A delay time td to re‐establish steady‐state neck motion is observed if a partially drawn nylon 6–10 monofilament is aged for a time ta at a stress σa below the stress used to propagate the neck in a tensilecreep experiment. The delay time increases as kσatam for long ta, where 0<m<1, as long as the neck is immobile. The factor k increases steadily with the aging temperature Ta below 40°C, the glass transition temperature Tg of amorphous nylon 6–10, but much more rapidly with Ta above Tg. Aging at σa>0 is necessary to observe a delay time and, in fact, the effects of aging at σa can be erased by subsequently aging a short time at zero stress. At aging stresses just below the propagation stress σp, the neck still moves, but at a lower velocity than at σp. In this regime the delay time increases as cp−σa), 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 104 Å3. 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.

1.
I.
Marshall
and
A. B.
Thompson
,
Proc. Roy. Soc. (London)
221A
,
541
(
1954
).
2.
P. I.
Vincent
,
Polymer
1
,
7
(
1960
).
3.
W. Whitney, Sc.D. thesis, Massachusetts Institute of Technology, 1965.
4.
W.
Whitney
and
R. D.
Andrews
,
J. Polymer Sci. Pt. C
16
,
2981
(
1967
).
5.
N.
Brown
and
I. M.
Ward
,
J. Polymer Sc. Pt. A‐2
6
,
607
(
1968
).
6.
R. C. Richards, M.S. thesis, Cornell University, 1970.
7.
D. H.
Ender
and
R. D.
Andrews
,
J. Appl. Phys.
36
,
3057
(
1965
).
8.
D. H.
Ender
,
J. Appl. Phys.
39
,
4877
(
1968
).
9.
M. L.
Williams
and
M. F.
Bender
,
J. Appl. Phys.
36
,
3044
(
1965
).
10.
E. J. Kramer and R. C. Richards (unpublished).
11.
B. N.
Dey
,
J. Appl. Phys.
38
,
4144
(
1967
).
12.
D. A.
Zaukelies
,
J. Appl. Phys.
33
,
2798
(
1962
).
13.
D.
Hansen
and
J. A.
Rusnock
,
J. Appl. Phys.
39
,
322
(
1965
).
14.
B. W.
Cherry
and
C. M.
Holmes
,
Brit. J. Appl. Phys.
2
,
621
(
1969
).
15.
H. W.
Starkweather
, Jr.
and
R. E.
Moynihan
,
J. Polymer Sci.
22
,
363
(
1956
).
16.
I.
Sandeman
and
A.
Keller
,
J. Polymer Sci.
19
,
401
(
1956
).
17.
The reason that one of the necks is usually stationary is that it has stress aged for a far longer length of time than the moving neck after a few stress aging cycles have been completed.
18.
Throughout this paper, stress will reported as the nominal stress, based on the cross‐sectional area of the unoriented monofilament, unless otherwise specified.
19.
L. E. Nielsen, Mechanical Properties of Polymers (Wiley, New York, 1960).
20.
H.
Eyring
,
J. Chem. Phys.
4
,
283
(
1936
).
21.
The shear activation volume defined here is the product of a (microscopic) stress concentration factor, an area on the shear plane and a shear distance.
22.
C.
Bauwens‐Crowet
,
J. C.
Bauwens
, and
G.
Homes
,
J. Polymer Sci. Pt. A‐2
7
,
735
(
1969
).
23.
This sample is the same as the one denoted 12–22 in Table I; V* is taken to be the average of the three determinations of V* for this sample.
24.
F. A. McClintock and A. S. Argon, Mechanical Behavior of Materials (Addison‐Wesley, Reading, Mass., 1966), p. 411.
25.
For example, if the stress concentration factor K is proportional to τpm where m is an exponent greater than zero, the actual shear stress change is proportional to (Δτp)m+1 and the measured activation volume is (m+1) times the real activation volume.
26.
D. H. Ender, MIT Final Tech. Rep. 68‐7‐CM, 1968, p. 137.
27.
W. G.
Johnston
and
J. J.
Gilman
,
J. Appl. Phys.
31
,
687
(
1960
).
28.
W. G.
Johnston
,
J. Appl. Phys.
33
,
2716
(
1962
).
29.
A. H. Cottrell, in Symposium on the Plastic Deformation of Crystalline Solids (Carnegie Institute of Technology and ONR, Pittsburgh, 1950), p. 60.
30.
A. H. Cottrell, Dislocations and Plastic Flow in Crystals (Oxford U.P., Oxford, 1953), p. 139.
31.
E.
Orowan
,
Proc. Phys. Soc. (London)
52
,
8
(
1940
).
32.
J. M.
Peterson
,
J. Appl. Phys.
37
,
4047
(
1966
).
33.
M. L.
Huggins
,
J. Org. Chem.
1
,
407
(
1936
).
34.
Reference 11, Fig. 7. Dey computes a value of 0.3 eV for ΔHeff*, but it is obvious on examination of Fig. 7 that this number is in error. The low (0.3 eV) value of ΔHeff* is one of Dey’s main arguments for his conclusion that dislocation motion is responsible for neck propagation.
35.
S. W.
Lasoski
and
W. H.
Cobbs
, Jr.
,
J. Polymer Sci.
36
,
21
(
1959
).
36.
L.
Pauling
,
J. Amer. Chem. Soc.
67
,
555
(
1945
).
37.
R.
Puffr
and
J.
Sebenda
,
J. Polymer Sci. Pt. C
16
,
79
(
1967
).
38.
J. J. Gilman, in Dislocation Dynamics, edited by A. R. Rosenfield et al. (McGraw‐Hill, New York, 1968), p. 3.
39.
R. E.
Robertson
,
J. Chem. Phys.
44
,
3950
(
1966
).
40.
M.
Goldstein
,
J. Polymer Sci. Pt. B
4
,
87
(
1966
).
41.
T. Ree and H. Eyring, in Rheology, edited by F. R. Eirich (Academic, New York, 1958), Vol. 2, p. 83.
42.
J. S.
Laruzkin
,
J. Polymer Sci.
30
,
595
(
1958
).
43.
R. E.
Robertson
,
J. Appl. Polymer Sci.
7
,
443
(
1963
).
44.
It is probably not true that the secondary flow unit has the same activation volume as the primary unit. In fact, one expects that both V2* and ΔH2* for the secondary flow unit are much smaller than these quantities for primary flow unit. In order to maintain consistency f2 must be defined then as the product of the actual secondary‐unit jump frequency and the number of secondary jumps necessary to produce a shear volume equal to V*, the activation volume of the primary flow unit.
45.
M. H.
Cohen
and
D.
Turnbull
,
J. Chem. Phys.
31
,
1049
(
1958
).
46.
F.
Bueche
,
J. Chem. Phys.
30
,
748
(
1959
).
47.
The secondary jump frequency f2 is not expected to be as strongly stress activated as f1 so that f2a≈f2p. Even if f2a/f2p is strongly stress dependent, it would be as exp[−V2*p−τa)/2kT] and V* would be replaced by (V* − V2*), which makes no qualitative difference to the argument presented.
48.
Strictly speaking, the temperature dependence of nm = vmq/vp should also contribute to the measured ΔH*. The fact that ΔH* is constant in the glass transition range where vm should vary non‐linearly with T implies that its contribution to ΔH* is small compared with the contribution due to ΔH1*.
49.
L. R. G. Treloar, The Physics of Rubber Elasticity (Oxford, U.P., Oxford, 1958), p. 264.
50.
W. A.
Johnson
and
R. R.
Mehl
,
Trans. AIME
135
,
416
(
1939
).
51.
M.
Avrami
,
J. Chem. Phys.
7
,
1103
(
1939
).
52.
M.
Avrami
,
J. Chem. Phys.
8
,
272
(
1940
).
53.
M.
Avrami
,
J. Chem. Phys.
9
,
177
(
1941
).
54.
N.
Saito
,
K.
Okano
,
S.
Iwayanagi
, and
T.
Hideshima
,
Solid State Phys.
14
,
343
(
1963
).
55.
An alternative assumption to the free volume assumption above may appear superficially attractive. Since the number of primary flow units varies, one is tempted to write that the force on a flow unit is proportional to the area between flow units (vpq/n)2/3 multiplied by the macroscopic shear stress rather than the implicit assumption made above that the force on each primary flow unit is independent of 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 τp than actually observed, i.e., it predicts td∝τp3/2 and vn∝τp2.
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