# Dislocation > *For other uses, see Dislocation (disambiguation).* A **dislocation** is a line defect in a [[Crystal_structure|crystal]]: the boundary between a region where the atoms have slipped by one lattice spacing and a region where they have not. It is the [[Crystallographic_defect|crystallographic defect]] that makes metals ductile. Shearing whole planes of a perfect crystal past one another at once would take several gigapascals in [[Copper|copper]]; a real copper crystal yields at about a megapascal, ten thousand times less, because a dislocation lets the plane slip one row of atoms at a time.[^frenkel1926][^hull-bacon-ch1] The slip a dislocation carries is its [[Burgers_vector|Burgers vector]] `b`; an edge dislocation has `b` perpendicular to its line and is an extra half-plane of atoms wedged into the crystal, a screw dislocation has `b` parallel to its line and turns the lattice planes into a helical ramp.[^hull-bacon-ch1] In the microsim below the reader raises the tensile stress `σ` on a single crystal of copper or aluminium. The sim resolves it onto the slip plane and slip direction by Schmid's law, `τ_R = σ·cos(φ)·cos(λ)`, and the instant `τ_R` reaches the critical resolved shear stress, 0.5 MPa for the copper preset and 1 MPa for aluminium, an edge dislocation glides across the three-dimensional lattice one Burgers vector per step.[^callister-ch7] A `when` layer swaps the perfect lattice for one carrying thermal [[Vacancy_defect|vacancies]]. On the [[Materials_science]] flagship this page serves Part II, Fundamentals › Structure, in the section *Defects and dislocations*, between diffraction and microstructure. ## History The elastic theory came first. In 1907 Vito Volterra worked out the elastic fields of the *distorsioni* of a cut and rejoined cylinder, which include the edge and screw dislocations, decades before anyone suspected they existed in crystals.[^volterra1907] In 1926 Yakov Frenkel estimated the stress needed to slide one plane of a perfect lattice rigidly over its neighbour as roughly `G/(2π)`, thousands of times larger than the stresses at which real crystals deform.[^frenkel1926] The resolution came in 1934 in three independent papers by Egon Orowan, Michael Polanyi and G. I. Taylor: crystals contain line defects along which slip is already partly done, and plastic flow is the motion of these lines at a small fraction of the theoretical stress; Taylor also showed that their mutual stresses explain [[Work_hardening|work hardening]].[^taylor1934][^orowan1934][^polanyi1934] J. M. Burgers added the screw dislocation and his vector in 1939; Frank and Read, and Peach and Koehler, followed in 1950.[^burgers1939][^frank-read1950][^peach-koehler1950] Direct evidence came in 1956, when Hirsch, Horne and Whelan in Cambridge and Walter Bollmann in Geneva watched dislocations move in thin foils in the transmission electron microscope.[^hirsch1956][^bollmann1956] ## Mechanisms Plastic flow needs dislocations to be created, to arrange themselves and to move. Those three questions are the mechanics behind the sim's single gliding line. ### Generating dislocations An annealed metal contains of the order of 10⁹ metres of dislocation line per cubic metre, a heavily deformed one 10¹⁵ to 10¹⁶: deformation manufactures dislocations.[^callister-ch4] The standard source is the Frank–Read mechanism: a segment of length `L` pinned at both ends bows out under a resolved shear stress `τ` against its line tension, an [[Elastic_energy|elastic energy]] of roughly `½·G·b²` per unit length.[^frank-read1950][^hull-bacon-ch8] The bowing becomes unstable when the radius of curvature falls to `L/2`, at `τ_FR ≈ G·b/L`; the loop then pinches off a complete ring and the segment repeats.[^hull-bacon-ch8] For copper, with `G ≈ 46 GPa` and `b = 0.256 nm`, a segment 1 µm long operates at about 12 MPa and one 100 nm long at 120 MPa, so the strength of a metal is set by the length of the free segments in it, which is what [[Precipitation_hardening|precipitation hardening]] shortens.[^callister-ch6] Dislocation-free [[Silicon|silicon]] grown by the [[Czochralski_method|Czochralski method]], having no sources at all, is correspondingly strong. ### Interaction and arrangement A dislocation's stress field falls off as `1/r`, so dislocations feel one another across hundreds of atomic spacings. Like-signed dislocations on the same plane repel and pile up against obstacles such as a [[Grain_boundary|grain boundary]], which is the origin of the grain-size dependence of yield strength; opposite-signed ones attract and annihilate, and like-signed edges on parallel planes settle one above another into walls, the low-angle boundaries of a [[Microstructure|microstructure]].[^hull-bacon-ch9] As the density `ρ` rises the dislocations obstruct each other and the flow stress follows the Taylor relation `τ = α·G·b·sqrt(ρ)`, with `α` about 0.2–0.5.[^taylor1934][^hull-bacon-ch10] Taking `α = 0.3` for copper gives 0.35 MPa at `ρ = 10¹⁰ m⁻²`, 35 MPa at 10¹⁴ and 110 MPa at 10¹⁵, the whole span from annealed to [[Cold_working|cold-worked]] copper, and [[Annealing_(materials_science)|annealing]] removes them again. ### Movement The sim's subject is glide, motion of a dislocation in the plane containing its line and its Burgers vector. In face-centred cubic metals such as copper and [[Aluminium|aluminium]] glide happens on the four {111} planes in the three <110> directions of each, the twelve slip systems that [[Slip_(materials_science)|slip]] can choose from.[^callister-ch7] A tensile stress `σ` along the crystal axis acts on a slip system only through its resolved shear component, which Schmid's law gives as `τ_R = σ·cos(φ)·cos(λ)`, where `φ` is the angle between the axis and the slip-plane normal and `λ` the angle between the axis and the slip direction.[^schmid-boas1935][^callister-ch7] The Schmid factor `cos(φ)·cos(λ)` cannot exceed 0.5, reached when both angles are 45°; slip begins on the system with the largest factor as soon as `τ_R` reaches the critical resolved shear stress `τ_CRSS`, of order 1 MPa for pure, well-annealed face-centred cubic crystals.[^callister-ch7][^hull-bacon-ch10] In the sim the control is `σ` and the readout `τ_R` on the most favourable system. For a crystal pulled along [001] each of the eight active systems has `cos(φ) = 1/√3` and `cos(λ) = 1/√2`, a Schmid factor of 0.408, so the copper preset, with `τ_CRSS = 0.5 MPa`, yields when `σ` reaches 1.23 MPa and the aluminium preset, with 1 MPa, at 2.45 MPa. Below the threshold nothing moves; above it the edge dislocation steps one `b` at a time, and each dislocation that leaves the crystal displaces the two halves by one Burgers vector, so a visible slip step is the passage of thousands. The preset `τ_CRSS` values are the row's presets rather than measurements on a particular crystal.[^crss-cn] The sim omits the [[Peierls_stress|Peierls stress]], the lattice's own resistance to the moving line, negligible for the wide dislocations of face-centred cubic metals and large for the narrow ones of body-centred cubic [[Iron|iron]] or covalent silicon.[^peierls1940][^nabarro1947] It also omits climb, in which an edge dislocation leaves its glide plane by absorbing or emitting vacancies; climb needs [[Diffusion|diffusion]] and matters only above roughly 0.4 of the melting temperature, where it controls [[Creep_(deformation)|creep]].[^hull-bacon-ch3] The vacancies exist in thermal equilibrium at a fraction `n_v/N = exp(−Q_v/(k·T))`, the [[Boltzmann_distribution|Boltzmann]] factor of the formation energy `Q_v`; for copper, with `Q_v = 0.9 eV`, that is about 10⁻¹⁵ at 300 K and a few times 10⁻⁴ near the melting point.[^likharev-sm2][^callister-ch4] ## Geometry The Burgers vector is defined by a circuit: a closed loop of lattice steps drawn atom to atom around the dislocation fails to close when repeated in a perfect crystal, and the closure failure is `b`.[^hull-bacon-ch1] For a full dislocation `b` is a lattice translation, `a/2<110>` in face-centred and `a/2<111>` in body-centred [[Cubic_crystal_system|cubic]] metals. The angle between `b` and the line classifies the dislocation. ### Edge In an edge dislocation `b` is perpendicular to the line. The picture is an extra half-plane of atoms inserted into the upper half of the crystal and ending at the line; above it the lattice is in compression, below it in tension, and the elastic energy of that field, larger than a screw's by the factor `1/(1 − ν)` with `ν` the [[Poisson's_ratio|Poisson's ratio]], is about `½·G·b²` per unit length: for copper 1.5 nJ per metre of line, or 2.4 eV per atomic spacing, which is why dislocations never form in thermal equilibrium.[^hull-bacon-ch4] The line and `b` fix one plane, so an edge dislocation has a unique glide plane and leaves it only by climb, which is why the sim's line is an edge. In the sim each glide step moves the half-plane's termination one atomic row along the slip plane, so the half-plane walks through the crystal although no atom moves more than a fraction of `b`. ### Screw In a screw dislocation `b` is parallel to the line. There is no extra half-plane; a circuit around the line climbs by one `b`, so the lattice planes normal to the line form a single helicoid. Every plane containing the line also contains `b`, so a screw dislocation has no unique glide plane and can cross-slip to another plane when the first is blocked, a freedom edge dislocations lack.[^hull-bacon-ch3] A surface step ending at a screw dislocation never grows out, which makes the screw a mechanism of [[Crystal_growth|crystal growth]].[^hull-bacon-ch1] ### Mixed A dislocation of general character has `b` at an angle `θ` to its line, and its field is the superposition of an edge component `b·sin(θ)` and a screw component `b·cos(θ)`. Since `b` is constant along a dislocation while the line can curve, a single loop is edge at two points, screw at two others and mixed elsewhere.[^hull-bacon-ch1] ### Partial In close-packed metals a full dislocation lowers its energy by splitting into two partials, whose Burgers vectors are not lattice translations, joined by a ribbon of [[Stacking_fault|stacking fault]]. The face-centred cubic reaction is `a/2[1̄10] → a/6[2̄11] + a/6[1̄21̄]`, into two Shockley partials; Frank's rule, that a reaction is favourable when the sum of the squares of the Burgers vectors falls, gives `a²/2 → a²/6 + a²/6 = a²/3`, and the partials repel until the fault ribbon between them, of energy `γ` per unit area, balances the repulsion.[^hull-bacon-ch5] Copper, with a low stacking-fault energy near 45 mJ/m², has widely split dislocations that cross-slip with difficulty; aluminium, near 160 mJ/m², has narrow ones that cross-slip easily.[^hull-bacon-ch5] ### Stair-rod and the Lomer–Cottrell junction When two dissociated dislocations gliding on intersecting {111} planes meet along the line of intersection, their leading partials can combine: `a/6[21̄1̄]` on (111) and `a/6[1̄21]` on (111̄) sum to `a/6[110]`, a stair-rod partial lying in neither plane. Frank's rule makes the reaction favourable, `a²/6 + a²/6 → a²/18`, and the product is sessile: the configuration, a Lomer–Cottrell lock, sits across both glide planes as a barrier to everything that follows.[^hull-bacon-ch5] Such locks account for the rapid linear hardening of the second stage of a single-crystal [[Stress–strain_curve|stress–strain curve]].[^hull-bacon-ch10] ### Jog When one dislocation cuts through another, each acquires a step equal to the other's Burgers vector. A step that takes the line out of its glide plane is a jog. On a screw dislocation a jog is a short edge segment whose glide plane is not the screw's, so it can keep up only by climb, creating or absorbing a vacancy at every step; a jogged screw therefore leaves a trail of point defects, one of the main ways deformation produces vacancies.[^hull-bacon-ch7] ### Kink A step that stays within the glide plane is a kink. Kinks are how a dislocation moves against a high Peierls barrier: a pair of kinks nucleates on the line and runs apart, advancing it by one row without the whole line jumping at once.[^hull-bacon-ch3] Kink-pair [[Nucleation|nucleation]] is thermally activated, which is why the yield stress of body-centred cubic metals such as iron and [[Steel|steel]] rises steeply as the [[Temperature|temperature]] falls.[^hull-bacon-ch3] ## Example in two dimensions (2D) In two dimensions dislocations can be seen whole. In 1947 Lawrence Bragg and John Nye floated a raft of equal soap bubbles, about a millimetre across, on a soap solution; the bubbles pack into a hexagonal lattice, and the sheared raft shows edge dislocations, grain boundaries and slip.[^bragg-nye1947] In a two-dimensional triangular lattice an edge dislocation is a bound pair of a five- and a seven-coordinated site, the two ends of the extra half-row, as seen in [[Colloid|colloidal]] monolayers. Two-dimensional crystals also melt through their dislocations: in the Kosterlitz–Thouless–Halperin–Nelson–Young theory, thermally created dislocation pairs unbind at a [[Phase_transition|transition]] that destroys translational but not orientational order, giving an intermediate hexatic phase between solid and liquid.[^halperin-nelson1978][^young1979] ## Observation Dislocations are too small for light and too rare for [[X-ray_crystallography|X-ray crystallography]] to average over, so seeing them needs a microscope that resolves the strain field or a chemical mark where a line meets a surface.[^hull-bacon-ch2] ### Transmission electron microscopy (TEM) The transmission electron microscope images a foil thinned to about 100 nm. Near a dislocation the planes are bent, so the local Bragg condition changes and the line shows as a dark or bright streak; because the contrast comes from the displacement field, a dislocation vanishes when the diffracting planes are not distorted by it, the `g·b = 0` invisibility criterion that identifies `b`.[^hull-bacon-ch2] Weak-beam imaging resolves the two partials of a dissociated dislocation, and high-resolution modes image the atomic columns, so the extra half-plane of an edge dislocation can be seen directly.[^hull-bacon-ch2] In-situ straining stages let dislocations be filmed moving, cross-slipping and multiplying.[^hirsch1956] ### Other methods Where a dislocation meets a surface the strained material dissolves faster, and a suitable etchant opens a pit at each intersection; counting etch pits gave the first dislocation densities and, on lithium fluoride, the first dislocation velocities.[^hull-bacon-ch2] X-ray topography images individual dislocations in nearly perfect crystals such as [[Semiconductor|semiconductor]] wafers, and electron channelling contrast in the scanning electron microscope now images them in bulk samples.[^hull-bacon-ch2] ## Dislocation forces The mechanics of a dislocation's motion is written as a [[Force|force]] per unit length of line: the derivative of the work the applied stress does when the line moves, not a force on any atom. ### Forces on dislocations Peach and Koehler showed that a dislocation with Burgers vector `b` and unit line direction `ξ` in a stress field `σ` feels a force per unit length `F = (σ·b) × ξ`.[^peach-koehler1950] For glide only the shear stress `τ` resolved on the glide plane along `b` matters, and the formula collapses to `F = τ·b`, normal to the line in the glide plane; a normal stress across the extra half-plane gives instead a climb force `σ_n·b`.[^hull-bacon-ch4] The numbers are small: at the copper preset's yield stress, `τ = 0.5 MPa` gives 1.3 × 10⁻⁴ N per metre of line, which suffices because the line has almost no inertia and, in a face-centred cubic metal, almost no lattice friction. ### Forces between dislocations Two parallel screw dislocations at separation `r` exert on each other a radial force per unit length `F = G·b²/(2π·r)`, repulsive for like signs and attractive for opposite ones.[^hull-bacon-ch4] For copper at `r = 100 nm` it is 4.8 × 10⁻³ N/m, equivalent to a shear stress of about 19 MPa, forty times the applied stress at which the crystal yields, so dislocations arrange themselves under their own stresses far more than under the load. Parallel edge dislocations on planes a distance `y` apart have a glide force `F_x = (G·b²/(2π·(1 − ν)))·x·(x² − y²)/(x² + y²)²`, which vanishes at `x = 0` and at `x = ±y`: like-signed edges are in stable equilibrium directly above one another, the configuration of a tilt wall, and unstable at 45°, while opposite-signed ones settle at 45° into dipoles.[^hull-bacon-ch4] These forces build the walls, cells and pile-ups of a deformed metal, and with them its hardening.[^hull-bacon-ch10] ### Free surface forces A free surface cannot support stress, so a dislocation near one is attracted to it as if by an image dislocation of opposite sign at the same depth on the far side; for a screw at depth `d` the force is `F = G·b²/(4π·d)`, and a dislocation that reaches the surface leaves a step and is gone.[^hull-bacon-ch4] The image force is why dislocations escape from thin foils and why the [[Strength_of_materials|strength]] of micrometre-sized pillars and whiskers rises as their diameter shrinks.[^hull-bacon-ch4] *See also:* [[Crystallographic_defect]] — the vacancy fraction `n_v/N = exp(−Q_v/(k·T))`, the sim's `when` layer · [[Burgers_vector]] · [[Slip_(materials_science)]] — the twelve face-centred cubic slip systems · [[Vacancy_defect]] · [[Frank–Read_source]] · [[Stacking_fault]] · [[Peierls_stress]] · [[Work_hardening]] · [[Grain_boundary]] · [[Strength_of_materials]] ## References [^hull-bacon-ch1]: Hull, D.; Bacon, D. J. *Introduction to Dislocations*, 5th ed. (2011), Butterworth-Heinemann, Ch. 1 Defects in Crystals (the Burgers circuit, edge, screw and mixed dislocations, the theoretical shear strength) (page to pin). [^hull-bacon-ch2]: Hull & Bacon (2011), Ch. 2 Observation of Dislocations (etch pits, TEM diffraction contrast and the `g·b = 0` criterion, weak-beam, X-ray topography, field-ion microscopy) (page to pin). [^hull-bacon-ch3]: Hull & Bacon (2011), Ch. 3 Movement of Dislocations (glide, climb, cross-slip, the Peierls stress, kinks) (page to pin). [^hull-bacon-ch4]: Hull & Bacon (2011), Ch. 4 Elastic Properties of Dislocations (line energy, the Peach–Koehler force, forces between dislocations, image forces) (page to pin). [^hull-bacon-ch5]: Hull & Bacon (2011), Ch. 5 Dislocations in Face-centred Cubic Metals (Shockley and Frank partials, Frank's rule, stacking-fault energies, the Lomer–Cottrell lock) (page to pin). [^hull-bacon-ch7]: Hull & Bacon (2011), Ch. 7 Jogs and the Intersection of Dislocations (page to pin). [^hull-bacon-ch8]: Hull & Bacon (2011), Ch. 8 Origin and Multiplication of Dislocations (the Frank–Read source, `τ ≈ G·b/L`) (page to pin). [^hull-bacon-ch9]: Hull & Bacon (2011), Ch. 9 Dislocation Arrays and Crystal Boundaries (pile-ups, tilt walls) (page to pin). [^hull-bacon-ch10]: Hull & Bacon (2011), Ch. 10 Strength of Crystalline Solids (critical resolved shear stress, Taylor hardening, the stages of single-crystal hardening) (page to pin). [^callister-ch4]: Callister, W. D.; Rethwisch, D. G. *Materials Science and Engineering: An Introduction*, 9th ed. (2014), Ch. 4 Imperfections in Solids: Example Problem 4.1 (copper, `Q_v = 0.9 eV/atom`) and §4.5 on dislocation densities of 10³ mm⁻² in carefully solidified crystals to 10⁹–10¹⁰ mm⁻² in heavily deformed metals (page to pin). [^callister-ch6]: Callister & Rethwisch (2014), Ch. 6 Mechanical Properties of Metals, Table 6.1 (copper shear modulus 46 GPa, Poisson's ratio 0.34) and Ch. 3 Table 3.1 (copper atomic radius 0.1278 nm, giving `a = 0.3615 nm` and `b = a/√2 = 0.256 nm`) (page to pin). [^callister-ch7]: Callister & Rethwisch (2014), Ch. 7 Dislocations and Strengthening Mechanisms, §7.4–7.6: slip systems (Table 7.1, twelve {111}<110> systems in FCC), Schmid's law `τ_R = σ·cos φ·cos λ`, and the critical resolved shear stress (page to pin). [^likharev-sm2]: Likharev, K. *Essential Graduate Physics, Part SM: Statistical Mechanics* (2013), Ch. 2 Principles of Physical Statistics, pp. 29–72 (the Gibbs/Boltzmann distribution `exp(−E/(k·T))` behind the vacancy fraction) (page to pin). https://open.umn.edu/opentextbooks/textbooks/part-sm-statistical-mechanics [^averill-ch12]: Averill, B.; Eldredge, P. *General Chemistry: Principles, Patterns, and Applications* (2011), Ch. 12 Solids, §12.4 Defects in Crystals (point, line and plane defects; dislocations and the ductility of metals) (page to pin). https://open.umn.edu/opentextbooks/textbooks/general-chemistry-principles-patterns-and-applications [^frenkel1926]: Frenkel, J. (1926). "Zur Theorie der Elastizitätsgrenze und der Festigkeit kristallinischer Körper." *Zeitschrift für Physik* 37: 572–609 (the theoretical shear strength `≈ G/(2π)`). [^volterra1907]: Volterra, V. (1907). "Sur l'équilibre des corps élastiques multiplement connexes." *Annales scientifiques de l'École Normale Supérieure* 24: 401–517. [^taylor1934]: Taylor, G. I. (1934). "The Mechanism of Plastic Deformation of Crystals. Part I. Theoretical." *Proceedings of the Royal Society A* 145 (855): 362–387. https://doi.org/10.1098/rspa.1934.0106 [^orowan1934]: Orowan, E. (1934). "Zur Kristallplastizität. I–III." *Zeitschrift für Physik* 89: 605–659. [^polanyi1934]: Polanyi, M. (1934). "Über eine Art Gitterstörung, die einen Kristall plastisch machen könnte." *Zeitschrift für Physik* 89: 660–664. [^burgers1939]: Burgers, J. M. (1939). "Some considerations on the fields of stress connected with dislocations in a regular crystal lattice." *Proceedings of the Koninklijke Nederlandse Akademie van Wetenschappen* 42: 293–325 and 378–399. [^frank-read1950]: Frank, F. C.; Read, W. T. (1950). "Multiplication Processes for Slow Moving Dislocations." *Physical Review* 79 (4): 722–723. https://doi.org/10.1103/PhysRev.79.722 [^peach-koehler1950]: Peach, M.; Koehler, J. S. (1950). "The Forces Exerted on Dislocations and the Stress Fields Produced by Them." *Physical Review* 80 (3): 436–439. https://doi.org/10.1103/PhysRev.80.436 [^hirsch1956]: Hirsch, P. B.; Horne, R. W.; Whelan, M. J. (1956). "Direct observations of the arrangement and motion of dislocations in aluminium." *Philosophical Magazine* 1 (7): 677–684. [^bollmann1956]: Bollmann, W. (1956). "Interference Effects in the Electron Microscopy of Thin Crystal Foils." *Physical Review* 103 (5): 1588–1589. https://doi.org/10.1103/PhysRev.103.1588 [^schmid-boas1935]: Schmid, E.; Boas, W. *Kristallplastizität mit besonderer Berücksichtigung der Metalle* (1935), Springer, Berlin (the resolved-shear-stress law). [^peierls1940]: Peierls, R. (1940). "The size of a dislocation." *Proceedings of the Physical Society* 52 (1): 34–37. [^nabarro1947]: Nabarro, F. R. N. (1947). "Dislocations in a simple cubic lattice." *Proceedings of the Physical Society* 59 (2): 256–272. [^bragg-nye1947]: Bragg, L.; Nye, J. F. (1947). "A Dynamical Model of a Crystal Structure." *Proceedings of the Royal Society A* 190 (1023): 474–481. [^halperin-nelson1978]: Halperin, B. I.; Nelson, D. R. (1978). "Theory of Two-Dimensional Melting." *Physical Review Letters* 41 (2): 121–124. https://doi.org/10.1103/PhysRevLett.41.121 [^young1979]: Young, A. P. (1979). "Melting and the vector Coulomb gas in two dimensions." *Physical Review B* 19 (4): 1855–1866. https://doi.org/10.1103/PhysRevB.19.1855 [^crss-cn]: *Citation needed.* The preset critical resolved shear stresses (copper 0.5 MPa, aluminium 1 MPa) are the values supplied by the M7 sim row. Hull & Bacon Ch. 10 and Callister Ch. 7 give the order of magnitude (about 1 MPa for pure, annealed FCC single crystals); a tabulated measurement for high-purity copper and aluminium single crystals with its source would settle the presets. ## External links - The Wikipedia pair's *External links* section lists the current dislocation-visualisation and simulation sites; none is reproduced here until its URL has been checked. - Averill & Eldredge, *General Chemistry*, Ch. 12 §12.4, and Likharev, *Part SM*, Ch. 2, are on the Open Textbook Library (links in the footnotes above).[^averill-ch12] <!-- MATTERSIM:BEGIN g33 — Matter & Energy Cluster microsim (framework build, specs/sims/Dislocation.json); do not hand-edit inside --> **Microsim — three.js (Wikitube framework):** *Dislocation* <div class="wt-sim" data-src="https://wikitube-3d-microsims.netlify.app/matter/Dislocation.html" data-title="Dislocation"></div> *Built from `MICROSIM_GUIDE/specs/sims/Dislocation.json`; part of the [[PORTAL_Matter|Matter portal]] spine (section sims and See-also variants).* <!-- MATTERSIM:END --> ## Wikipedia : Wikitube **Strict pair:** [Wikipedia](https://en.wikipedia.org/wiki/Dislocation) : [Wikitube](https://en.wikitube.io/wiki/Dislocation) · pinned revision [1368483965](https://en.wikipedia.org/w/index.php?oldid=1368483965) · 2026-09-11 ## Previous hub tags Hubs: `Life_Physics`. Portals: [[PORTAL_Materials_science]]. --- *Matter & Energy Cluster child articles, wave 1 · 2026-09-11 · drafted · Materials_science row M7 · sim pending (matter/Dislocation).*