Dislocation substructures in pure aluminium after creep deformation as studied by electron backscatter diffraction (original) (raw)
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Journal of Applied Crystallography, 2023
The peak broadening in neutron diffraction experiments on tensile specimens of pure Al (99.8%) and an Al-Mg alloy pre-deformed at different creep strains is analysed. These results are combined with the kernel angular misorientation of electron backscatter diffraction data from the creep-deformed microstructures. It is found that differently oriented grains possess different microstrains. These microstrains vary with creep strain in pure Al, but not in the Al-Mg alloy. It is proposed that this behaviour can explain the power-law breakdown in pure Al and the large creep strain observed in Al-Mg. The present findings further corroborate a description of the creep-induced dislocation structure as a fractal, predicated on previous work.
Dislocation Substructure Gradient Formation in Aluminum by Creep under Weak Potential
Arabian Journal for Science and Engineering, 2011
Transmission diffraction electron microscopy of thin foils was used to study the dislocation substructure gradient of aluminum destroyed during creep. Creep under +1 V potential resulted in the formation of a dislocation substructure gradient, which was observed as a regular change in quantitative structural characteristics upon moving away from the sample fracture surface.
Alloying effects on dislocation substructure evolution of aluminum alloys
International Journal of Plasticity, 2004
The constitutive response of aluminum alloys is controlled by the evolution of dislocation substructure including mobile and forest dislocation density, cell size distribution and morphology, and misorientation angle between neighboring cells. The present study focuses upon the small strain regime and compares the measured microstructural evolution of 3003, 5005, and 6022 aluminum alloys during deformation. Room temperature tensile deformation experiments were performed on industrially manufactured specimens of each alloy and the evolving microstructure was compared with the mechanical response. The dislocation structure evolution was characterized using transmission electron microscopy and orientation imaging of deformed specimens. It was observed that structural evolution is a function of lattice orientation and the character of neighboring grains. In general, the dislocation cell size and misorientation angle between dislocation cells evolves systematically with deformation at relatively small strain levels. #
Metallurgical and Materials Transactions A, 2002
Creep experiments were conducted on aluminum single crystals and copper polycrystals deformed within the five-power-law regime. The dislocation structure of copper, which has not been extensively characterized in the past, consists of less-well-defined subgrain walls of relatively low misorientation, typically between 0.1 and 0.3 deg, with a Frank network of dislocations within the subgrains. The aluminum, as expected, consisted of well-defined subgrain boundaries with a typical misorientation between 1.0 and 2.0 deg. The subgrains were probed from one boundary to another in copper and aluminum using convergent-beam electron diffraction (CBED). This allowed a determination of any changes in the lattice parameter, which would indicate the presence of any internal stresses. Earlier investigations by others suggested that internal stresses may be high in the vicinity of the "hard" subgrain boundaries in both loaded and unloaded specimens, based on a variety of techniques including X-ray diffraction (XRD), stress-dip tests, as well as some preliminary CBED. It was determined in this work that the lattice parameter was unchanged at the equilibrium or stress-free value within the interior of the subgrains and along (within a one-beam diameter) the subgrain boundaries.
Acta Metallurgica, 1982
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 sour= 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. R&um&-Nous avons itudie le comportement de la sous-structure de dislocations pendant le fluage de I'aluminium aux temperatures moyennes, dans le but de determiner quels sont les mtcanismes microscopiques qui contr&nt la vitesse de deformation. Ce premier article d&it les prop&es gtomttriques des rtseaux de dislocations apres fluage. Lea sous-joints sont constitues par les disiocations qui se sont multipliees sous l'effet de la contrainte locale; ils sont en majorite mixtes, et composes de 3 vecteurs de Burgers coplanaires a 120"; leur champ de contrainte a longue distance. s'il existe, est infirieur ou tga) a la contrainte appliqufe; enfin, ils sont constituts de segments de dislocations situ& dans des plans de glissement ce qui impose certaines restrictions sur la gtometrie des reseaux possibles. Dans les articles suivants, ces resultats permettront de comprendre le comportement dynamique de la sous-structure. au tours d'expi$ences de fluage in situ dam un microscope tlectronique a haute tension, et d'ttablir un nouveau modele de fluage aux temperatures moyennes. zllgmmenfasatmg-Das Verhalten der Substruktur in Aluminium wurde bei mittleren Temperaturen untersucht, urn die mikroskopischen Mechanismen zu bestimmen, die die Dehngeschwindigkeit kontrollieren. Diese erste Arbeit beschreibt die einxelnen geometrischen Eigenschaften der nach dem Kriechversuch beobachteten Versetzungsnetzwerke. Die Subkorngrenxen bestehen aus Versetzungen meist gemischten Charakters, die von Quellen nach deren Aktivierung durch die lokale Spannung emittiert werden. Drei koplanare Bergersvektoren unter 120" treten auf. Das weitreichende Fcld dieser Subkomgrenzenist, wenn iiberhaupt vorhanden, kleiner oder gleich der Kriechspannung. Die kleinen Versetzungsegmente liegen in ihrer Gleitebene, welches die miiglichen Strukturen der Subkorngrenzen einschriinkt. In einer nachfolgenden Arbeit sind diese Ergebnisse wichtig fur das Verstlndnis des dynamischen Verhaltens der Substruktur wahrend der in siru-Experimente im Hochspannungseiektronensikroskop und fur die Erarbeitung eines neuen Kriechmodelles fur mittlere Temperaturen. 93 Al 99.33:: rolling by 75", Al 99.37; annealed ih at 500cC Al 99.9% annealed 1 h at 350°C Al 99.96"/, annealed 2hat3OO"C Al 99.6"/, annealed 2h at 550°C
Dislocation substructure evolution on Al creep under the action of the weak electric potential
Materials Science and Engineering A-structural Materials Properties Microstructure and Processing, 2010
The dislocation substructure evolution on Al creep under the action of the weak electric potential is established by methods of transmission diffraction electron microscopy. It is shown that change of the electrical potential of the Al sample surface is accompanied by the increase of dislocation substructure self-organization degree.
Compact and Dissociated Dislocations in Aluminum: Implications for Deformation
Physical Review Letters, 2005
Atomistic simulations, confirmed by electron microscopy, show that dislocations in aluminum can have compact or dissociated cores. The calculated minimum stress (P) required to move an edge dislocation is approximately 20 times smaller for dissociated than for equivalent compact dislocations. This contradicts the well accepted generalized stacking fault energy paradigm that predicts similar P values for both configurations. Additionally, Frank's rule and the Schmid law are also violated because dislocation core energies become important. These results may help settle a 50-year-old puzzle regarding the magnitude of P in face-centered-cubic metals, and provide new insights into the deformation of ultra-fine-grained metals.
In situ observation of dislocation nucleation and escape in a submicrometre aluminium single crystal
Nature Materials, 2009
Smaller is stronger' does not hold true only for nanocrystalline materials 1 but also for single crystals 2-5. It is argued that this effect is caused by geometrical constraints on the nucleation and motion of dislocations in submicrometre-sized crystals 6,7. Here, we report the first in situ transmission electron microscopy tensile tests of a submicrometre aluminium single crystal that are capable of providing direct insight into sourcecontrolled dislocation plasticity in a submicrometre crystal. Single-ended sources emit dislocations that escape the crystal before being able to multiply. As dislocation nucleation and loss rates are counterbalanced at about 0.2 events per second, the dislocation density remains statistically constant throughout the deformation at strain rates of about 10 −4 s −1. However, a sudden increase in strain rate to 10 −3 s −1 causes a noticeable surge in dislocation density as the nucleation rate outweighs the loss rate. This observation indicates that the deformation of submicrometre crystals is strain-rate sensitive. Deformation of miniaturized single crystals with physical dimensions in the micrometre to submicrometre range consistently shows that plastic deformation of smaller crystals requires higher stress 1-5. In addition, their plastic flow deviates from the continuous strain hardening of bulk crystals, with the emergence of intermittent slip bursts 8. These findings strongly suggest that the crystal dimension affects the motion or nucleation of lattice dislocations at this length scale. On the basis of phenomenological considerations, the 'dislocation starvation' theory has been proposed to account for these effects 3,6. The basic idea is that in small crystals dislocations escape more quickly through free surfaces than they can multiply, leaving the sample free of dislocations. Plastic deformation is thus controlled by dislocation nucleation, which is rarely a continuous process. Simulation studies of discrete dislocation dynamics followed to rationalize the mechanical size effects by taking into account the length scale effects on dislocation plasticity, such as: stochastic distribution of dislocation source lengths 7,9 , dislocation escape through free surfaces (surface annihilation) 9-11 and the truncation of dislocation source operation by free surfaces 10-12. However, no direct observations of dislocation mechanisms during the deformation of (sub)micrometre crystals have been reported, leaving a missing gap between the mechanical test data and the simulation perspectives. Post-mortem transmission electron microscopy (TEM) characterization of deformed crystals may be useful for tracking the evolution of dislocation density with strain 13,14 except for LETTERS NATURE MATERIALS
Materials Science and Engineering: A, 1989
ABSTRACT The detailed arrangement of dislocations in cells and sub-boundaries is ruled by the dislocation energy minimization principle in a recovery process. In order to determine the exact process, the mechanical spectroscopy technique has been chosen because the dislocation structure is subjected to stress in conditions where irreversible changes do not appear. A relaxation peak appears at around 0.45Tm(T m is the melting temperature) but only if cell walls are present. It disappears in the well-organized dislocation structure as occurs after creep tests. This peak is used to follow the reorganization of the dislocation during thermomechanical treatments.