A delay time *t _{d}* to re‐establish steady‐state neck motion is observed if a partially drawn nylon 6–10 monofilament is aged for a time

*t*at a stress σ

_{a}_{a}below the stress used to propagate the neck in a tensilecreep experiment. The delay time increases as

*k*σ

*for long*

_{a}t_{a}^{m}*t*, where 0<

_{a}*m*<1, as long as the neck is immobile. The factor

*k*increases steadily with the aging temperature

*T*below 40°C, the glass transition temperature

_{a}*T*of amorphous nylon 6–10, but much more rapidly with

_{g}*T*above

_{a}*T*. Aging at σ

_{g}_{a}>0 is necessary to observe a delay time and, in fact, the effects of aging at σ

_{a}can be erased by subsequently aging a short time at zero stress. At aging stresses just below the propagation stress σ

_{p}, the neck still moves, but at a lower velocity than at σ

_{p}. In this regime the delay time increases as

*c*(σ

_{p}−σ

_{a}), where

*c*is markedly increased by decreases in temperature. It is proposed that stress aging is caused by the stress‐induced formation of small regions of better interchain packing (microcrystals) in nominally amorphous portions of the nylon and that the delay time is the time necessary to break up the microcrystalline structure induced by stress aging into the structure characteristic of steady‐state flow. The breakup process is thermally activated with an activation enthalpy of 3.85 eV and stress activated with a shear activation volume of approximately 10

^{4}Å

^{3}. A modified Eyring model of flow is developed which assumes that the microcrystals are primary flow units which use up free volume in the amorphous regions as they are formed. This model accounts qualitatively for all, and quantitatively for some, of the features of the dependence of delay time on propagation and aging variables. Other theories, in particular free‐volume collapse and dislocation dynamics theories, cannot be reconciled with all the observed characteristics of stress aging.

## REFERENCES

*Mechanical Properties of Polymers*(Wiley, New York, 1960).

*Mechanical Behavior of Materials*(Addison‐Wesley, Reading, Mass., 1966), p. 411.

*K*is proportional to $\tau pm$ where

*m*is an exponent greater than zero, the actual shear stress change is proportional to $(\Delta \tau p)m+1$ and the measured activation volume is $(m+1)$ times the real activation volume.

*Symposium on the Plastic Deformation of Crystalline Solids*(Carnegie Institute of Technology and ONR, Pittsburgh, 1950), p. 60.

*Dislocations and Plastic Flow in Crystals*(Oxford U.P., Oxford, 1953), p. 139.

*Dislocation Dynamics*, edited by A. R. Rosenfield et al. (McGraw‐Hill, New York, 1968), p. 3.

*Rheology*, edited by F. R. Eirich (Academic, New York, 1958), Vol. 2, p. 83.

*T*implies that its contribution to $\Delta H*$ is small compared with the contribution due to $\Delta H1*.$

*The Physics of Rubber Elasticity*(Oxford, U.P., Oxford, 1958), p. 264.

*n*. It is true that this alternative hypothesis would give rise to an initial jump frequency after aging that is smaller than the steady-state jump frequency, but it also predicts a much weaker dependence of $td$ and $vn$ on $\tau p$ than actually observed, i.e., it predicts $td\u221d\tau p3/2$ and $vn\u221d\tau p2.$