Velocities of individual dislocations have been measured in LiF, covering a range of twelve orders of magnitude in velocity, from 10−7 cm/sec to 105 cm/sec. The velocity is extremely sensitive to applied stress at low velocities, and for each crystal there exists a minimum stress for dislocation motion, below which dislocations do not move. The edge components of dislocation loops move considerably faster than the screw components. The upper limit for dislocation velocity appears to be the velocity of sound in the crystal. The effects of temperature, impurities, and radiation damage on dislocation velocity are described. These variables affect the dynamic resistance to motion encountered by a moving glide dislocation. The growth of total dislocation density, the growth of individual glide bands, and the distribution of glide dislocations during plastic deformation are described. The yield stress of LiF is determined by the resistance to motion encountered by a glide dislocation in moving through an otherwise dislocation-free region of the crystal. The yield stress is independent of the dislocations present in an undeformed crystal, and the state of pinning and geometrical arrangement of such dislocations do not affect the yield stress. Stress-strain curves have been calculated from the data on dislocation mobility and dislocation density, and the calculated and measured curves are compared. At low strains the flow stress can be predicted from measured dislocation properties.

Stress-strain curves and transient creep curves for single crystals are obtained from calculations based on the observed behavior of dislocations in LiF, and the strain rate equation, ε̇=bnv. The pronounced yield drops and apparent ``delay times'' predicted by the calculations are observed experimentally. This agreement between calculation and experiment implies that further consideration of the calculations may give some insight into the yield point and transient creep behavior of single crystals. Several parameters are varied in the calculation in order to determine the effect of such things as testing machine speed, machine hardness, rate of dislocation multiplication, work-hardening rate, number of mobile dislocations initially present, and the dependence of dislocation velocity on stress. The parameter that most strongly influences the yield points and transient creep behavior of a given material is n0, the number of mobile dislocations initially present. The wide range of observed yield point and transient creep behavior of various materials can be rationalized in terms of how dislocation velocity varies with stress; the less sensitive is the velocity to the applied stress, the more pronounced will be the yield drop or delay time. The calculation is applicable only when n0≠0; and the yield points and yield drops develop continuously and relatively gradually as dislocation motion and multiplication begin before the upper yield point is reached. The upper yield stress is not closely related to the stress required to unpin dislocations. It is necessary to distinguish between this type of yield point, which is commonly observed, and the Cottrell type of yield point in which dislocation motion begins at the upper yield stress, and in which the upper yield stress is related to the unpinning stress.

Experimental observations are presented of dislocation multiplication, of the defect structure left behind by a moving dislocation, and of cross-glide of individual dislocations in LiF crystals. New dislocation loops form at many different sites in the wake of a moving dislocation. These loops have the same Burgers vector as the parent dislocation but do not, in general, lie on the same atomic plane. The rate of formation of new loops depends upon the magnitude of the applied stress. Such creation of new loops leads eventually to the formation of a wide glide band. A moving screw dislocation trails many line defects behind it that lie parallel to its direction of motion. The existence and nature of these trails and the observed dislocation multiplication can be explained in terms of a mechanism which involves the formation, by cross-glide, of jogs on a screw dislocation. This cross-glide multiplication mechanism was originally proposed by Orowan and by Koehler. It is demonstrated that cross glide occurs easily in LiF, so that this mechanism is plausible. Some interesting complications arise when jogs are formed that are longer than several atomic spacings but less than several hundred. The defect trails exert a dragging of the screw dislocations that is not negligible compared to the yield stress of a crystal.

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Dislocation etch pits can be formed on LiF by a dilute aqueous solution of FeF 3 • In this report the etch pit formation is described in detail, and the mechanism for pit formation is discussed. The nature of the etch pits depends on the character of the dislocation, and on the exact composition of the etchant. Edge dislocations and screw dislocations etch slightly differently; the former produce deeper pits. The etching is inhibited by some segregated impurities at dislocations, therefore aged dislocations and fresh dislocations etch much differently. Etch pit formation is probably due to the preferential nucleation of' unit pits one molecule deep at a dislocation, and the movement of the monomolecular steps across the surface. The relative rates of these two processes determine the shape of the etch pits. The nucleation rate for unit pits depends upon the dislocation energy, hence upon the character of the dislocation and the impurity content as suggested by Cabrera. The nucleation rate is faster at edge dislocations, because of their higher energy. The nucleate rate is low at dislocations with segregated impurities, because the impurities lower the dislocation energy. The ferric ion is adsorbed on the surface and inhibits the motion of steps, so that steeper, more visible pits are produced as the iron content is increased.