Orientation dependence of dislocation structures in cyclically ...

Orientation Dependence of Dislocation Structures in Cyclically Deformed Cu-16 At. Pct Al Alloy Single Crystals X.W. LI, X.M. WU, Z.G. WANG, and Y. UMAKOSHI The dislocation structures induced by the cyclic deformation of Cu-16 at. pct Al alloy single crystals oriented, typically, as [123] for single slip and [023] and [117] for double slip, were studied by transmission electron microscopy (TEM) and compared with the results of Cu single crystals. Completely unlike the dislocation structures of Cu single crystals of corresponding orientations, the dislocation structures of these oriented Cu-16 at. pct Al alloy single crystals show a typical planar morphology. As the applied-plastic-stain amplitude increases, the dislocation configuration changes, on the whole, from multipolar arrays to dislocation tangles in the primary slip plane and from lowdensity planar slip bands to well-developed persistent Lu¨der’s bands (PLBs) in the planes normal to the primary slip plane, respectively. Secondary slips can be clearly observed to activate from very low plastic-strain amplitudes in all three Cu-16 at. pct Al single crystals investigated. Interestingly, the crystallographic orientation has almost no effect on the dislocation structure of Cu-16 at. pct Al single crystals.

I. INTRODUCTION

IN the past three decades, much development in understanding the cyclic deformation mechanism of fcc single crystals has occurred, and research was devoted mainly to wavy-slip pure metals.[1–4] It is now well known that the cyclic stress-strain (CSS) curve of single-slip Cu single crystals consists of three parts, including a clearly extended plateau. In the plateau region, the dislocation substructure can be well described as two-phase structures, i.e., persistent slip-band (PSB) ladders and matrix veins.[5] The plateau behavior is accomplished by adjusting the amount of PSBs to correspond to the composed plastic-shear-strain amplitude. Although fundamental investigations of the cyclic deformation mechanism in some single crystals of wavy-slip pure metal (e.g., Cu) have yielded a wealth of experimental and theoretical results, there is as yet no general and unequivocal knowledge of the fatigue mechanism of low-stacking-faultenergy (low-SFE) or planar-slip alloy single crystals. Some primitive studies on dilute Cu alloys[6–9] demonstrated that such solid-solution alloys behave with great similarity to their base metal, namely Cu, both in cyclic and dislocation behavior. However, with the increase in solute content, the SFE decreases and the change of slip mode from wavy slip to planar slip would, thus, take place, causing the cyclic behavior to become different from that of pure metals.[8,10–14] X.W. LI, formerly Assistant Professor, State Key Laboratory for Fatigue and Fracture of Materials, Institute of Metal Research, Chinese Academy of Science, Shenyang, 110016, People’s Republic of China, is JSPS Postdoctoral Fellow, Department of Materials Science and Engineering, Graduate School of Engineering, Osaka University, Osaka 565-0871, Japan. X.M. WU, formerly Ph.D. Student, State Key Laboratory for Fatigue and Fracture Materials, Institute of Metal Research, Chinese Academy of Science, is Senior Researcher, Liaoning Electric Power Research Institute, Shenyang, 110006, People’s Republic of China. Z.G. WANG, Professor, is with the Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Science. Y. UMAKOSHI, Professor, is with the Department of Materials Science and Engineering, Graduate School of Engineering, Osaka University. Contact e-mail: umakoshi@ mat.eng.osaka-u.ac.jp Manuscript submitted May 6, 2002. METALLURGICAL AND MATERIALS TRANSACTIONS A

For example, the early work of Youssel[12] and Lukas and Klesnil[13,14] demonstrated that the ladder structures of PSBs, which can be seen often in cycled Cu single crystals, were not found in Cu-30 at. pct Zn alloy single crystals fatigued under a stress-controlled condition, and that the dislocation structures on {111} slip planes were revealed to consist of dislocation pileups or dislocation tangles. Further, Abel et al.[8] reported that the cyclic hardening behavior of singleslip-oriented Cu-Al alloy single crystals with Al contents less than 4 at. pct is similar to that of Cu single crystals, whereas the crystals with Al contents more than 11 at. pct do not exhibit a clear cyclic saturation phenomenon. Later, some investigators[15–21] examined emphatically the fatigue behavior of a typical low-SFE material, i.e., Cu-16 at. pct Al single crystals with a single-slip orientation. It was found that persistent Lu¨der’s bands (PLBs) consisting of densely dislocated arrays on closely spaced parallel primary-slip planes act as the main form of strain localization in fatigued Cu-16 at. pct Al single crystals,[15,16] which is different from the case of copper single crystals, where strain localization takes place in the form of the PSBs featuring the ladder structure of dipolar walls. Inui et al.[21] detected systematically the dislocation structures in single-slip-oriented Cu-16 at. pct Al crystals fatigued at different stress levels under a total-strain-control condition. They reported that primary edge dislocations dominate the structure at all stages of deformation, forming dipolar and multipolar arrays. With increasing stress level, the main change of dislocation structure is an increase in density with little change of form, and, concurrently, the spacing of the activated primary slip planes decreases. Quite recently, Gong et al.[22] reported that the dislocation structures in cyclically deformed Cu-30 wt. pct Zn crystals oriented for single slip have two basic configurations depending on the imposed strain amplitude, namely, dislocation segments and multipoles at low strain amplitudes and planar dislocation loops and tangles at high strain amplitudes. Much of the previous work on the fatigue behavior of fcc low-SFE alloy single crystals has been done mainly on single-slip orientations. However, to the best of our knowledge, VOLUME 34A, FEBRUARY 2003—307

Table I. The Cyclic Deformation Testing Condition for Cu-16 At. Pct Al and Cu Single-Crystal Specimens Observed by TEM Orientation [123] [023] [117]

Fig. 1—Stereographic triangle showing the crystallographic orientations investigated in the present work as indicated by open circles.

there has been little public report on the orientation dependence of dislocation behavior in single crystals of alloys with a really planar-slip character cyclically deformed at a constant-plastic-strain amplitude. Li et al.[23–28] investigated and summarized in a relatively systematic way the CSS response and dislocation structures in differently oriented Cu single crystals and found a strong crystallographic orientation dependence of cyclic deformation and dislocation features of Cu single crystals oriented for double or multiple slip. Recent experimental results by Wu et al.[29] have shown that the crystallographic orientation has nearly no influence on the cyclic deformation and strain-burst behavior of Cu16 at. pct Al single crystals. The aim of the present study is to explore how the crystallographic orientation affects the dislocation structures in cyclically strained Cu-16 at. pct Al alloy single crystals. II. EXPERIMENTAL PROCEDURES Master ingots of a Cu-16 at. pct Al alloy were prepared by melting Cu and Al of 99.999 pct purity in an induction furnace in vacuum. Single-crystal bars of 34 mm in diameter and 65 mm in length were grown directionally from the ingots by the Bridgman method, and the tensile axis was controlled by the use of a copper seed crystal oriented, typically, as [123] for single slip or [023], or [117] for double slip. Fatigue specimens were cut with a spark cutter to a gage shape of a square, 6 ⫻ 4 mm2 in cross section and 16 mm in length, from the as-grown crystals. The orientation of the specimens was determined by the Laue back-reflection technique with an accuracy of ⫾2 deg Figure 1 presents the orientations of Cu-16 at. pct Al crystals investigated in the present work. Before fatigue tests, the specimens were homogenized at 850 ⬚C for 12 hours in vacuum and then electropolished to remove surface strain. Push-pull fatigue tests were performed at room temperature in air using a Shimadzu servohydraulic testing machine under computer control. A triangular waveform signal with a frequency of 0.5 Hz was used for the constant-plastic-strain control. When 308—VOLUME 34A, FEBRUARY 2003

2.7 1.1 6.4 7.4 6.3 7.8 6.2

␥pl ⫻ 10⫺4 ⫻ 10⫺3 ⫻ 10⫺3 ⫻ 10⫺5 ⫻ 10⫺3 ⫻ 10⫺5 ⫻ 10⫺3

N 23,016 12,500 12,000 51,600 11,000 50,400 12,000

␥pl, cum 24.9 55.0 307.2 15.3 277.2 15.7 297.6

the strain burst occurred, the enlarged plastic-strain range can be recovered automatically into the controlled value by using a self-designed computer program. To reveal clearly the three-dimensional distribution of dislocation structures, transmission electron microscopy (TEM) thin foils were prepared parallel to the three specific crystallographic planes (111), (101), and (121), which are mutually perpendicular. These thin foils were first sliced from the gage part of the fatigued specimens by spark cutting parallel to the aforementioned three planes, then mechanically thinned to dozens of microns thick and, finally, polished by a conventional twin-jet method. The TEM observations were carried out using a JEOL*-2000FX II electron micro*JEOL is a trademark of Japan Electron Optics Ltd., Tokyo.

scope operated at 200 kV. III. RESULTS A. The CSS Behavior The fatigue condition for Cu-16 at. pct Al single crystal specimens adopted for TEM observations is listed in Table I, where ␥p l is the applied resolved shear-plastic-strain amplitude and ␥pl, cum (␥pl, cum ⫽ 4N␥;p l , N is being the total cyclic number) is the cumulative plastic strain. It must be pointed out that the strain burst occurred repeatedly during cycling of Cu-16 at. pct Al crystals, and shear stress, thus, fluctuates within a certain stress range. The cyclic hardening behavior of these three oriented crystals has been described in detail in previous work.[29–32] Here, their CSS curves are reproduced in Figure 2, where the typical CSS curve of singleslip-oriented Cu single crystals is also included for comparison. Since Cu-16 at. pct Al alloy single crystals do not show a clear saturation phenomenon, the shear stress at half life is adopted to establish the CSS curve, as shown in Figure 2. For these three oriented Cu-16 at. pct Al crystals, their CSS curves seem to show a roughly analogue plateau, although there are slight differences among their plateau stress levels. Clearly, the plateau region for Cu-16 at. pct Al crystals is shorter than that for Cu single crystals oriented for single slip. To sum up, the crystallographic orientation has no significant effect on the CSS behavior of Cu-16 at. pct Al alloy single crystals.[29] B. Dislocation Structure 1. The [123] Cu-16 at. pct Al single crystal Figure 3 shows typical micrographs of the dislocation structure on the (111) primary slip plane of the [123] Cu16 at. pct Al crystal specimen cyclically deformed at a METALLURGICAL AND MATERIALS TRANSACTIONS A

Fig. 2—The CSS curves of three differently oriented Cu-16 at. pct Al alloy single crystals. The dash line represents the typical CSS curve of Cu single crystals oriented for single slip.[1]

low plastic-strain amplitude of 2.7 ⫻ 10⫺4 Three typical structural regions were found, as follows. (1) Traces of all three secondary slip planes (critical, conjugate, and crossslip planes) were clearly shown on the primary slip plane, although the dislocation density is not so high, as shown in Figure 3(a), where the incident beam B ⫽ 111 and reflection vector g ⫽ 022. (2) A variety of configurations of linear multipoles can be detected in Figure 3(b): point A indicates densely spaced dislocation dipoles on very close neighboring slip planes, point B shows closely spaced multipoles in the form of dislocation trains, and point C demonstrates multipoles with much larger spacing of the dipoles, although the distance of their planes is small. (3) Lots of stacking faults appeared, as shown in Figure 3(c). These three structural regions are observed as being characteristic of this low strain amplitude. The majority of dislocations formed are present in dipolar configurations, and traces of slip planarity have been clearly seen. For constructing three-dimensional dislocation structures, we adopted the (101) and (121) sections for observations, which intersect primary slip planes orthogonally and permit the planar slip bands to be viewed in elevation, with a high probability of intersection. Figure 4 illustrates dislocation structures in these two foils cycled at 2.7 ⫻ 10⫺4. Apparently, the dislocation structures observed in both the (101) and (121) foils are similar, which are mainly characterized by some primary and secondary slip bands with a low dislocation density. With the increase of ␥p l up to 1.1 ⫻ 10⫺3, various kinds of dislocation features can be detected on the primary slip plane (111), as shown in Figure 5(a). At point A, closely spaced multipoles of primary edge dislocations are aligned in the direction parallel to the primary Burgers vector, b; these multipoles show a somewhat irregular pattern as compared with those formed at a lower strain amplitude of 2.7 ⫻ 10⫺4 (Figure 3(b)), which is similar to the dislocation arrangement of kinked dipoles found by Hong and Laird.[19,20] Very few groups of single, primary edge dislocations can be seen in the form of trains, as indicated at point B. In some local regions (e.g., at point C), an interesting structural type, i.e., a regular fish scale–like dislocation net, is seen strictly along the traces of the cross-slip planes. Rare METALLURGICAL AND MATERIALS TRANSACTIONS A

Fig. 3—Dislocation structure in the primary slip plane (111) of the [123] Cu-16 at. pct Al single crystal cyclically deformed at ␥p l ⫽ 2.7 ⫻ 10⫺4. (a) B ⫽ 111 and g ⫽ 022. (b) and (c) B ⫽ 112 and g ⫽ 111. Note that (a) through (c) represent some typical dislocation features observed in different fields of view, respectively (refer to text).

primary screw dislocations can be also observed at point D, which is relatively long and straight in most cases. In addition, some dislocation segments of mixed character distribute in a free state between slip planes, as indicated at point E. It appears that all dislocations are aligned along single slip planes within the available resolution. Figures 5(b) and (c) illustrate dislocation structures in (101) and (121) foils cycled at 1.1 ⫻ 10⫺3. The dislocations are arranged in planar arrays, and edge dislocations dominate the structures. It is VOLUME 34A, FEBRUARY 2003—309

clearly seen in Figures 6(b) and (c), they may exist within those dense slip bands. Clearly, the dislocation structures described previously are, in essence, similar to those observed in single-slip Cu16 at. pct Al single crystals cyclically deformed at constanttotal-strain amplitudes by Inui et al.[21] It is well known that the dislocation structures in cyclically deformed Cu single crystals are susceptible to the imposed plastic-strain amplitude.[1,2,3] However, in the present Cu-16 at. pct Al alloy single crystals with a low SFE, the dislocation structures induced by cyclic deformation are mainly governed by the stress level rather than the plastic-strain amplitude, because no saturation behavior can be found. In the process of cyclic testing of Cu-16 at pct Al crystals under total-strain-amplitude control, the change of plastic-strain amplitude has no obvious effect on the formed dislocation structure. Therefore, it is easy to understand that the dislocation structures revealed by TEM are basically similar under plastic- and total-strain-control conditions, and that the effect of the increase in plastic-strain amplitude on the dislocation structure in the present work is analogous to that of the increase in stress level in the work by Inui et al.[21]

Fig. 4—Dislocation structure in the normal of the primary slip plane (111) of the [123] Cu-16 at. pct Al single crystal cyclically deformed at ␥p l ⫽ 2.7 ⫻ 10⫺4. (a) (101) plane, B ⫽ 101, and g ⫽ 111. (b) (121) plane, B ⫽ 121, and g ⫽ 202.

also evident from these two figures that the dislocation density of the primary slip bands increases markedly compared to that seen in Figure 4, indicating that the slip deformation tends to take place mostly along the primary slip plane. Figure 6 shows dislocation structures in the [123] crystal specimen cyclically deformed at a high strain amplitude of 6.4 ⫻ 10⫺3. The micrographs are taken from the (111), (101), and (121) foils, respectively. Different from the case of lower ␥p l (Figures 3 and 5(a)), the dislocation structure on the primary slip plane (111), typically formed at ␥p l ⫽ 6.4 ⫻ 10⫺3 (Figure 6(a)), shows specific characteristics: many different sizes of dislocation loops are observed to be inhomogeneously distributed and pile up to form dislocation tangles or nets in local regions. Typical structures in (101) and (121) foils normal to the primary slip plane are given in Figures 6(b) and (c), which demonstrate that the dislocation density increases notably and most dislocations are arranged in planar arrays concentrated parallel to the primary slip planes, the spacing of which has obviously decreased compared to those seen in Figure 4 and Figures 5(b) and (c). Furthermore, some well-developed slip bands can be observed along the primary slip planes. Such bands are quite similar to those reported by Buchinger et al.,[16] who termed them persistent Lu¨der’s bands (PLBs). Although dislocation activities on the conjugate and critical planes cannot be 310—VOLUME 34A, FEBRUARY 2003

2. The [023] Cu-16 at. pct Al single crystal Figure 7 indicates the dislocation structure in a [023] Cu16 at. pct Al crystal specimen cycled at ␥p l ⫽ 7.4 ⫻ 10⫺5. The dislocation structure in the primary slip plane, shown in Figure 7(a), is mainly composed of low-density dislocation loops and nets, and some traces of secondary slips are faintly visible. The dislocation structure in the (101) and (121) foils normal to the primary slip plane is presented in Figures 7(b) and (c). Analogues to the case of the [123] crystal (Figure 4), the dislocations are arranged well in planar arrays, and the activity of secondary dislocation is evident. However, it can be seen from Figure 7(b) that some dislocations bands arrange uniformly in rows over a long distance along the conjugate and critical slip planes. The width of these bands was determined to be about 0.6 ␮m, and the spacing between these bands is about 1.0 ␮m. It appears that the density of dislocations on the conjugate and critical slip plane becomes much higher compared to that in Figure 4(a). Careful TEM observations reveal that, for the [023] crystal, there is little change in dislocation configuration with increasing plastic-strain amplitude, except for an increase in density. A typical example at a high ␥p l of 6.3 ⫻ 10⫺3 is given in Figure 8, the micrographs of which are taken from three mutually perpendicular foils. The morphology of the dislocation distributions shown in Figure 8 is very similar to that of the corresponding data in Figure 6. 3. The [117] Cu-16 at. pct Al single crystal Figure 9 shows the dislocation structures in the [117] Cu-16 at. pct Al crystals cycled at different plastic-strain amplitudes. It is very interesting to find that the dislocation structures induced by cyclic deformation of [117] crystals with a conjugate double-slip orientation are completely analogue to those obtained with critical double-slip-oriented [023] crystals. The dislocation structure observed from the (111) foils (not shown here) also consists of low-density dislocation tangles and nets. The dislocations are found to arrange well in planar arrays in the (101) and (121) planes at a low ␥p l of 7.8 ⫻ 10⫺5, and the activity of secondary dislocations is also already evident, as shown in Figure METALLURGICAL AND MATERIALS TRANSACTIONS A

Fig. 5—Dislocation structure in the [123] Cu-16 at. pct Al single crystal cyclically deformed at ␥p l ⫽ 1.1 ⫻ 10⫺3. (a) (111) foil, B ⫽ 111, and g ⫽ 022. (b) (101) foil, B 101, and g ⫽ 111. (c) (121) foil, B ⫽ 121, and g ⫽ 202.

9(a). Apparently, the dislocations can move on the crossslip planes over a long distance and occur along conjugate and critical planes in the form of bands. Such a planarity of the dislocation arrangement is preserved up to a high ␥p l of 6.2 ⫻ 10⫺3; a good example is given in Figures 9(b) through (d). In a word, the dislocation density increases greatly and the crystal is increasingly filled with activated slip planes as the applied plastic-strain amplitude increases. IV. DISCUSSION A. Slip Deformation Mode Combining the existing results on pure Cu crystals[1–4,25–28] with the present results on Cu-16 at. pct Al crystals, it can be realized that large differences in the resultant cyclic dislocation structures exist between these two kinds of crystals. For Cu single crystals, various kinds of characteristic structures such as veins, PSB ladders, labyrinths, and cells can be observed and form in such a way that three-dimensional tangles containing both primary and secondary dislocations are arranged in a low-energy configuration. On the other hand, the overall dislocation structures of Cu-16 at. METALLURGICAL AND MATERIALS TRANSACTIONS A

pct Al single crystals exhibit a planar morphology regardless of applied plastic-strain amplitude. Table II shows the detailed comparison of dislocation structures between these two crystals. Note that the results of [034] Cu single crystals[33] are adopted, since there is a lack of results for [023] Cu single crystals. Such obvious differences are closely related to the difference in the slip deformation mode, i.e., wavy slip or planar slip. Some previous work,[7,21,34–38] which emphasized different aspects of the subject, indicated that there exist some important factors influencing the slip mode, such as the SFE, the short-range order (SRO), the yield stress, the electronatom ratio, and so on. The SFE of Cu decreases by more than one order of magnitude when alloyed with 16 at. pct Al.[39,40] It is well known that a higher SFE leads to smaller separation between the partial dislocations.[39] During the deformation of fcc metals, cross-slip may occur by the pinching (trapping and packing) of partial dislocations of opposite sign in their original slip planes and their subsequent extensions on the cross-slip plane.[41] Therefore, the requirement of pinching the partial dislocations for cross-slip dominates the dependence of dislocation structures on SFE. For pure VOLUME 34A, FEBRUARY 2003—311

Fig. 7—Dislocation structure in the [023] Cu-16 at. pct Al single crystal cyclically deformed at ␥p l ⫽ 7.4 ⫻ 10⫺5. (a) (111) foil, B ⫽ 111, and g ⫽ 022. (b) (101) foil, B ⫽ 101, and g ⫽ 111. (c) (121) foil, B ⫽ 121, and g ⫽ 202. Fig. 6—Dislocation structure in the [123] Cu-16 at. pct Al single crystal cyclically deformed at ␥p l ⫽ 6.4 ⫻ 10⫺3. (a) (111) foil, B ⫽ 111, and g ⫽ 022. (b) (101) foil, B ⫽ 101, and g ⫽ 111. (c) (121) foil, B ⫽ 121, and g ⫽ 202.

Cu crystals, a smaller interpartial spacing corresponding to a high SFE facilitates partial dislocation pinching and, consequently, a greater possibility for cross-slip, resulting in a wavy dislocation substructure. In regard to Cu-16 at. pct Al crystals, on the other hand, a large separation between the partials associated with a low SFE inhibits cross-slip 312—VOLUME 34A, FEBRUARY 2003

and causes dislocations to organize themselves into planar arrays or planar slip bands (e.g., PLBs). Gerold and Karnthaler[36] pointed out emphatically that SRO seems to act as the most important role in slip planarity. They explained the effect of SRO on planar slip as a phenomenon of “glide-plane softening.” Obviously, this phenomenon should be satisfied in the present Cu-16 at. pct Al alloy single crystals.[42,43] More recently, Wang et al.[37] suggested that the electron-atom ratio is a very accurate parameter to determine slip mode. Hong[38] proposed that the slip planarity METALLURGICAL AND MATERIALS TRANSACTIONS A

ing that the Cu-16 at. pct Al single crystal is a typical planarslip material, and the planarity of slip can be regarded as its typical deformation characteristic. In Sections IV–B through IV–E, some dislocation features of cyclically deformed Cu16 at. pct Al single crystals will be discussed and compared with the results of Cu single crystals. B. Plastic-Strain-Amplitude Dependence of Dislocation Structures

Fig. 8—Dislocation structure in the [023] Cu-16 at. pct Al single crystal cyclically deformed at ␥p l ⫽ 6.3 ⫻ 10⫺3. (a) (111) foil, B ⫽ 111, and g ⫽ 022. (b) (101) foil, B ⫽ 101, and g ⫽ 111. (c) (121) foil, B ⫽ 121, and g ⫽ 202.

is promoted not only by a low SFE, but also by increases in shear modulus, atomic-size misfit, and solute content, and he predicted very well the slip mode in other alloys by taking the Cu-Al alloy as a model alloy, according to his new mode. In summary, all the aforementioned experimental phenomena and analyses provided convincing evidence indicatMETALLURGICAL AND MATERIALS TRANSACTIONS A

By comparison, we can note that in fatigued Cu and Cu16 at. pct Al crystals, the variation of dislocation structures with increasing ␥p l is quite distinctive, as shown in Table II. For Cu-16 at. pct Al crystals, the variation can be summarized as follows: planar-slip dislocations with low density (multipolar arrays and planar slip bands) → planar-slip dislocations with high density (dislocation tangles and well-developed slip bands, e.g., PLBs), although there exists a slight difference in dislocation features in the (111) plane at the low-␥p l region between single slip [123] and double slip [023] and [117] crystals. On the whole, such a variation can be regarded as orientation-independent for the Cu-16 at. pct Al crystal. However, for pure Cu crystals, the variation process is not only orientation dependent, but also obviously different from that of Cu-16 at. pct Al crystals, which can be clearly shown by taking single-slip crystals and [117] crystals for examples, as follows: (1) veins → PSBs and veins → labyrinths and cells (single-slip crystals)[1–5] and (2) veins → walls and veins with a reduced size → labyrinths and cells ([117] conjugate double-slip crystals).[28] Undoubtedly, the difference of the change of dislocation structures with varying ␥p l is attributable to the notable difference of slip deformation mode between Cu and Cu-16 at. pct Al crystals. Moreover, the TEM observation revealed an interesting experimental fact that secondary slips could be operated even at a very low plastic-strain amplitude (⬃10⫺4) in the present three differently oriented Cu-16 at. pct Al crystals. The observations on the (101) foil are taken for example. It is evident that for these three crystals, secondary slips have left tangible traces on the (101) plane, as shown in Figures 4(a), 7(b), and 9(a). This phenomenon is evidently different from those observed in Cu single crystals, for which only the primary slip can be detected to activate during cycling at very low ␥p l levels even in multiple- or doubleslip-oriented crystals.[44–47] Gong et al.[22] also found that secondary slips can be activated from a very low-plasticstrain amplitude of 3.8 ⫻ 10⫺5 in a cyclically deformed, low-SFE Cu-30 wt. pct Zn single crystal oriented for single slip. They suggested that the local operation of secondary slips even at very low ␥p l levels is closely related to the low dislocation-generation stress and weak latent hardening effect resulting from heterogeneous deformation in low-SFE materials. Apparently, their explanations are also applied to the case of Cu-16 at. pct Al crystals. We can, thus, deduce a conclusion that the easy operation of secondary slips even under low-strain-amplitude cycling may be a common phenomenon in low-SFE fcc alloy single crystals. C. Strain Localization It is worth mentioning that strain localization of both Cu and Cu-16 at. pct Al crystals takes place in different forms. VOLUME 34A, FEBRUARY 2003—313

Fig. 9—Dislocation structures in the [117] Cu-16 at. pct Al single crystals cyclically deformed at different strain amplitudes. (a) (101) foil, ␥p l ⫽ 7.8 ⫻ 10⫺5, B ⫽ 101, and g ⫽ 111. (b) (111) foil, ␥p l ⫽ 6.2 ⫻ 10⫺3, B ⫽ 111, and g ⫽ 022. (c) (101) foil, ␥p l ⫽ 6.2 ⫻ 10⫺3, B ⫽ 101, and g ⫽ 111. (d ) (121) foil, ␥p l ⫽ 6.2 ⫻ 10⫺3, B ⫽ 121, and g ⫽ 202.

Regarding Cu-16 at. pct Al crystals, dislocations cannot easily cross-slip, and their movements are mostly restricted to its primary slip bands. The applied strain is carried alternatively by different slip bands during cycling, leading to the strain localization in PLBs consisting of densely dislocated arrays on closely parallel primary slip planes, as indicated in Figures 6, 8, and 9. More-clearly-discernible pictures showing such PLBs are presented in Figure 10. These slipdislocation bands or PLBs can be considered as a “pileup” of parallel dislocations. Such particular arrays of dislocations play an important role during cyclic deformation of the crystal. However, a quite pronounced feature of strain localization in cyclically deformed Cu single crystals is the occurrence of PSBs consisting of well-known ladder structures of dipolar walls, by which most plastic strains are carried during cyclic deformation. Buchinger et al.[16] have described in detail the difference between PSBs and PLBs. Here, it needs to be emphasized that the formation of PSBs in fatigued Cu single crystals is governed strongly by the different dislocation reactions associated with the crystallographic orientations, while the PLBs are expected to form in variously oriented Cu-16 at. pct Al crystals because of the planarity of the dislocation structures. 314—VOLUME 34A, FEBRUARY 2003

D. On the Formation of Deformation Twins As is well known, twinning can contribute to the plastic deformation in addition to the slip mechanism when a material is plastically deformed at certain special testing conditions.[40] The nucleation of deformation twinning in fcc crystals is associated with the activation of (a/6)具112典 Shockley partial dislocations. In general, lowering the SFE enhances the propensity of fcc crystals to form deformation twins.[48] Quite recently, Niewczas and Saada[48] observed deformation twinning in Cu-8 at. pct Al single crystals deformed at tension at room temperature and examined closely its nucleation mechanism. However, in the present fatigued Cu-16 at. pct Al crystal samples, nearly no deformation twins were found. In the following text, we will probe tentatively into the formation of deformation twins from two possible influencing factors. One is the critical stress and/or strain needed for twin nucleation. Niewczas and Saada[48] have found in unidirectionally deformed [541] Cu-8 at. pct Al single crystals that, when the change in the dominant mode of deformation from glide to twinning occurs, the resolved stress on the [121](111) twin system is about 110 MPa, which was regarded as the critical stress for twin nucleation, and at this METALLURGICAL AND MATERIALS TRANSACTIONS A

METALLURGICAL AND MATERIALS TRANSACTIONS A

VOLUME 34A, FEBRUARY 2003—315

critical double slip conjugate double slip nearly no

nearly no

[117]

wave slip

obvious

obvious

obvious

nearly no

[117]

conjugate double slip

planar slip

Secondary Slips

within the standard triangle [034]**

[023]

critical double slip

single slip

[123]

single slip

Slip Type

veins

walls; veins

veins

low-density dislocation loops and nets; planar slip bands low-density dislocation tangles and nets; planar slip bands

multipolar arrays; planar slip bands

Dislocation Structures

PSB ladders; veins; labyrinths; cells veins; walls

PSB ladders; veins

—

irregular multipoles; dislocation segments and nets; planar slip bands with a increasing dislocation density —

Intermediate ␥pl dislocation tangles; well-developed slip bands (PLBs)

High ␥pl

labyrinths; cells

walls; cells

labyrinths; cells

dislocation tangles; well-developed slip bands (PLBs)

dislocation tangles; well-developed slip bands (PLBs)

Change of Dislocation Structures with ␥pl

no PSBs

PSBs

PSBs

PLBs

PLBs

PLBs

Strain Localization

strong

nearly no

Orientation Dependence of Dislocation Structures

28

33

1 through 5

present work

Reference*

*Note that the references relevant to experimental results by individual investigators are also given in the table. **Since there are no experimental results reported on the [023] critical double-slip-oriented Cu single crystals, the results of Cu single crystal with [034] orientation, which is quite near the [023] orientation, are adopted in the table for comparison. It is believed that the dislocation behavior of [023] Cu single crystals should be quite similar to that of [034] crystals.

Cu

Cu-16 at. pct Al

Crystal

Low ␥pl

Comparison of Dislocation Behavior during Cyclic Deformation for Cu-16 At. Pct Al Alloy and Pure Cu Single Crystals

Slip Deformation Orientation Mode

Table II.

Fig. 10—PLBs observed in variously oriented Cu-16 at. pct Al crystals. (a) [023] crystal, (101) foil, ␥p l ⫽ 6.3 ⫻ 10⫺3, B ⫽ 101, and g ⫽ 111. (b) [117] crystal, (121) foil, ␥p l ⫽ 6.2 ⫻ 10⫺3, B ⫽ 121, and g ⫽ 202.

point, the crystal specimen has been deformed to 73 pct strain. However, in the fatigued [123], [023], and [117] crystals investigated here, the maximum resolved stress on the possible [121](111) and [211](111) twin systems is about 50 MPa, and the order of magnitude of the strain amplitude imposed is only around 1 pct. Obviously, these values are far less than the aforementioned critical values needed for the nucleation of deformation twins. Despite the fact that the lower SFE of the Cu-16 at. pct Al alloy than that of the Cu-8 at. pct Al alloy would possibly reduce the critical value for twin nucleation, such an effect of the SFE is believed to be limited. Therefore, in the strain-controlled fatigue testing, a sufficient driving force for the formation of deformation twins is hard to provide, due to the limitation of the applied-strain amplitude and the comparatively low resolved stress on twin systems. Another possible influencing factor is related to the lattice shearing and rotation. It is commonly recognized that the formation of deformation twins would cause the shearing deformation of the crystal lattice and the orientation of twinned regions is always different from that of surrounding matrix, indicating an obvious rotation of the crystal. During a unidirectional tensile or compressive test, the irreversible rotation of an fcc crystal easily takes place, whereas such lattice rotation is difficult under a symmetrical tension-compression fatigue loading, although a slight lattice rotation (about 6 deg) has been detected in cyclically deformed Al and Cu single crystals.[49.50] Therefore, in this sense, the deformation under a unidirectional loading condition rather 316—VOLUME 34A, FEBRUARY 2003

than a symmetrical fatigue-loading condition is expected to induce more readily the formation of deformation twins. In making this analysis, it is likely understood that the deformation twins were seldom observed in the cyclically strained Cu-16 at. pct Al single crystals. However, on the other hand, the Cu-16 at. pct Al alloy is a low-SFE material, and slip is accommodated primarily by glide of extended dislocations consisting of pairs of partial dislocations bounded by a stacking-fault ribbon. Under high appliedstrain amplitudes, the continued activity of glide of extended dislocations on parallel sets of planes would eventually be expected to produce overlapping stacking faults, which might ultimately lead to configurations akin to mechanical twins. For example, as shown in Figure 8(c), slip-dislocation bands seem to show somewhat similar configurations to twins. Laird et al.[51] also presented a TEM photograph showing that a PLB emanates from a twin in a cyclically deformed polycrystalline Cu-16 at. pct Al alloy. Therefore, some further experimental efforts are still needed to reveal the possibility of the formation of deformation twins as well as their nucleation mechanism in fatigued low-SFE materials. E. Orientation Dependence of Dislocation Structures As the experimental results of the [123], [023], and [117] orientations indicated, the crystallographic orientation has nearly no effect on the dislocation structures of Cu-16 at. pct Al crystals. For the critical double-slip [023] and conjugate METALLURGICAL AND MATERIALS TRANSACTIONS A

double-slip [117] Cu-16 at. pct Al crystals, the reaction between primary and secondary dislocations would produce, theoretically, Lomer–Cottrell locks and immobile sessile jogs, respectively. However, due to the difficulty of crossslip, the special planar-slip nature would compel the slip deformation to take place mostly along the primary slip plane, inevitably suppressing the spatial dislocation reactions between different slip systems and resulting in the formation of dislocation bands restricted to a few parallel primary slip planes. Thereby, it is easy to understand why the [023] and [117] crystals behave more like a single-slip-oriented crystal (e.g., the [123] crystal) during cyclic deformation. The CSS experimental results (Figure 2) also showed that the crystallographic orientation, indeed, has almost no effect on the cyclic deformation behavior of Cu-16 at. pct Al single crystals,[29] which can be explained well by such an orientation independence of dislocation structures. In contrast, the crystallographic orientation has a strong effect on the dislocation structures in cyclically deformed Cu single crystals.[25–28] For example, a PSB ladder structure can be often seen in cyclically deformed Cu single crystals oriented for single slip; however, it does not occur in some cyclically deformed double-slip Cu crystals (e.g., the [117] crystal) due to the involvement of the secondary slip system. For wavy-slip Cu single crystals, dislocations can be tangled three-dimensionally and can, thus, form various arrangements during cyclic deformation, primarily depending upon the crystallographic orientation, the accumulated plastic strain, and the applied plastic-strain amplitude, etc. In addition, it is worth pointing out that the conclusion on the aforementioned orientation dependence is just based on the experimental results of several differently oriented Cu-16 at. pct Al single crystals. To date, the knowledge of the fatigue behavior and dislocation structures of typical multiple-slip-oriented (i.e., [001], [011], and [111]) and coplanar double-slip-oriented Cu-16 at. pct Al single crystals is still lacking. Accordingly, a more comprehensive conclusion on the orientation dependence requires further experimental efforts. The present results and further efforts are believed to be very helpful to further our understanding of the cyclic plasticity of low-SFE materials. V. SUMMARY AND CONCLUSIONS To clarify the effect of the crystallographic orientation on the dislocation structures in low-SFE materials as well as the role which the slip-deformation mode plays in the dislocation behavior during cyclic deformation of fcc crystals, the dislocation structures of cyclically deformed Cu-16 at. pct Al alloy single crystals, oriented, typically, as [123] for single slip and [023] and [117] for double slip, were studied and discussed in comparison with results of Cu single crystals. The following conclusions can be drawn. 1. Unlike the case of Cu single crystals, the crystallographic orientation has almost no effect on the dislocation structures of Cu-16 at. pct Al alloy single crystals; this is true at least for the present three differently oriented crystals. The difference of cyclic dislocation structures between Cu and Cu-16 at. pct Al crystals can be interpreted in terms of different slip-deformation modes due to the addition of the alloying element Al. 2. The resultant dislocation structures [123], [023], and METALLURGICAL AND MATERIALS TRANSACTIONS A

[117] Cu-16 at. pct Al alloy single crystals present a typical planar morphology. In the primary slip plane, the dislocation structures consist mainly of multipolar arrays at low plastic-strain amplitudes and develop into dislocation tangles at high plastic-strain amplitudes, whereas the dislocation structures in (101) and (121) foils normal to the primary slip plane change from low-density planar slip bands into well-developed PLBs with increasing plastic-strain amplitude. 3. Secondary slips can be activated from a very low plasticstrain amplitude in any of the present three differently oriented Cu-16 at. pct Al single crystals investigated. This behavior is believed to be a common phenomenon in low-SFE fcc alloy single crystals. In addition, PLBs, where the applied plastic strains are highly localized, are expected to form in variously oriented Cu-16 at. pct Al crystals because of the planarity of the dislocation structures.

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