Al 2.3 [2] Planar ,L -Cu-I 5at.%A1 1.1 [2] Planar ~, Yes Cu-I 7at.%Al 0.9 [2] Planar High Yes Cu-13at.%Zn 8.1 [2] Wavy/planar Low -Cu-20at.%Zn 5.0 [2] Planar J. Yes Cu-22at.%Zn 4.6 [2] Planar 9 Yes Cu-25at.%Zn 3.9 [2] Planar .[ Yes Cu-30at.%Zn 3.3 [2] Planar 15 Yes 7 = stacking fault energy, G = shear modulus, )'srto= diffuse antiphase boundary energy [13]. aA dash means that no detailed observationswere recorded.

of SRO and the occurrence of planar slip. In addition, the appearance of strain bursts in fatigue experiments under plastic strain control directly correlates with the existence of a measurable degree of SRO. In a strain burst, groups of dislocations become suddenly active on new slip planes, giving rise to an overshooting of the desired plastic strain amplirude. C o m m o n fatigue machines are too slow to react sufficiently fast to avoid this type of instability which most likely is due to the creation and multiplication of a high n u m b e r of dislocations. The collective motion of these newly created dislocations during a strain burst is another indication that the increase of the friction stress due to segregation to the stacking fault does not control the planarity of the glide, During a strain burst there is little time for such segregation to occur and one could accordingly expect that the fresh dislocations could glide without any restrictions and produce wavy glide. Since this is not observed, it can be deduced that a mechanism based on glide plane softening is more likely to control the planarity of the glide, i.e. the destruction of SRO on the active glide planes reduces the local resistance to dislocation motion. As a consequence, the dislocation activity is restricted to a few planes and the glide becomes planar. This is in agreement with other investigations [3, 4], and the same reasoning is also applicable to precipitate-hardened alloys [5, 26]. There are also some instructive differences in the fatigue response of the two alloys of the present study. The degree of slip planarity does not seem to be identical in both alloys. Macroscopically, evidence of the higher slip planarity in Ni-20 a t . % C r can be found from the larger stress drops during strain bursts under plastic strain control and the lower cyclic hardening rates. As Fig. 1 shows, the cyclic hardening rate in Cu-17 a t . % M n is about four times higher than that in Ni-20 at.%Cr. Since the degree of SRO in both alloys appears to be similar [12], the

observed differences have to be related to the differences in ~ and r/, which are substantial (Table 1) and could lead to additional effects in Cu-17 a t . % M n (Table 1). A stronger interaction between foreign atoms and dislocations could lead to a more pronounced cyclic hardening [Fig. l(b)]. In Ni-20 at.%Cr, the values of 3 and ~/are considerably lower than in Cu-17 a t . % M n , and little cyclic hardening can be found in this alloy up to ~,p].... values of about 300 and up to values of 7p~= 1.7 x 10 -3 [Fig. l(a)]. It is therefore suggested that a lower degree of planar slip may be related to a stronger influence of dislocation--dislocation interactions. In pure Cu or Ni fatigued under plastic strain control, high cyclic hardening rates and, subsequently, local cross-slip but no strain bursts have been found [6, 27]. In addition, differences in the development of the dislocation structure were observed in the two alloys. Microscopic investigations of the present specimens deformed into the late stages of fatigue life show characteristics usually found in materials with wavy slip, i.e. extrusions and three-dimensional arrays of dislocations. These structures are found to coexist with the planar slip arrays typical for earlier stages of deformation. As Fig. 3(a) shows, extrusions can be found in specimens of Ni-20 at.%Cr, i.e. there is collective strain localization similar to that found in wavy-slip materials. In the present case most of the extrusions are seen close to cracks [Fig. 3(b)] which could be interpreted as an indication that a transition from planar to wavy slip is only possible under complicated stress states such as those at crack tips. Observations in polycrystals show that extrusions can also appear close to grain boundaries where again for compatibility reasons several glide systems must be activated [15]. Similar micro-PSBs have been found in front of cracks in underaged A1-Li alloys [28]. With increasing deformation the dislocation pile-ups typical of planar glide are progressively substituted by three-dimensional arrays as shown in Fig. 5.

WOLF et al.: CYCLICALLY DEFORMED SRO ALLOYS However, a global transition from planar to wavy slip might not be possible in single crystals of Ni-20 at. %Cr owing to the high contribution of SRO (compared to the additional microstructural contributions, e.g. from solid solution hardening) to the resistance to dislocation motion, which reduces the effects of dislocation interactions, i.e. distances between active planes lower than 100 nm (Fig. 2) are necessary to produce a transition. In order to achieve this density of active planes in cyclic deformation, a higher cumulative deformation is required but cannot be achieved because of crack initiation and propagation. A lower contribution of SRO to the slip resistance compared to that resulting from solid solution hardening effects apparently reduces the cumulative plastic deformation necessary for a transition of the glide behavior. In Cu-17 at.%Mn it is possible to find extrusions which are not in the vicinity of cracks (Fig. 7). In addition, as Fig. 6 shows, threedimensional dislocation patterns develop in the late stages of fatigue, giving rise to a configuration resembling the vein structure observed in pure metals [6]. This suggests that in the case of a wellpronounced dislocation~tislocation interaction (or by lowering the contribution of SRO) a transition from planar to wavy slip can be observed for high cumulative strains, i.e. when the distances between active glide planes become small enough to allow local cross-slip and wavy slip to occur (around 100 nm according to Fig. 2). In summary, the present results show that planar glide can be destroyed (i.e. the formation of three-dimensional dislocation structures becomes predominant) by sufficiently strong dislocation interactions, whose efficiency can be related to the contribution of SRO to the total hardening. 5. CONCLUDING REMARKS In agreement with past investigations, SRO has been identified as the controlling factor determining planar glide behavior. However, there are differences in the cyclic behavior of the two alloys compared here. Extrusions and three-dimensional arrays of dislocations were observed in the late stages of fatigue deformation in C u - 1 7 a t . % M n single crystals oriented for single glide. The three-dimensional dislocation arrays developed from the originally planar arrays, as revealed by TEM. The alloy Ni-20 at.%Cr showed a similar tendency but extrusions could only be found near cracks, i.e. under complex stress states. This was interpreted as arising from the difference in the planarity of the glide between the two alloys. This difference also gave rise to differences in the cyclic hardening rate and the characteristics of strain bursts. The individual features of each alloy can be explained if the factors characterizing solid solution hardening are taken into account.

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Acknowledgements--The authors wish to thank Mr E.

Fischer for assistance in growing the single crystals, Mr P. W/igli for the help in SEM investigations and Mrs R. B/inninger for laboratory assistance.

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