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Universal strain-temperature dependence of dislocation structure evolution in face-centered-cubic metals
Abstract
The combined effect of strain and temperature on the microstructural evolution of plastically deformed face-centered-cubic (fcc) metals is explored systematically. In particular, the detailed nanoscale, internal structure of dislocation boundaries is determined in pure polycrystalline aluminum, nickel and gold and compared to earlier results in copper. In all the metals studied, dislocations within the boundaries tend to rearrange themselves with increasing strain in the same sequence from tangles into dislocation cells with tangled boundaries, followed by dislocation boundaries consisting of wavy, parallel dislocations and finally into arrays of parallel dislocations. The strain at which rearrangement occurs decreases with increasing temperature. The results are represented by microstructural maps on the strain–temperature plane. The topology of the microstructural maps is found to be similar for all metals studied, suggesting a universal strain–temperature dependence in deformed fcc metals.
... For copper subjected to a large strain by quasi-static deformation or nickel subjected to a strain of ∼0.3 by dynamic plastic deformation, the microstructure of the as-rolled sample includes two kinds of dislocation populations [16,[29][30][31]. As shown in Fig. 3(a), the first dislocation population is expressed as lamellar boundaries comprising the DDWs and MBs [16]. ...
... The microstructural evolution is affected by the recovery, which reduces the stored energy in the deformed metals by the removal or rearrangement of dislocations. Fig. 3 (d) illustrates that some DDWs contain an array of parallel dislocations that are generally formed in typical pure FCC metals under large plastic strains [29]. The selected area electron diffraction (SAED) suggests that Ag-base phase [27][28][29][30][31][32][33][34] precipitates form in the region shown in Fig. 3(d). ...
... Fig. 3 (d) illustrates that some DDWs contain an array of parallel dislocations that are generally formed in typical pure FCC metals under large plastic strains [29]. The selected area electron diffraction (SAED) suggests that Ag-base phase [27][28][29][30][31][32][33][34] precipitates form in the region shown in Fig. 3(d). Straumal et al. [32,33] suggested that in addition to leading to strong grain refinement in the materials, SPD can also increase the instability of supersaturated solid solution. ...
... Since the feathers and morphology of boundaries strongly depend on the orientations or the active slip systems of the grains, the detailed structure of dislocation boundaries in the nanoscale has gradually been in a hot research at present [4,[8][9][10][11][12][13][14][15][16]. In order to determine the activated slip systems in a grain or different cell blocks, it is essential to identify the different types of dislocations in the dislocation boundaries. ...
... In order to determine the activated slip systems in a grain or different cell blocks, it is essential to identify the different types of dislocations in the dislocation boundaries. The detailed structure of dislocation walls in face-centered-cubic (fcc) metals, such as Al, Ni, Au and Cu, has been discussed only in a few reports [13][14][15]. Hong et al. have paid special attentions on the dislocation content of GNBs aligned with slip planes in rolled aluminium [13]. ...
... After rolled by 10% in thickness reduction, dislocation networks in the boundaries were identified in GNBs [13]. Meanwhile, without considering the orientation effects and the morphology of dislocations, it is found that dislocations within the boundaries tend to rearrange themselves with increasing strain in the same sequence from tangles into dislocation cells with tangled boundaries, followed by dislocation boundaries consisting of wavy, parallel dislocations and finally into arrays of parallel dislocations [14]. However, this topic has not been conducted in bodycentered-cubic (bcc) metals. ...
The evolution of dislocation microstructure in electron beam melted Ta-2.5W alloy was investigated by transmission electron microscope (TEM). Long straight dislocations and dislocation loops are formed in Ta-2.5W alloy cold-rolled by 5%. A set of long, continuous extending planar boundaries (EPBs) are formed when the reduction reaches 20%. In the early stage of development, EPBs are fragmented, diffused and curved, which are connected by non-crystallographic cells boundaries to maintain their continuity. The straight segments of EPBs are usually parallel with the trace of {110}, and incline at about 25–35° to the rolling direction (RD). Two groups of EPBs are formed in a grain when the reduction is larger than 30%. The dislocations within EPBs tend to rearrange themselves with increasing strain in a sequence, from tangled dislocations, followed by parallel long straight screw dislocations and finally into dislocation nets, which are composed by 1/2 < 111 > and [100] type dislocations. The relaxation process of dislocations and the interaction of dislocations with EPBs make EPBs appear wavy and deviate from the trace of slip planes.
... Microstructure in the form of dislocation cells is observed in deformed medium to high stacking fault energy (SFE) metals (e.g. Cu, Ni and Al) [3][4][5][6][7]. In contrast, deformation twins and second-generation microbands dominate the microstructure in deformed low SFE metals [8,9]. ...
... These phenomena have been summarized in straintemperature maps of the dislocation structures for copper, nickel and gold [7]. Each point in the straintemperature plane corresponds to a set of different 1359-6462/$ -see front matter Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. ...
... The regions in the map correspond to distinct dislocation morphologies separated by "phase lines". Figure 1 illustrates the microstructural evolution and the evolution of the detailed nanoscale internal structure of dislocation boundaries in deformed nickel [7]. The strain-temperature maps in all the fcc metals studied are similar topologically and are divided into five different regions, each corresponding to a different nanoscale morphology, which is illustrated in the inset of Figure 1: dislocation tangles (Fig. 1a), dislocation cells with tangled walls (Fig. 1b), dislocation cells with wavy dislocations (Fig. 1c), dislocation cells with ordered walls (Fig. 1d) and a recovered zone. ...
The universal topology of experimental strain–temperature maps of dislocation structures of face-centered cubic metals allows
the ordering of dislocation structure forming processes in these metals, which is not consistent with the stacking fault energy or the melting temperature. Using dimensional analysis, it is shown that the metals can be ordered by the activation energy for cross slip. The experimental maps are scaled by the cross-slip activation energy to form a universal strain–temperature map. The implications for dislocation rearrangement mechanisms are discussed.
... Formation of dislocation loops/ forests/ accumulations seems to be the characteristics of the alloy of low to medium strain hardening exponent [33,[40][41]. In contrast, the dislocation tangles, slip bands, cells, Taylor lattice structures are observed in the alloy pertaining to the medium to high strain hardening exponent [40][41][42]. Hence, many of the precipitates and dispersoids [4,10,20,[27][28]43], their size, shape, morphology, and orientations may have a strong influence on nature of deformation as well as the stress corrosion cracking performances [4,10]. ...
... It was also noticed that the three stages of the dislocation reorganization/ re-arrangement from dislocation loops to tangled dislocation, tangled dislocation to the Taylor lattice structures, and Taylor lattice to the dislocation cell structures (Figs. [3][4][5] are due to mutual interactions and re-organization of the dislocation structures during the processing [42]. A high UTS value as compared to the YS (Table 3) indicates that material can bear maximum load during plastic deformation. ...
The effect of tensile straining on the precipitation and dislocation behavior of the over-aged 7075 aluminum alloy at T7352 temper have been investigated. Microstructures depict dissolution of the precipitates at low tensile straining but reprecipitation occurs at a higher amount of strain. Such precipitation improves the resistance against stress-corrosion cracking as well as other mechanical properties. Detailed characterization displayed the formation of dislocation loops, forest dislocation along with tangled dislocations at a low amount of straining corresponding to a true strain of~ 0.02. In contrast, low-density, high-density Taylor lattices, and dislocation cell structures were seen at higher true straining of 0.06, and 0.1 respectively. Dissolution, reprecipitation of second phases, changes in orientation, and development of dislocation structures are contributing to the two-slope deformation behavior of AA7075T7352 alloy.
... 21,22 Landau et al. have investigated the evolution of dislocation structure in fcc metals. 23,24 It was found that the dislocations in fcc metals within the boundaries tend to rearrange themselves with increasing strain in a sequence: from tangles to wavy, parallel dislocations, and finally into arrays of parallel straight dislocations or dislocation nets. 23 However, this topic has not been conducted in bcc metals by now. ...
... 23,24 It was found that the dislocations in fcc metals within the boundaries tend to rearrange themselves with increasing strain in a sequence: from tangles to wavy, parallel dislocations, and finally into arrays of parallel straight dislocations or dislocation nets. 23 However, this topic has not been conducted in bcc metals by now. It is not clear whether the same sequence would be found in bcc metals. ...
The texture and deformation microstructures of Ta-2.5W alloy were investigated during cold rolling process. The microhardness can reach 280 HV when the reduction was 40%. Meanwhile, the mature body-centered cubic rolling texture was developed. The dislocation configuration appeared in a sequence from long straight dislocations and dislocation loops, followed by dislocation tangles and finally to cells boundaries and long, continuous planar boundaries. Microbands did not appear until the reduction reached 20%. The density of microbands increased with increasing reduction. The dislocations within the boundaries of microbands tended to rearrange themselves with increasing strain in a sequence from tangled dislocations, followed by parallel dislocations and finally into dislocation nets. Meanwhile, the boundaries had at least one primary set of parallel dislocations lying along the longitudinal direction of the boundaries during the whole cold-rolled process. The formation of microbands based on the double cross-slip of long straight screw dislocations was confirmed.
... DDWs block the motion of glide dislocation on the active slip system. Simultaneously, it also traps the dislocation Fig. 9 SEM micrographs: a-c fracture behavior of tensile specimen S1, d-f fracture behavior of specimen S2, and g-i fracture behavior of specimen S3 inside the dislocation walls [42,43], hence causes very high hardening of the alloy by activating the multiple slip system. ...
Microstructure evolution and their effects on mechanical behaviors of the AA7075T7352 aluminum alloy are reported. Phase analysis was done by X-ray diffraction and transmission electron microscope. The presence of the GP-Zones, ɳ', and ɳ, along with Al2Cu, Al2CuMg, and Al3Zr, was noticed. Mechanical characterizations were done with the help of a tensile test and Vickers microhardness. Flow behaviors were studied to evaluate the impact of second-phase particles in the properties. Strain hardening exponents along with UTS/YS ratio have been calculated. Flow curve fitting follows Ludwigson relationship with two distinct slopes. Dislocation loops and forest dislocation were noticed in the low strain range, while dense dislocation walls in the high strain range. Variation in flow parameters is due to the random spread of precipitate particles in the matrix. The material fails by mixed mode of ductile and brittle fractures.
... As a result of the deformation, the Al cell size decreased to 0.3-0.5 μm, as compared to the as-fabricated state. This is consistent with the reduction of cell size due to deformation of pure aluminum [33]. Moreover, the dislocation density within the eutectic phase increased dramatically, while the dislocation density of the Al cells remained low, showing only a few threading dislocations across the cells. ...
Compressive creep properties of AlSi10Mg parts produced by additive manufacturing selective laser melting (AM-SLM) were studied using a spark plasma sintering (SPS) apparatus capable of performing uniaxial compressive creep tests. Stress relief-treated specimens were tested under an applied stress of 100-130 MPa in the 175–225 °C temperature range. Utilizing two different configurations, the creep tests were conducted either with or without a low-density electric current (˜2.63–3.26 A/mm2) flowing through the test specimens. The results revealed that the creep rate increased under the influence of an applied electric current. The creep parameters (i.e., stress exponent n and apparent activation energy Q), were empirically determined. The stress exponent values were found to be 19.6 ± 1.2 and 16.2 ± 1.4 with and without current, respectively, while apparent activation energy was found to be 142 ± 9 kJ/mol and 150 ± 13 kJ/mol with and without current, respectively. The experimental results, together with microstructural examination of specimens, indicate that plastic deformation was controlled by dislocation activity. Furthermore, it is suggested that the annihilation process of dislocations during creep was enhanced by the electric current.
... Several groups have explored a variety of configurations to examine the possible impact of dislocation distribution on mechanical behaviour. For example, strainetemperature dependence of dislocation pattern has been rationalized in terms of a map, which depends on metals and alloys (Steeds, 1966;Anongba et al., 1993;Landau et al., 2011Landau et al., , 2012. Other groups have focused their attention on the impact of grain orientation and/or single crystal orientation on the dislocation structures obtained under tensile and cyclic loading (Kawasaki, 1979, Kawasaki andTakeuchi, 1980;Hansen and Kuhlmann-Wilsdorf, 1986;Feaugas, 1999;Buque et al., 2001;Holste, 2004;Huang and Winther, 2007;Feaugas and Haddou, 2007;Girardin et al., 2015;Jiang et al., 2015). ...
An extensive collection of experimental data was gathered to improve the well-known self-organization process of dislocation resulting from a tensile strain. Statistical analyses by Transmission Electronic Microcopy were performed over a large area in order to obtain additional insights into the different relationships between the structural parameters and the flow stress. A scaling behaviour was established for the size of quasi equiaxed dislocation cells over a large plastic strain range. Additionally, the dislocation organization reveals a variety of scaling laws relative to shear stress, cell size, dislocation wall thickness and dislocation densities in dislocation walls and cells. The physical bases of these laws were demonstrated and their consequences on plastic strain behaviour are discussed.
... The GNB alignment and misorientation as well as the Burgers vectors in that network are similar to the presently analysed observations in aluminium but none of the dislocations were screws. Variations in the visual sharpness of the boundaries in fcc metals have further been reported to scale with the propensity for cross slip [42,43], whereas the grain orientation and slip system dependence of the GNB alignment in general is similar [14]. ...
For the specific slip geometry of two sets of coplanar systems (a total of four systems) in fcc metals, the range of dislocation networks in boundaries aligned with one of the two active slip planes is predicted from the Frank equation for boundaries free of long-range elastic stresses. Detailed comparison with experimental data for eight dislocation boundaries in cold-rolled aluminium grains of the 45° ND rotated Cube orientation is conducted. It is concluded that the boundaries are Low-Energy Dislocation Structures, which are in good agreement with the Frank equation while also lowering the energy by dislocation reactions. Cross slip plays a role in the boundary formation process.
... During plastic deformation of metals of medium to high stacking fault energy, the gliding dislocations interact to give work-hardening and to form dislocation boundaries, which develop into a regular deformation microstructure within each grain. The microstructure evolves with strain and depends on temperature, but a range of fcc [1][2][3][4] and bcc [5] metals exhibit microstructures with common characteristics. The dislocation boundaries in the microstructure may be parallel planar dislocation boundaries or cell boundaries [6]. ...
Previous studies have revealed that dislocation structures in metals with medium-to-high stacking fault energy, depend on the grain orientation and therefore on the slip systems. In the present work, the dislocations in eight slip-plane-aligned geometrically necessary boundaries (GNBs) in three grains of near 45° ND rotated cube orientation in lightly rolled pure aluminium are characterized in great detail using transmission electron microscopy. Dislocations with all six Burgers vectors of the ½1 1 0 type expected for fcc crystals were observed but dislocations from the four slip systems expected active dominate. The dislocations predicted inactive are primarily attributed to dislocation reactions in the boundary. Two main types of dislocation networks in the boundaries were identified: (1) a hexagonal network of the three dislocations in the slip plane with which the boundary was aligned; two of these come from the active slip systems, the third is attributed to dislocation reactions (2) a network of three dislocations from both of the active slip planes; two of these react to form Lomer locks. The results indicate a systematic boundary formation process for the GNBs. Redundant dislocations are not observed in significant densities.
... At this stage one must note that the comparison deals with polycrystals which were strained to only 25%, compared with 50% for the single crystals. The strain which is required to produce a similar microstructure is higher for single crystals, for which the operation of multiple slip systems occurs at higher strain compared with polycrystals, where several slip systems are readily active with the onset of plasticity [24,29]. ...
The thermo-mechanical response of single crystal and polycrystalline high purity copper is systematically compared at low and high strain rates. The mechanical response of each type of material is very different in terms of strain hardening, although both are distinctly
strain rate sensitive. A simplified interpretation of the Taylor–Quinney coefficient, in which the strain dependence is not considered, shows a clear (almost linear) increase of this factor with the strain rate, while the two types show distinct trends. This factor increases with the strain rate but remains markedly lower than the classical value of 0.9. The stored energy of cold work is found to be relatively independent of the strain rate, with the polycrystal storing more energy than the single crystal. A microstructural study (transmission electron microscopy) of representative specimens of each type at low and high strain rates reveals a basically similar microstructure,
despite dissimilar values of energy storage. It is proposed that a higher level of storage of the energy of cold work by polycrystalline copper is due to the presence of grain boundaries in this group.
... Recently, cryogenic rolling of copper has drawn special attention due to formation of bimodal microstructure after appropriate annealing of the deformed sample, which results in excellent combination of strength and ductility456. Extensive studies have been conducted on deformation behaviour of FCC materials at low temperature789. It has been demonstrated that deformation twinning at low temperature governs the plastic deformation, particularly in the case of low and medium stacking fault energy materials like copper [10]. In contrast to FCC materials, experimental studies on low-temperature deformation behavior of BCC materials are limited111213. ...
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Research on the deformation behavior of ultrafine-grained (UFG) pure aluminum (Al) sheets was conducted at a temperature of −196 °C to provide a new strategy for forming complex components with UFG sheets. In this work, a UFG pure Al sheet with an average grain size of 0.9 μm was prepared by cold rolling (CR) and recovery annealing. Then, the mechanical behaviors of the UFG pure Al sheet were evaluated by uniaxial tensile tests and bulging tests. The fracture morphology, microstructure evolution and surface topography were observed by scanning electron microscopy (SEM), transmission electron microscopy (TEM) and atomic force microscopy (AFM). The UFG pure Al exhibits considerable strain hardening and superior ductility at −196 °C, significantly different from the obvious strain softening behavior that occurs at room temperature (RT). As the temperature is decreased from RT to −196 °C, the uniform elongation (UEL) is improved 24.2 times from 1.2% to 30.2%, and the total elongation (TEL) is improved 3.3 times from 11.8% to 51.2%. The limiting dome height (LDH) and the ultimate bulging rate, K, of the UFG pure Al sheet at −196 °C are 1.6 times and 2.9 times greater than those at RT, respectively. Microstructure observations show that many dislocation substructures such as dislocation pile-ups, dislocation cells and dislocation networks are formed in UFG pure Al deformed at −196 °C, that would otherwise be absent in its counterpart deformed at RT. This unusual phenomenon is attributed to the remarkable transition of the dominant deformation mechanism of UFG pure Al from significant dynamic recovery accompanied by grain boundary sliding (GBS) to dislocation multiplication and accumulation with decreasing temperature.
In this paper, we investigated the effects of building orientation and heat treatment temperature on the anisotropy of SLM-formed AlSi10Mg specimens. Three different deposition orientations of 0°, 45° and 90° and heat treatment temperatures of 270°C, 300°C and 330°C were used to hold the specimens for 2 h, followed by furnace cooling. The metallographic organization, fracture morphology, grain size and morphology, grain orientation and distribution were characterized by optical microscopy (OM), scanning electron microscopy (SEM) and electron backscatter diffraction (EBSD). SEM metallographic shows that the white cell structure is eutectic Si where Si grid-like structure distributes around α-Al dendrites. EBSD results show that all specimens consist of a large number of grains growing along <001> direction, which is the preferred orientation. The organizational changes in grain size, distribution and low-angle grain boundaries (LAGBS) caused by orientation and heat treatment were analyzed to confirm the anisotropy of the SLMed AlSi10Mg alloy and to analyze and explain the reasons for the poor plasticity of the 90° orientation. The present results provide new insights into the effects of deposition direction and heat treatment parameters on the microstructure and properties of AlSi10Mg alloy prepared by SLM.
The FCC Fe40Mn40Co10Cr10 (at. %) high entropy alloy exhibits deformation twinning for room temperature deformation and deformation-induced HCP transformation for subzero deformation. Since the systematic investigation of temperature-dependent dislocation structures is not available, we present an in-depth characterization of the defects involved in deformation at room temperature (298 K) and subzero temperature (223 K) of the cold-rolled and annealed Fe40Mn40Co10Cr10 alloy. The material deformed at 223 K shows a higher strain hardening rate than the material deformed at room temperature while both materials show large ductility, 48 % for the 223 K deformation and 55 % for the 298 K deformation. The main deformation mechanisms of the investigated HEA include the development of inhomogeneous dislocation structures and interaction between dislocations and deformation twin/mechanically induced HCP martensite. The stacking fault energy measured using TEM weak-beam dark-field imaging of dissociated dislocations is 20±9 mJ/m² at 298 K. The Fe40Mn40Co10Cr10 alloy exhibiting a positive temperature dependence of SFE leads to a decrease of SFE as deformation temperature decreases from 298 K to 223 K. The decrease of SFE results in the transition from deformation twinning to deformation-induced HCP transformation. Further, at higher strains at 223 K, kink banding of HCP and reverse transformation from HCP to FCC were observed, which could account for strain accommodation and stress relaxation, and the large ductility. The 298 K deformation leads to various types of dislocation structures: Extended dislocations, Taylor lattice of perfect dislocations, dislocation loops, highly dense dislocation walls, cell blocks, and cell structures. The observed dislocation structures at 298 K and 223 K are similar suggesting the minor effect of SFE on dislocation structures.
X-ray diffraction peak broadening is discussed in terms of line broadening and rocking-curve broadening in a novel theoretical description. The nonlocal strain tensor is factorized by using the method of polar decomposition instead of the more conventional separation into symmetrical and antisymmetrical components. A number of X-ray line-broadening and rocking-curve experiments on the same single crystals or individual grains in bulk polycrystals prove that plastic deformation produces strained subgrains mutually rotated by rigid-body rotations. The novel theoretical description appropriately accounts for the rigid-body rotation and strain at the same time and provides straightforward separation of the two effects of line and rocking-curve broadening in the radial and normal directions of the diffraction vector. The mathematical results are discussed in terms of experiments of X-ray diffraction, Laue asterism and electron backscatter diffraction. From the experimental results it is shown that the simultaneous evaluation of line and rocking-curve broadening provides qualitative information about the redundant and geometrically necessary character of dislocations, not available if only one or the other is accessible.
The combined effect of strain and temperature on the microstructure and detailed internal structure of dislocation boundaries was systematically studied in compressed pure polycrystalline copper and nickel and compared to the microstructure of compressed polycrystalline aluminum. Below 0.5Tm the microstructure of Cu and Ni consists of dislocation cells, however, only in Cu second generation microbands are formed. In Cu and Ni, the dislocations inside the boundaries rearrange themselves from tangles to ordered arrays of parallel dislocations following interplay between strain (requirement for cross slip) and temperature (dislocation mobility and ease of cross slip). The ordered detailed structure is similar to that observed in Al deformed at room temperature and lower strain levels. The amount of strain and temperature applied to Cu and Ni in order to achieve the same detailed structure formed in Al depends on the stacking fault energy (SFE) of the metal- higher strain and temperature as the SFE is lower.
Deformation of metals from medium to high strain significantly affect the deformation structure as well as the recovery and recrystallization behaviour when deformed samples are annealed. This behaviour is illustrated for FCC metals of medium to high stacking fault energy, with emphasis on the behaviour of aluminium and aluminium alloys deformed by cold rolling to large strain. The analysis encompasses hardness testing, EBSD and TEM. The deformation microstructure is a lamellar structure of dislocation boundaries and high angle boundaries where the percentages of the latter increases to about 60-80% at large strain. The macrotexture is a typical rolling texture, which is composed of individual texture components present as micrometre and submicrometre size volumes. In the lamellar structure correlations have been established between microstructural parameters and local orientations showing for example variations in stored energy between the texture components and large variations in the spatial distributions of the high angle lamellar boundaries. Such local variations can affect the structural coarsening during recovery at low temperature leading to significant structural difference on a local scale. The local variations in the deformed structure can also significantly affect the structural changes taking place locally during high temperature annealing thereby affecting the evolution of the structure and texture on a macroscopic scale.
The effect of deformation-induced disorientations on work-hardening of metals is modelled by dislocation dynamics. By incorporating excess dislocations related to disorientations, Kocks' dislocation model describing stage III hardening is extended to stage IV. Disorientations evolving from purely statistical reasons still lead to stage III behavior and a saturation of the flow stress, but deterministic contributions to the development of disorientations, as differences in activated slip systems across boundaries, cause a linear increase of the flow stress at large strains. Such a constant work-hardening rate is characteristic for stage IV.
The texture development during rolling of three Ni-Co alloys, with 31, 41 and 61wt%Co, has been investigated using electron backscattering diffraction. Already after 20 and 35% reduction the textures in the three alloys are clearly different. The texture development in Ni-30%Co is towards a copper-type texture, and the texture development in Ni-60%Co is towards a brass-type texture. In Ni-40%Co the texture development is in between those in the other two alloys.
The effects of strain and temperature on the microstructure and the detailed structure of dislocation walls in compressed pure polycrystalline copper were systematically studied. The microstructure observed consisted of dislocation cells and second generation microbands (MB2) and is presented in a microtructural map. The detailed structure of the cell's boundaries evolves with increasing strain and/or temperature from tangled dislocations into arrays of parallel dislocations. Interplay between strain and temperature controls the microstructure and the detailed structure of the dislocation boundaries. Above 0.5 T-m only MB2 are observed; a different type of MB2 is also observed, one order of magnitude wider than the one observed at the lower temperature. With increasing strain and temperature the MB2 density increases. It seems that dislocation cells and MB2 are competitive dislocation patterns and at elevated temperatures the MB2 is the preferred dislocation pattern.
The effect of strain on the microstructure and detailed internal structure of dislocation boundaries in pure polycrystalline fcc metals (aluminum, copper, nickel and gold) was systematically studied and compared as a function of strain following compression at room temperature. At low strains all metals form a cellular structure. The dislocations in the cell walls tend to rearrange themselves from tangles to ordered arrays of parallel dislocations as the strain is increased. However the rearrangement does not correlate well with the stacking fault energy. A good correlation is found with the cross-slip activation energy.
The evolution of dislocation patterns in single crystal aluminum was examined using transmission electron microscopy (TEM). In-situ tensile tests of single crystals were carried out in a manner that activated double slip. Cross slip of dislocations, which is prominent in all stages of work hardening, plays an important role in dislocation motion and microstructural evolution. In spite of the limitations of in-situ straining to represent bulk phenomena, due to surface effects and the thickness of the samples, it is shown that experiments on prestrained samples can represent the early stages of deformation. Transition between stage I and stage II of work hardening and evolution during stage III were observed.
The role of stacking fault energy (SFE) in deformation twinning and work hardening was systematically studied in Cu (SFE ∼78
ergs/cm2) and a series of Cu-Al solid-solution alloys (0.2, 2, 4, and 6 wt pct Al with SFE ∼75, 25, 13, and 6 ergs/cm2, respectively). The materials were deformed under quasi-static compression and at strain rates of ∼1000/s in a Split-Hopkinson
pressure bar (SHPB). The quasi-static flow curves of annealed 0.2 and 2 wt pct Al alloys were found to be representative of
solid-solution strengthening and well described by the Hall-Petch relation. The quasi-static flow curves of annealed 4 and
6 wt pct Al alloys showed additional strengthening at strains greater than 0.10. This additional strengthening was attributed
to deformation twins and the presence of twins was confirmed by optical microscopy. The strengthening contribution of deformation
twins was incorporated in a modified Hall-Petch equation (using intertwin spacing as the “effective” grain size), and the
calculated strength was in agreement with the observed quasi-static flow stresses. While the work-hardening rate of the low
SFE Cu-Al alloys was found to be independent of the strain rate, the work-hardening rate of Cu and the high SFE Cu-Al alloys
(low Al content) increased with increasing strain rate. The different trends in the dependence of work-hardening rate on strain
rate was attributed to the difference in the ease of cross-slip (and, hence, the ease of dynamic recovery) in Cu and Cu-Al
alloys.
Deformation microbands coincident with {111} slip planes are observed to be associated with impact craters in large-grain (0.8 mm) copper (stacking fault energy [SFE] 80 mJ/m2) while microtwins coincident with {111} slip planes are observed in large-grain (∼1.1 mm) brass (70 Cu–30 Zn; SFE ∼10 mJ/m2). Taken together with other impact crater studies, there is solid evidence for a strong influence of SFE in the microband–microtwin transition.
The evolution of the cold deformation microstructure is described for medium to high stacking fault energy, single phase fcc metals. Macroscopic strain accommodation for polycrystalline metals is considered, and it is suggested that grains subdivide during deformation on a smaller and smaller scale, and that each volume element is characterised by an individual combination of slip systems. A number of microstructural observations (especially of aluminium, nickel, and copper) are described, and dislocation arrangements are discussed on the basis of the general principle that they are low energy dislocation structures. It is shown that the microstructural evolution is quite similar in polycrystalline metals and in single crystals deforming by multislip, and ways in which metallurgical parameters such as stacking fault energy and grain size can affect the microstructure are examined. The general principle of grain subdivision during cold deformation is discussed with reference to the microstructural observations.
Transient creep curves were used to compare cell structures in polycrystalline Al produced by monotonic steady state creep at various temperatures and by tensile stress cycling at room temperature. Large plastic strain at a given shear modulus-corrected stress results in a dislocation structure which depends neither on temperature nor on deformation mode.
The effects of deformation on the structure of copper and its low stacking-fault energy alloys are reviewed and it is shown that these may be classified according to the method of examination used. In rolled low stacking-fault energy alloys the deformation sequence involves the formation of stacking faults, mechanical twins, and shear bands that are at first restricted to individual grains but subsequently extend from one surface to the opposite surface. In copper the deformation sequence involves the formation of equiaxed cells of dislocations, microbands, clustering of microbands, and shear-band formation.
The configuration of dislocation wall structures and the interactions between dislocations and dislocation walls play, a significant role in the understanding of deformation processes in metals. Samples of single-crystal aluminum deformed by tensile-straining (15%) were analyzed using TEM. In tensile-deformed (15%) single crystal aluminum, a cell structure is well developed and dislocations in the cell boundaries consist of either one set of Burgers vector or two sets of Burgers vector. The three-dimensional image of cell wall structure, misorientation angle across the cell boundaries and the Burgers vectors of dislocations in the cell boundaries are characterized.
Dislocation dynamics simulations of multiple slip in f.c.c. crystals lead to the formation of patterned microstructures. The mechanisms participating to dislocation storage and dynamic recovery are investigated and discussed. Cross-slip and short-range interactions are found to govern the bifurcation from uniform to ordered microstructures.
This paper concerns a study of the nature of shear bands in Cu-Al single crystals. The experimental results show that during the rolling deformation two families of shear bands form independently of the crystal orientation. However, initial orientation of the crystal was found to influence both the onset of band formation and orientation of the first family of bands with respect to the rolling direction. It appears from the TEM observations that shear bands result from correlated multiple coarse slip events. The results show that formation of macroscopic shear bands on an apparently nonrational plane results from the instability of pre-existing structures and the occurrence of a sequence of coarse slip events plus accommodating localized plastic flow.
From a widespread convention, glide in crystalline materials is classified as 'planar' versus 'distributed'. At elevated temperatures, 'distributed glide' is favored and is almost usually found at medium to high strains in pure metals, in low-concentration homogeneous alloys, and in a preponderance of multiphase alloys, provided temperatures are not too low. While at low temperatures, 'planar glide' is widely observed and favored in solid solution alloys. In this paper, a general explanation is presented over the proposed additional causes for planar versus distributed glide, based on the LEDS theory.
The relation between the polycrystal deformation and single crystal deformation has been studied for pure polycrystalline copper deformed in tension. The dislocation microstructure has been analyzed for grains of different orientation by transmission electron microscopy (TEM) and three types of microstructures have been identified. A correlation is found between microstructure and grain orientation, which agrees well with earlier observations in tensile deformed aluminum polycrystals and copper single crystals. The stress–strain curve of the copper polycrystal is calculated with good accuracy from single crystal data, which are weighted according to the volume fractions of the three different types based on a quantitative texture measurement of the polycrystal.
Torsion deformation was used to investigate dislocation substructure evolution at large strains in high purity nickel and NiCo solid solutions. Observations of small strain dislocation structures formed in stage III revealed that the laminar dislocation structure observed after large strains in stage IV develops from short paired dislocation sheets within the tangled dislocations of an equiaxed cell wall. The development of these short paired dislocation sheets into long microbonds occurs gradually by a multiple-slip process in accordance with the principles of low energy dislocation structures and without the occurrence of a shear instability. The plane of these sheets and /or microbands does not correspond to a {111} slip plane. As these microbands form, a misorientation between the interior of the paired sheets and the surrounding matrix develops and increases with increasing strain.
The development of dislocation cells in polycrystalline aluminium and nickel has been evaluated over a tensile strain range, 0.05–0.30, using transmission electron microscopy. The results of the two test series are compared. Measurements of both the cell sizes and relative cell misorientations have been made, the latter using microdiffraction. The results are then compared with previous studies from the literature.
Secondary creep has most generally been associated with a rather steady structure. Many models have been suggested to explain the constant strain rate in terms of the effective stress which is determined by the structurally-dependent internal stresses. The internal stresses deduced macroscopically have been of the order of half the applied stress. In this article, by pinning the dislocations under load in an Al-Zn alloy, the evolution of the structure and local effective stresses with strain has been identified by electron microscopy. Values of local effective stresses at the subgrain boundaries ranging between 10-20 times the applied stress have been measured. The emission of dislocations from these boundaries and the evolution of substructure within the subgrain interior indicate that the controlling mechanism during the creep process is the relaxation of internal stresses by this emission. At the same time, the subboundary stress fields existing in different subgrains determine their different behaviour as a function of time. Hard and soft subgrains alternate in the deformation process to produce overall uniform strain.
Plastic deformation leads in many metals to a continuous refinement of the microstructure. An analysis of certain key microstructural parameters reveals a scaling behavior, reflecting a continuity of the basic processes underlying plastic deformation over a very wide range of strain, and hence for structures from the micro- to nano-scale dimensions.
Work-hardening phenomena are based on the very fundamental principles (i) that at the position of every dislocation axis the respective resolved shear stress cannot exceed the friction stress, including the self-stress of bowing dislocations, and (ii) that always that structure forms which among those accessible by the dislocations minimizes stored energy per unit length of dislocation line. Such dislocation structures have been named LEDSs. The corresponding work-hardening theory, the mesh length theory, is applicable to all materials deforming via gliding dislocations and to all types of deformation. Results previously achieved with the mesh length theory are summarized, and a number of new developments are discussed. Depending on the dislocation structures formed, the work-hardening behavior differs. Easily intersecting glide causes dislocation cell structures with almost dislocation-free cell interiors delineated by dislocation rotation boundaries. Pronounced planar glide causes Taylor lattices characterized by local planar order parallel to the one or perhaps two most highly stressed glide plane(s), no systematic lattice rotations, and overall uniform dislocation density. The most widely observed basic features of work hardening are explained in general terms. Specific applications are indicated for layer-type crystals, h.c.p. single crystals, single-crystal and polycrystalline pure f.c.c. metals and α-brass-type alloys, precipitation-hardened materials and steels. Included are the different stages of work hardening, dynamical effects in low temperature plasticity, the general characteristics of grain boundary strengthening and the Hall-Petch relationship. In addition, proposed explanations for (i) glide system interactions in polyslip resulting in microbands and affecting texture formation, and (ii) creep without stress dependence of dislocation density, are discussed.
Layered cell structures in deformed copper single crystals with the [112] and the [415] tensile axes were observed by transmission electron microscopy on slices with (\bar{1}\bar{1}1), (1\bar{1}0), and (001). It was confirmed that cell walls are formed on planes rotated around the axis on the active slip plane to a definite direction with respect to the tensile axis. The rotation angles were measured as a function of tensile stress. The slip line lengths on the two side planes were measured by optical microscopy. The slip line lengths and the ratios of the slip line lengths of edge dislocations to those of screw dislocations agree approximately with the slip distances estimated from the spacing and the rotation angle of the cell walls, assuming that the cell walls of the primary system are obstacles for primary dislocations. It is concluded that the slip line length is mainly determined by the layered cell structure.
The microstructural evolution is followed in pure aluminium and nickel cold-rolled over a large strain range. A number of dislocation configurations are characterized and classified and it is found that dislocation rotation boundaries are the dominant feature which subdivide the grains on a finer and finer scale as the strain is increased. These configurations of dislocation boundaries are analyzed on the basis of the LEDS hypothesis for dislocation structures and agreement is found. The strengthening effect of dislocation boundaries is discussed and equations are suggested for the relationship between flow stress and microstructural parameters.
Copper single crystals with the [145] and the [112] axes were deformed in tension at room temperature and dislocation structures were observed by transmission electron microscopy. In the later stage of deformation, layered cell (carpet) structures were formed nearly parallel to the active slip planes. Close examination showed, however, that cell boundaries inclined from the slip plane by several degrees around a axis on the slip plane. The cell boundary densities were approximately proportional to tensile stress and the proportional constants in both crystals were nearly the same. The observed results were consistent with the fundamental assumptions in Takeuchi's theory of work hardening.
Past research leads to the conclusion that dislocation cell formation in work-hardened crystalline materials occurs when dislocations assemble into low energy configurations. On the assumptions that the dislocation cells formed in f.c.c. metals in early stage II also approach the lowest energy for a given dislocation content in the material and that the initial dislocation arrangement before cell formation consists of linear dipolar mats, i.e. sets of similar edge dislocations in coplanar arrays but alternating sign from one mat to the next, the structure of the resulting cells is investigated. It is found that the initial pile-up-like arrays should transform into tilt walls with or without some twist component, such that the axis of relative misorientation is roughly parallel to the original edge dislocations. By a simple consideration of energies it is found that cell formation should begin at or below about 1.2τ0 where τ0 is the initial critical flow stress. Again if the minimum energy is considered, it is found that for low stacking fault energy materials Lomer-Cottrell locks should form prominently, such that the primary dislocations become rotated roughly normal to the Lomer-Cottrell locks. These results are in good agreement with available experimental evidence.
The aim of this study is to work out a microscopic picture of the second stage of creep of aluminium at intermediate temperatures. The substructure and its evolution during creep having been described in the two former articles, we detail here the movement of individual dislocations responsible for the main part of deformation, during in situ experiments in a high voltage electron microscope. The mobile dislocations are emitted by sources situated inside a few subgrains. They have to cut through subboundaries, and this appears to be the rate controlling process. This mechanism which is described as insertion followed by extraction, is studied in detail, as well as annihilation of mobile dislocations in other subboundaries. It is easier than the Orowan process, and it involves cross slip, but no climb. However, the stress necessary is higher than the applied stress, so that extraction requires high local forward internal stresses in subboundaries. The different types of recovery observed—static and dynamic—are discussed, and it is concluded that cross slip is the rate controlling mechanism of creep of aluminium at intermediate temperatures.
The substructure behaviour of Aluminium has been studied at intermediate temperatures in order to determine the microscopic mechanisms which control the strain rate. This first article describes the detailed geometrical features of the dislocation networks after the creep test. The subboundaries are made of the dislocations emitted by the sources which are activated by the local stress; most of them are of mixed character, exhibiting 3 coplanar Burgers vectors at 120°; their long range stress field, if any, is smaller or equal to the creep stress; the small dislocation segments are situated in their respective glide planes, which brings some restrictions on the possible network geometries. In the subsequent articles, these features will appear as essential to understand the dynamic properties of the substructure during in situ creep experiments in a high voltage electron microscope, and to work out a new picture of creep at intermediate temperatures.
It has been seen that there is a clear correlation between the type of microstructure and the grain orientation, which was found to be, to a great extent, comparable with that obtained in tensile strained single crystals of copper where correlation was observed between the crystallographic orientation and the microstructural evolution. In tensile strained polycrystalline copper specimens, a grain to grain variation in the deformation microstructure has also been observed. For example, equiaxed cells and parallel dislocation walls have been found in different grains of an oxygen-free high purity copper. These variations have been related to the slip patterns predicted on the basis of a Schmid factor analysis. However, only a limited number of grains have been examined and the exact grain orientation data were not included. Thus, a clear relationship between the microstructure and grain orientation has not been established in polycrystalline copper, and this has been the aim of the present study.
The evolution of the cold deformation microstructure is described for medium to high stacking fault energy, single phase fcc metals. Macroscopic strain accommodation for polycrystalline metals is considered, and it is suggested that grains subdivide during deformation on a smaller and smaller scale, and that each volume element is characterised by an individual combination of slip systems. A number of microstructural observations (especially of aluminium, nickel, and copper) are described, and dislocation arrangements are discussed on the basis of the general principle that they are low energy dislocation structures. It is shown that the microstructural evolution is quite similar in polycrystalline metals and in single crystals deforming by multislip, and ways in which metallurgical parameters such as stacking fault energy and grain size can affect the microstructure are examined. The general principle of grain subdivision during cold deformation is discussed with reference to the microstructural observations.MST/1290
Plane wave shock loading produces twins or twin faults in many metals and alloys, and, especially for fcc materials with decreasing stacking fault energy (SFE), a critical twinning pressure (CTP), and crystallographic orientations which control the onset and extent of twinning. The CTP increases with increasing SFE, and is lowest in [001] orientations for fcc. The deformation associated with plane shock is generally within the realm of the plastic regime of the stress–strain diagram. Correspondingly, spherical shock, characteristic of impact cratering, produces large strains through solid state flow and sliding of overlapping shear bands composed of dynamically recrystallised grains, and is therefore only partly encompassed in the realm of the stress–strain diagram. Below this zone for impact craters, and similar to plane shock loading, there is a slip–twinning or slip–microbanding transition in fcc materials, which depends on the SFE; twinning persists for low SFE while microbands dominate in high SFE fcc materials such as nickel or copper. A simple model is used to explain differences in residual microstructures in plane shock loading and spherical shock or impact cratering as they relate to SFE and shock wave geometry. Simple schematics are used to explain the connection between plane shock and spherical shock loading in relation to the stress–strain diagram as a paradigm.
The microstructural evolution during polyslip in f.c.c. metals is investigated by the examples of Al, Ni, Ni-Co alloys and an Al-Mg alloy, deformed at room temperature either by rolling or by torsion. The principles governing this evolution appears to be the following: (a) There are differences in the number and selection of simultaneously acting slip systems among neighboring volume elements of individual grains. In any one volume element (called a cell block), the number of slip systems falls short of that required for homogeneous (Taylor) deformation, but groups of neighboring cell blocks fulfil the Taylor criterion collectively. (b) The dislocations are trapped into low-energy dislocation structures in which neighbor dislocations mutually screen their stresses. The microstructural evolution at small strains progresses by the subdivision of grains into cell blocks delineated by dislocation boundaries. These boundaries accommodate the lattice misorientations, which result from glide on different slip system combinations in neighbouring cell blocks. The cell blocks are subdivided into ordinary cells and both cell blocks and cells shrink with increasing strain. All observations appear to be in good accord with the theoretical interpretation. However, some problems remain to be solved quantitatively.
The aim of this study is to work out a microscopic picture of the second stage of creep at intermediate temperatures. The dislocation networks having been described in a previous article, we report here subboundary properties under load during in situ experiments on predeformed specimens, in the high voltage electron microscope. Subboundary formation, migration and destruction take place during creep. The two former processes occur frequently and ensure a constant average subgrain size under constant stress. They also play a key-role during transients. For subboundaries with three coplanar Burgers vectors, and in the case of pure tilt, migration occurs by glide of the dislocation segments. Those with two orthogonal Burgers vectors and equal dislocation spacings remain immobile or are destroyed in the absence of significant climb. Triple junctions are also mobile. Subboundary migration accounts for about 10% of the creep strain. The creep rate is therefore bound to the movement of individual dislocations which will be described in another article.
Stage IV has become the accepted name for that work-hardening stage within which large plastic strains can occur at a very
low, virtually constant work-hardening rate, as exemplified by cold rolling and wire drawing. By contrast, in the preceding
stage III, the work-hardening rate decreases sharply with strain, whereas in the still earlier stage II, the work-hardening
rate is also almost constant but has a high value. The classical paper by Langford and Cohen on drawn iron wire is now recognized
as one of the earliest studies of stage IV. Already in 1970, a detailed theoretical analysis of that work based on the mesh
length theory was presented[2] which has stood the test of time, although in it the Langford and Cohen experiments were considered to represent stage II
on account of the operation of similitude and the almost constant work-hardening rate. The present paper re-examines the 1970
theoretical interpretation in terms of stage IV behavior, which necessitates reinterpretation of stage III. Included in the
present interpretation are more recent insights regarding dislocation behavior in so-called LEDS, low-energy dislocation structures.
It is concluded that stages II and IV differ, because in stage II, cross slip is insignificant, while in stage IV, it is unlimited.
Accordingly, cross slip is gradually established in the course of stage III. However, similitude appears to operate in all
three stages. By extension of the argument regarding stages III and IV, it is seen that stages V and VI could follow, including
similitude, through the establishment of climb.
Plane-wave shock deformation has been shown to produce deformation twins or twin-faults in essentially all metal and alloys. In FCC metals and alloys twinning depends upon stacking-fault free energy (SFE) and a critical twinning pressure; which increases with increasing SFE. For impact cratering where the shock wave is spherical and a prominent deviatoric (shear) stress is involved, metals and alloys with high SFE form microbands coincident with {111} plane traces while low SFE metals and alloys either form mixtures of twins and microbands or microtwins. Oblique shock loading of copper also produces mixtures of twins and microbands. Both microtwins and microbands increase in volume fraction with increasing grain size. BCC iron is observed to twin in both shock loading and as a result of impact cratering. Impact craters, shaped charges, and other examples of extreme deformation and flow at high strain rates exhibit various regimes of shear bands and dynamic recrystallization as a mechanism for solid-state flow. Deformation twins and microbands are also often precursors to this process as well. Examples of these phenomena in FCC materials such as Al, Ni, Cu, stainless steel and brass, and BCC materials such as Fe, W, Mo, W-Ta, and Ta are presented; with emphasis on optical metallography and transmission electron microscopy.
A systematic examination of microbands developed in various materials, including pure Al, Cu, Ag, Nb metals and Al-Mg, 6061 Al and Al-Li-Cu alloys, deformed dynamically or quasi-statically to intermediate strains has been conducted. Based on extensive characterization of microbands using transmission electron microscopy, the characteristics of microbands do not appear to be strongly dependent on crystal structure, material properties, strain level or deformation path. This finding suggests that the formation mechanism of microbands may be similar in a variety of f.c.c. and b.c.c. metals and alloys. A possible microband formation mechanism is proposed, based on a concept of the further development of coarse slip bands (or dislocation tangles on glide planes), which is consistent with experimental observations. The model involves first the generation of polarized dislocations on primary slip systems, followed by an annihilation process for the primary dislocations in the central portion of a band structure, forming double dislocation walls parallel to the primary slip planes. Misorientation inside the double walls is believed to be caused by the directional shear of primary dislocations. Secondary slip is then induced by the internal stresses in the region between the double walls. Finally a stable dislocation configuration is created as a result of the interaction between the primary and secondary dislocations. The proposed model is in agreement with several observed phenomena, including the constancy of microband thickness, tilt axis, shear sense, and uniform shear strain over the cross-section of a microband. The roles of stacking fault energy, solute atoms and precipitates on microband formation are also discussed.
Gross low temperature plastic deformation of metals results from the movement of large numbers of dislocations. This movement is characterized by dislocation-dislocation interaction events statistically distributed in time and space and by the continuing trend of the dislocation density to rearrange into low energy configurations. The various approaches proposed to link flow stress and strain hardening with the evolving substructure may be grouped into families, emphasizing one or the other of those aspects.The first examines possible low energy dislocation configurations and derives the observed flow stress and strain hardening from the characteristics of the proposed dislocation architecture. The second phase relates flow stress evolution to the kinetics of dislocation movement, i.e. the build-up of the substructure as resulting from successive interaction events.The two approaches have been developed independently. It is the aim of the present paper to examine to what extent experimental observations fit into the framework of the existing models (the “kinetic” model being somewhat modified by an addition proposed in this paper) and to what extent the models are complementary in covering different aspects of the one same truth or are mutually exclusive.
The generation of the fine grained, dynamically recrystallized microstructure has been studied in hot rolled copper and α brasses and in 70:30 brass deformed by hot torsion. The new grains, which developed preferentially at grain boundaries and inhomogeneities of deformation, contained none of the deformation features present in unrecrystallized parts of the microstructure. This observation is contrary to theories of dynamic recrystallization which imply that the microstructure contains a spectrum of grains ranging from just recrystallized to severely deformed. Texture studies showed cold rolled textures at low rolling temperatures but in the case of copper the texture became quite flat after rolling at 350 and 425 °C. Flat textures were associated with a minimum grain size. At higher rolling temperatures the textures were again typical of cold rolled material. It is suggested that while normal slip processes are operating at all temperatures there is an additional contribution from grain boundary deformation processes associated with fine grained microstructures. Such processes would account for the absence of normal deformation features from dynamically recrystallized microstructures.
Deformation microtwins characteristic of Neumann bands dominate the residual microstructures beyond a dynamic recrystallization and highly deformed regime below the crater wall of impact craters in polycrystalline bcc iron targets impacted by 3.2 mm diameter iron projectiles at velocities ranging from 0.5 to 3.8 km s−1. Corresponding impact craters in polycrystalline fcc 304L stainless steel targets impacted by 3.2 mm diameter stainless steel projectiles at velocities ranging from 0.5 to 3.9 km s−1 were observed to have a more limited dynamic recrystallization zone at the crater wall followed by a highly deformed transition into a region of mostly microtwins, but with some intermixing of grains containing microbands.