Front Cover
 Table of Contents
 List of Tables
 List of Figures
 Previous investigations
 Experimental procedures
 Experimental results
 Conclusions and observations
 Biographical sketch
 Back Cover

Title: Strain aging in nickel 200
Full Citation
Permanent Link: http://ufdc.ufl.edu/UF00097522/00001
 Material Information
Title: Strain aging in nickel 200
Physical Description: xii, 144 leaves : ill. ; 28 cm.
Language: English
Creator: Cribb, Walter Raymond, 1949-
Publication Date: 1975
Copyright Date: 1975
Subject: Nickel   ( lcsh )
Crystals -- Defects   ( lcsh )
Dislocations in metals   ( lcsh )
Materials Science and Engineering thesis Ph. D
Dissertations, Academic -- Materials Science and Engineering -- UF
Genre: bibliography   ( marcgt )
non-fiction   ( marcgt )
Thesis: Thesis--University of Florida.
Bibliography: Bibliography: leaves 135-143.
Additional Physical Form: Also available on World Wide Web
General Note: Typescript.
General Note: Vita.
Statement of Responsibility: by Walter Raymond Cribb.
 Record Information
Bibliographic ID: UF00097522
Volume ID: VID00001
Source Institution: University of Florida
Holding Location: University of Florida
Rights Management: All rights reserved by the source institution and holding location.
Resource Identifier: alephbibnum - 000171010
oclc - 02943131
notis - AAT7431


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Table of Contents
    Front Cover
        Page i
        Page i-a
        Page ii
        Page iii
    Table of Contents
        Page iv
        Page v
        Page vi
    List of Tables
        Page vii
    List of Figures
        Page viii
        Page ix
        Page x
        Page xi
        Page xii
        Page 1
        Page 2
        Page 3
    Previous investigations
        Page 4
        Page 5
        Page 6
        Page 7
        Page 8
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        Page 28
        Page 29
        Page 30
        Page 31
        Page 32
        Page 33
        Page 34
    Experimental procedures
        Page 35
        Page 36
        Page 37
        Page 38
        Page 39
        Page 40
        Page 41
        Page 42
    Experimental results
        Page 43
        Page 44
        Page 45
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        Page 119
        Page 120
    Conclusions and observations
        Page 121
        Page 122
        Page 123
        Page 124
        Page 125
        Page 126
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        Page 143
    Biographical sketch
        Page 144
        Page 145
        Page 146
    Back Cover
        Page 147
Full Text







S1 ilI II II I II III i lH 1 I8 11I
3 1262 08552 4618

To Mom and Dad


Sincere appreciation is due many people in this department for

their help during my entire stay at the University of Florida. Most

sincere thanks are due Professor Robert E. Reed-Hill whose continued

guidance and encouragement made this dissertation possible.

Many thanks to Professor F.N. Rhines who first encouraged me and

gave me confidence to strive for a higher degree in metallurgy and

whose continued interest in my program is appreciated.

I would also like to thank the members of my committee,

Drs. Martin A. Eisenberg, Craig S. Hartley and John J. Hren for fruitful

discussions of my work.

Many thanks to my colleagues Messrs. Juan R. Donoso, R.M. Chhatre,

Francisco Boratto and to the laboratory assistants C. Barnes and

M. Brimanson who spent many hours of discussion and who cooperated in

the collection and interpretation of experimental data. The preparation

of the final manuscript by Elizabeth Seville is also greatly appreciated.

The financial support of the Army Research Office (Durham), the

International Nickel Company, and the Energy Research and Development

Administration is greatly appreciated.

Finally, I thank my wife, Kathie, whose patience, encouragement

and understanding during my course of study helped make it all possible.



ACKNOWLEDGMENTS ................................................... iii

LIST OF TABLES .................................................... vii

LIST OF FIGURES ............... .............................. viii

ABSTRACT .............................................. .......... xi

INTRODUCTION ...................................................... 1


I PREVIOUS INVESTIGATIONS ................................... 4

1.1 Static Strain Aging .................................. 4

1.1.1 Historical Aspects ............................ 4

1.1.2 Mechanisms of Static Strain Aging in Metals
Alloys ........................................ 5

1.1.3 Summary of Important Mechanisms of Dislocation
Locking During Aging .......................... 10

1.1.4 Aspects of the Static Strain Aging Experiment.. 11

1.1.5 Static Strain Aging Stages in BCC Metals ...... 13

1.1.6 Static Strain Aging in Nickel and FCC Alloys... 17

1.2 Dynamic Strain Aging ................................. 19

1.3 Work Hardening in Metals and Alloys .................. 23

1.4 Anelastic Phenomena in Nickel and FCC Ferrous Alloys.. 27

II EXPERIMENTAL PROCEDURES ................................... 35

2.1 Materials ............................................ 35

2.2 Experimental Techniques .............................. 37


2.2.1 Swaging ................... .................... 37

2.2.2 Annealing .................................... 37

2.2.3 Specimen Profile Measurements ................ 37

2.2.4 Tensile Testing .............................. 38

2.2.5 Static Aging Experiments ...................... 38

III EXPERIMENTAL RESULTS .................... .................. 43

3.1 The Behavior of the Lower Yield Stress Increase, Ao... 43

3.2 The LUders Extension, EL ............................. 50

3.3 The Hardening Component, AoH ....................... 50

3.4 Activation Energies .................................. 50

3.5 The Dependence of Ao and EL on Prestrain ............ 53

3.6 Comparison of Nickel 270 and Nickel 200 Static
Strain Aging ....................................... 55

3.7 The Stress-Strain Behaviors .......................... 57

3.8 The Work Hardening Behaviors of Nickel 270 and of
Nickel 200 ........................................... 69

IV DISCUSSION ..................... .......................... 80

4.1 Rationale for Static Strain Aging in Nickel 200 ...... 80

4.2 The Mechanism for Static Strain Aging Exhibited in
Nickel 200 .......................................... 82

4.2.1 The Distribution of Vacancies, Carbon Atoms
and Dislocations After Plastic Deformation .... 82

4.2.2 Vacancy Trapping by Carbon Atoms .............. 82

4.2.3 The Concentration of Carbon-Vacancy Pairs ..... 83

4.2.4 Theory of Initial Schoeck Locking by Carbon-
Vacancy Pairs ..................... ........ 85

4.2.5 The Mechanism Controlling the Increase in Ao
with Time ....................... ............. 96

4.2.6 Regarding the Behavior of Nickel 200 After
the Peak in Ao ...................... ...... 107

4.3 Summary .............................................. 'l

4.4 Comments on the Relationship Between Static Strain
Aging and Dynamic Strain Aging in Nickel 200 ......... 115

CONCLUSIONS AND OBSERVATIONS ...................................... 121


FCC METALS ...................................... ... ..... 125

ATMOSPHERE ............................................... 132

BIBLIOGRAPHY ............................ ......................... 135

BIOGRAPHICAL SKETCH ............................................... 144


Table Page

1 Recognized Aspects of Strain Aging ..................... 2

2 The Diffusivity of Carbon in Nickel .................... 21

3 Estimates of the Rate of Vacancy Production During
Plastic Deformation .................................... 32

4 Alloy Compositions ..................................... 36

5 Least Squares Parameters for Static Strain Aging Data
Assuming Ao Is a Function of In t ...................... 45

6 Least Squares Parameters for Static Strain Aging Data
Assuming a Log Ao-Log t Linear Relationship ........... 49

7 The Slopes and Intercepts (at t = Is) of Ao Versus In t
Curves Calculated from Eq. 56 ......................... 108

8 Interaction Energies for Tetragonal Defects in the
FCC Structure ........................................ 128


Figure Page

1 General aspects of the classical static strain
aging test .................... ...... ............... 12

2 An example of the stages of the yield return in Nb-O
alloys [8]; (a) the increase in Ao with time; (b) the
components of AM and their dependence on aging time..... 16

3 Schematic example of stage behavior in polycrystalline
fcc metals ............................................. 25

4 Schematic description of the work hardening behavior in
a metal using a log 9-log a plot ....................... 25

5 Illustrating the method used to determine the aging
parameters EL, Aa= Ly o and AcH = oExt oo. Dashed
loading line indicates the approximate loading line
which would have been observed in the absence of
misalignment of the test specimen .................... 41

6 Selected load-time curves obtained after restraining
a series of specimens 5% at 273"K, aging at 4080K for
the times indicated, and restraining at 2730K ......... 42

7 The time and temperature dependence of the return of
the lower yield stress in Nickel 200. Specimens were
prestrained to a stress level of 265 MPa. The dashed
curves are approximate corrected curves which account
for specimen heat-up in the aging baths ............... 44

8 Normalized aging curves for Nickel 200. The 3730K
curve was normalized to an assumed maximum of 28.5 MPa.
The dashed curves reflect approximate corrections for
the heat-up time of the specimens .................... 47

9 Illustrating the approximate t1/7 power law relation
governing the aging of Nickel 200 at temperatures
below 4480K. Data for the 448K (shown) and 473K
cases do not fit this relation well ................... 48

10 The dependence of the Luders strain on time and
temperature in Nickel 200 ............................ 51

Figure Page

11 The approximate behavior of the secondary hardening
component of the lower yield stress increase
(AcH = "Ext "o) ................. ...................... 52
12 The dependence of Ao and EL on prestrain. Nickel 200
specimens were prestrained at 2730K, aged for
6000 seconds at 4480K, and restrained at 2730K ........... 54

13 (a) The yield return of a Nickel 270 specimen aged for
a time to achieve a maximum in Aofor Nickel 200.
(b) Yield return for a Nickel 200 specimen aged only
one-half as long ......................................... 56

14 True stress-true plastic strain curves for Nickel 270
(- = 4.2 x 10-4 s-l) .................................... 58

15 Variation of the 0.2% yield stress and the ultimate
tensile strength with temperature of Nickel 270
(6 = 4.2 x 10-4 s-1) ............... ..................... 59

16 Variation of the uniform and total elongation with
temperature in Nickel 270 (- = 4.2 x 10-4 s-1) .......... 60

17 True stress-true plastic strain curves for Nickel 200
( = 4.2 x 10-4 s-1) ................. ... ............ 62

18 The temperature dependence of the 0.2% yield stress and
ultimate tensile strength of Nickel 200 ................. 63

19 The temperature and strain rate dependence of the 0.2%
yield stress in Nickel 200 on an expanded stress axis .... 64

20 Variation of the stresses at 5, 11, 19 and 30% plastic
strain with temperature in Nickel 200 .................. 66

21 Variation of the uniform and total elongations with
temperature in Nickel 200. Also shown are the
approximate temperature ranges over which serrations
were observed at the respective strain rates.............. 67

22 Variation of reduction in area with temperature for
Nickel 200 and Nickel 270 ................... ............. 68

23 The log O-log o curves of Nickel 270 (f = 4.2 x 10-4 s ). 70

24 The log e-log a curves of Nickel 200 (C = 4.2 x 10-4 s-1) 71

25 The variation of mil and mIII with temperature in
Nickel 270 (e = 4.2 x 10-4 -1) ....................... 73

Figure Page

26 The variation of mil and mIII with temperature in
Nickel 200 (E = 4.2 x 10-4 s-1) ......................... 74

27 The variation of the work hardening parameter (05%-c0.5%)
with temperature (C = 4.2 x 104 s )................. 76

28 The variation of e2 and e3 (the approximate strains at
which Stages II and III, respectively, begin) with
temperature for Nickel 270 and Nickel 200 deformed at
-l -
a strain rate of 4.2 x 10- s ......................... 78

29 The dependence of E3 on temperature and strain rate in
Nickel 200 .............................................. 79

30 A schematic illustration of the assumed configuration
of the carbon-vacancy pair and the three possible
independent orientations that it may assume ............. 87

31 The concentrations of dipoles in each of the three
possible orientations (B' = 0.3 eV (6900 kcal/mole),
A = 0.2 eV (4600 kcal/mole)) ............................ 91

32 Schematic illustration of the growth of a saturated
carbon-vacancy atmosphere. Rs is time dependent and the
concentration within R is assumed to be a fraction, f,
of the carbon concentration ............. ............ 99

33 The aging curves obtained from the model for strain
aging in Nickel 200 (Eq. 56); the dashed lines are the
experimental data (Figure 7) ..... ..................... 109

34 This diagram illustrates the aging stages of Nickel 200.
The solid line represents the experimental scope of
the present investigation .............................. 116

35 The interaction potential u. of a carbon-vacancy dipole
with a screw dislocation for r = b and A/b = 0.2 eV ..... 131

Abstract of Dissertation Presented to the
Graduate Council of the University of Florida in
Partial Fulfillment of the Requirements for
the Degree of Doctor of Philosophy



Walter Raymond Cribb

August, 1975

Chairman: Robert E. Reed-Hill
Major Department: Materials Science and Engineering

Dynamic and static strain aging were observed in commercially

available Nickel 200 which contains principally 1000 ppm carbon as an

alloying impurity. Static strain aging tests were conducted on

annealed tensile specimens which were prestrained at 2730K to a stress

level of 265 MPa (approximately 0.05 strain) at a nominal strain rate

of 4.2 x 103 s-1. Under these conditions, homogeneous plastic flow

was guaranteed to occur. Specimens were aged immediately after pre-

straining for different times at 373, 408, 428, 448 or 4730K and the

time dependence of the return of the lower yield stress was observed.

The return of the yield experiments indicated that AB increased as

In t or approximately as t/7 kinetically and behaved in accordance with

an activation energy of 25 kcal/mole before the observed peak in Ao.

It is demonstrated that the defect responsible for this anomalous increase

in Ac may be the rotation of carbon-vacancy pairs in the strain fields

of dislocations. A quantitative model is derived for the increase in

Ao before the aging peak and it is concluded that several important

stages in the aging of Nickel 200 may occur: (a) the formation of

carbon-vacancy pairs and their initial ordering, (b) the migration

of vacancies in the strain energy gradients of dislocations and the

consequent formation of more carbon-vacancy pairs near dislocations,

(c) the growth of an ordered carbon-vacancy dipole atmosphere,

(d) depletion of free vacancies in the remainder of the lattice

which decreases the flux to the ordered atmosphere and results in a

Ac maximum, (e) the migration of bound vacancies to dislocation sinks

and the resulting decrease in Ac, (f) the migration of carbon atoms in

the strain fields of dislocations and the growth of a Cottrell atmosphere,

and (g) precipitation of graphite during averaging. Items (f) and (g)

are only speculated to occur. This model is different from the Cottrell-

Bilby model and can account for the kinetics and activation energy for

strain aging observed in Nickel 200.

Tensile tests were conducted between 77 and 800oK at nominal strain

rates of 4.2 x 10-5, 10-4, 10-3, and 10-2 s-1. The results of these

experiments confirm that dynamic strain aging (DSA) in Nickel 200 is

exhibited over a temperature interval between 273 and 5750K at

4.2 x 10- s Over the DSA interval, the following phenomena were

exhibited and depended upon the strain rate: the Portevin-Le Chatelier

Effect, yield stress plateaus, ultimate stress peaks, reduction in area

minima and mild ductility minima. An analysis of work hardening indicates

that anomalous work hardening over the DSA interval is very weakly

exhibited. The mechanism for discontinuous yielding is rationalized

to be dynamic Snoek ordering of carbon-vacancy pairs during plastic

deformation and can account for the anomalously low temperature interval

(with respect to the expected mobility of carbon) over which DSA is

observed to occur.


Currently, eight aspects of strain aging are recognized [1] as

playing a major role in the deformation of polycrystalline metals

(Table 1). The first two are static strain aging phenomena which

are obtained by restraining a set of prestrained specimens that have

been aged at an elevated temperature. The last six aspects listed in

Table 1 are characteristic of dynamic strain aging, i.e., aging which

occurs during plastic deformation. Dynamic strain aging can occur in

both substitutional and interstitial alloys. The most interesting

cases of dynamic strain aging have normally involved interstitial

solutes in transition metals.

Most research on the role of interstitial impurities in the

mechanical behavior of transition metals has been conducted using the

body-centered cubic class of metals such as Fe, Nb, Mo, Ta, 14 and V [2].

The principal interstitial impurities in these metals which are responsible

for strain aging are N, 0, C and H.

Nickel is the only metal of the commercially important Period IV

transition series of the Periodic Table that is face-centered cubic. It

is also the only fcc transition metal widely used for constructional

purposes. The other fcc transition metals Rh, Pd, Ir, and Pt are less

abundant and have not been used as major construction materials. As such,

in-depth investigations of their mechanical properties have not been




Recognized Aspects of Strain Aging

1. Yield Points

2. Strengthening

3. Discontinuous Yielding

4. Strain Rate Sensitivity Minimum

5. Ductility Minimum

6. Abnormal and Rate Dependent Work Hardening

7. Yield Stress Plateaus

8. Flow Stress Transients on Changes in Strain Rate


The principal purpose of the present investigation was to

characterize the strain aging phenomena of commercially available

Nickel 200. This alloy contains as its principal strengthening agent

solid solution interstitial carbon (0.15 w/o maximum). To the best

knowledge of the author, a complete classical static strain aging

investigation has never been conducted using an interstitial solid

solution fcc alloy.

A prime goal in the present investigation was to develop a quanti-

tative model that could explain the kinetics and energetic of the return

of the lower yield stress. Furthermore, the tensile behavior of Nickel 200

during constant strain rate tests conducted over a wide range of

temperatures and strain rates was investigated in order to better define

the dynamic strain aging phenomena in Nickel 200.



1.1 Static Strain Aging

1.1.1 Historical Aspects

It has been recognized for a long time that the yield phenomenon

in iron and other bcc metals is closely related to the presence of

interstitial impurity atoms such as carbon or nitrogen. Most strain

aging investigations have centered about the iron and steel industry

since the 1930s when the phenomenon in low carbon steel first became

a major commercial nuisance.

The first metallurgical investigation of aging in mild steel was

conducted by Davenport and Bain [3] in 1935, who noted that heterogeneous

flow occurred in both annealed and deformed materials after having been

"aged" by storing before working. Subsequent work by Gensamer and

Low [4] in 1944 related the strain aging and yield point to the presence

of trace amounts of nitrogen and carbon. Since the time of these early

investigations, much interest has continued to center on iron and

other commercially significant body-centered cubic metals such as

vanadium [5], niobium [6], tantalum [7], and molybdenum [8]. At the

present time, very little research effort has been directed to the study

of static strain aging phenomena in face-centered cubic and hexagonal

metals containing interstitial impurities.


1.1.2 Mechanisms of Static Strain Aging in Metals and Alloys

Three main dislocation pinning mechanisms have been postulated

on the basis of experimental evidence in metals and alloys. These are

Cottrell pinning, Suzuki locking and Schoeck locking.

In.all of the previously mentioned investigations of bcc metals,

the most plausible explanation for static strain aging is due to Cottrell

and Bilby [9] who attributed the effect to the diffusion of interstitial

atoms in solution (e.g., carbon, nitrogen, oxygen or hydrogen) to

dislocations. Their concept relates the increase in flow stress and yield

point return after aging to the migration of solute atoms to the

tensile region about an edge dislocation. The effect of this segregation

is to locally lower the strain energy of the system and to consequently

stabilize the dislocation to the point where an increased flow stress

is required to remobilize the dislocation or to generate mobile dislocations.

The Cottrell mechanism is of major importance in causing the return

of the sharp yield point in steel while an increase in the steel's

ability to work harden and a reduction in ductility (in the later stages

of aging) are probably associated with precipitation of carbides and

nitrides [10]. The major contribution of the Cottrell-Bilby work was

to solve the problem related to the diffusion of an interstitial atom

in the stress field of a dislocation. The solution predicts the time-

temperature dependence of the rate of impurity migration as inferred

from internal friction measurements [11]. Cottrell and Bilby derived

the following relationship for n(t), the number of atoms arriving at the

dislocation in the time t per unit length,

n(t) = 3 ()l1/3 no(Tt)2/3 (1)

where no is the average number of solute atoms per unit volume and

the parameter A is the interaction constant which describes the

tendency for a solute atom or center of dilatation to be attracted

by an edge dislocation's hydrostatic stress field, D is the diffusivity

of the solute and k and T have their usual meanings. The principal

characteristics of Cottrell pinning as manifested in the static strain

aging experiment are (a) a t2/3 time dependence of the lower yield

stress return, and (b) an activation energy for the yield return

approximately equal to that for the migration of interstitial solute

atoms. The model assumes long range migration of solute and probably

involves about 103 atoms jumps [12] (or a net rms displacement of 30 to

50b). An empirical result is that the increase in stress necessary

to free a dislocation from its atmosphere as measured by Ao is directly

proportional to the number of atoms, n(t), which have arrived at the

dislocation. Thus, the strain energy decrease associated with long

range impurity migration is directly proportional to n(t). This model

with modifications [13,14] has survived for twenty-five years without

its concepts being significantly altered. Excellent reviews 'of the

Cottrell-Bilby theory are available in many places [15-19].

Suzuki [20] has pointed out that in face-centered cubic metals

containing extended dislocations, a completely different form of

interaction between dislocations and impurity atoms can exist. Since

the stacking fault has a locally different crystal structure from the

matrix, the solid solubility of impurities contained in the matrix can

differ appreciably within the stacking fault and outside. Consequently

a chemical potential exists across the fault, resulting in the

binding of impurity atoms to the stacking fault. Pinning is a result

of the accumulation of solute at the fault. Thus, this type of inter-

action should be characterized by an activation energy due to solute

migration. Unfortunately, while the magnitude of the locking stresses

has been calculated and applied with some success to solid solution

alloys, the kinetics of migration to the faults have not been studied [21].

Nickel has a stacking fault energy [22,23] of approximately 400 dynes/cm.

The equilibrium separation [24] between two partial dislocations is

estimated to be only 3b. The magnitude of the yield point produced by

segregation to stacking faults is related to the fault area. Hence, in

other metals such as Ag-6 w/o Al [25] where faults are estimated to

be 30b wide the effect is more important. Thus, one would not expect

Suzuki locking to be a very important pinning mechanism in nickel [26].

The third type of pinning is sometimes called short range order

locking and was proposed by Schoeck [27] and later expanded upon by

Schoeck and Seeger [28]. Schoeck and Seeger considered a bcc lattice

in which the concentration of interstitials is low enough to keep the

interaction between interstitials small. Snoek originally proposed [29] that

small atom impurities in solid solution occupy the octahedral interstices

at the center of an edge or the center of a face of the unit cell in a

bcc metal. Such sites have tetragonal symmetry since two of the six

solvent atoms surrounding the interstitial site are closer than the other

four. As a result, the octahedral sites may be classified into three

groups depending upon which one of three mutually perpendicular <100>

directions the two nearest neighbor solvent atoms are aligned along.

Thus, the three types of interstitial sites correspond to the three

directions of tetragonality and if no applied stress is acting, the

three kinds of interstitial sites will be occupied by the same fraction

of interstitials; namely, one-third will be in each of the three types

of sites. One may visualize each type of site occupied by an inter-

stitial as a dipole. The principal axis of the tetragonal distortion

gives the orientation of the dipole. If an applied (non-hydrostatic)

stress is acting, the energy of interaction between the stress and the

dipoles will in general depend on the orientation of the dipoles

(i.e., the types of sites occupied by an interstitial atom). As a

consequence, an applied stress will cause a redistribution of dipole

orientations and the population of the sites with lower energy will

increase, whereas the population of the sites with higher energy will

decrease. This process is known as the Snoek effect [29] and it gives

rise to a well established internal friction peak [30,31,32]. The

activation energy associated with stress induced ordering of interstitial

solute in bcc metals is normally that associated with diffusion of the

impurity [33]. Schoeck [27] in 1956 pointed out that a similar redis-

tribution of dipole orientations could be effected by the strain field

of a dislocation. By such a process, the energy of the system is lowered

in a period of time approximately that required for one interstitial

atom jump and, therefore, the dislocation becomes locked. Whereas the

locking due to atmosphere formation (Cottrell) requires diffusion of

interstitials over long distances, the locking due to stress induced

ordering of interstitial dipoles is accomplished merely by atomic

rearrangement between neighboring lattice sites and, therefore, takes

place in times which are orders of magnitude faster.

Schoeck and Seeger [28] examined the process in considerable detail

in 1959. Starting with the interaction energy between the interstitial

solute atoms and the dislocation and assuming the concentration of solute

was small, they showed that the line energy of a dislocation surrounded

by a Snoek ordered atmosphere is decreased by an amount U0 given by

mc A2
o = 3LkT
U= In (2)


c = total concentration of interstitials

A = an interaction constant

L = cut-off radius

kT = thermal energy

They next derived that the extra applied stress, AT, necessary to

pull the dislocation from the ordered atmosphere is given by

U kT
AT 2bA (3)

A more complete derivation of these results is carried out in the

discussion relative to the carbon-vacancy pair (Section 4.2.4) and is

related directly to pinning in Nickel 200.

The possible contribution to the rise in yield stress made by

ordering of solute atoms in the stress fields of dislocations has

generally been ignored probably because it occurs very quickly at the

temperatures that have usually been investigated. Snoek ordering,

however, can explain the rapid static strain aging phenomena observed

by Wilson and Russell [34] in tensile tests on a low carbon steel and

similar observations on a range of materials (for example, Carpenter

on tantalum-oxygen [12] and niobium-oxygen [35]; Owen and Roberts [36]

on martensite; Rose and Glover [37] in stainless steel). Support for
this view comes from an investigation by Nakada and Keh [38] of rapid

strain aging in iron-nitrogen alloy single crystals.

Wilson and Russell [34] verified that the rise in yield stress at

261K in iron specimens (containing 0.039 w/o carbon) prestrained 4%

was 63% complete in 100 seconds and noted that this time is in reasonable

agreement with relaxation times observed in the case of the elastic

after-effect due to the ordering of carbon in iron. Thus, the process

is complete in roughly the jump time of a carbon atom. It must be

noted that parts of their data were taken with a reduced applied load

which was generally between 80 and 90% of the load at the end of

prestrain. Aging while applying a load has been shown to influence the

size of the yield point and depends strongly upon the fraction of the

prestrain load that is used in aging [39,40,41,42,43]. A similar case

occurs in the data or Nakada and Keh [38] who used single crystal

Fe-0.l a/o C and N in their investigation.

Although Quist and Carpenter [35] did not conduct the usual

static strain aging experiments, their investigation of dislocation

pinning in Nb-0 alloys during internal friction measurements is note-

worthy. They conducted their experiments between 273 and 3130K and

attributed damping phenomena to the pinning of dislocations by Snoek

ordering of oxygen interstitial atoms in the strain fields of dislocation

line segments. They observed that pinning was effectively completed

in a period of one oxygen atom jump time.

1.1.3 Sumnaryof Important Mechanisms of Dislocation Locking During Aging

The two most important mechanisms of dislocation locking that may

occur in metals containing dissolved interstitial impurities are related

to (a) the Cottrell-Bilby model and, (b) the Schoeck-Seeger model.

The principal features of the Cottrell-Bilby model are

(1) Solute atoms migrate toward the dislocation over long distances

under the influence of the gradient in elastic interaction energy

between the dislocation and the solute.

(2) t2/3 aging kinetics are predicted by the model and observed


(3) The activation energy for the yield return predicted by the model

and observed experimentally is that for the diffusion of

interstitial solute.

The principal features of the Schoeck-Seeger model are

(1) Interstitial atoms with tetragonal strain fields reorient in

the strain field of a dislocation.

(2) The aging process by Snoek ordering is completed in approximately

the time required for one atom jump.

(3) The activation energy predicted by this model is that for

diffusion of the interstitial solute.

1.1.4 Aspects of the Static Strain Aging Experiment

Figure 1 illustrates the mechanics of the classical static strain

aging test for a specimen deformed in tension (or compression) at a

constant strain rate. The initial prestrain and unloading cycle gives

the specimen a known deformation history and internal state, i.e.,

a higher "fresh" dislocation density than that in the annealed specimen.

If the material is immediately restrained after unloading, the stress-

strain curve returns to the curve which would have been attained had

the specimen not been unloaded. However, by aging under the proper

conditions (e.g., higher temperature and/or longer times) a yield point

occurs and is followed by a period of Luders flow at constant load

Oh 7

UnSoad, age, retest


Figure 1. General aspects of the classical static strain
aging test.

before work hardening resumes. In fact, the process of aging results

in a gradual transition from a smooth reloading curve for very short

aging times to a curve similar to that shown in Figure 1.

The important parameters of the reloading curve are oU, oL and

OExt' the upper yield stress, the lower yield stress and the stress
increment obtained by extrapolating the post-yielding curve, respectively.

For short aging times, OExt is equal to o the value of the stress

before unloading. The extrapolated stress, Ext, is determined by

the intercept of the flow curve, i.e., that portion of the stress-strain

curve where uniform strain hardening is present, with the pre-yield

or "elastic" portion of the reloading curve. The experimental parameter,

Ao = L 'o, is the parameter that is normally associated with strain

aging as it is experimentally the easiest to determine.

Accompanying the return of the yield point is the reappearance of

heterogeneous deformation, i.e., the passage of a LUders band down the

specimen gage length, at the lower yield stress, which is also charac-

teristic of annealed metal. During the initial stages of aging, the

lower yield stress increases with aging time as does the size of the

Liders strain. The rate of change of these properties generally increases

with aging time and temperature. After aging for somewhat longer times,

depending on the metal and its history, the variation of Ao and the

LUders strain with time in bcc metals becomes much slower, in many cases

exhibiting a slight decrease with aging time.

1.1.5 Static Strain Aging Stages in BCC Metals

Five stages of aging during static strain aging have been identified

for bcc metals containing interstitials. The first stage has been

explained on the basis of observations of very rapid returns of yield

points in interstitial iron alloys [34,38,44] and internal friction

experiments in other bcc metals [12,35]. The explanation is that very

rapid pinning may be attributed to stress induced ordering of inter-

stitials in the strain fields of dislocations as previously modeled

by Schoeck and Seeger [28] (Section 1.1.2).

The last four stages have been explained [10,45,46] on the basis

of a Petch equatior of the form

o = o. + 2k d-1/2 (4)

where a is the lower yield stress, a. the lattice friction stress,

k a dislocation locking parameter and 2d is the grain size. This

equation was developed originally by Petch [47,48,49] in order to

provide a method of separating the factors contributing to the lower

yield strength of polycrystalline iron. During the LUders band pro-

pagation, it was believed that unpinned sources release many dislocations

which pile up at the grain boundaries. Thus, a feature of the model

is the grain size and the boundaries are pile-up sites which act as stress

concentrations. The pile-ups are controlled by the grain size and act

in conjunction with the applied stress to unpin nearby dislocations in

a neighboring grain. The friction stress is represented by oi and is

the stress to move an unbound or free dislocation through the lattice.

Rosenfield and Owen [50] formulated the aging phenomena in terms

of an equation of the form

Ao = AoH + 2k d-1/2

where Ao is the gain in the lower yield stress after aging, AoH the

gain in the hardening component of the lower yield stress increase,

as determined by extrapolating the load-time curve after the Luders

strain back to the reloading curve,and k and d have the usual meanings.
Szkopiak [6] performed static strain aging experiments on niobium-

oxygen alloys and separated the two components of the yield stress increase of

Eq. 5 as shown in Figure 2. In Figure 2a, the typical return of the

yield stress experiment on a bcc metal shows that at small aging times

the increase in lower yield stress is very rapid (depending upon

temperature of aging) and approaches a maximum. At longer aging times,

the lower yield stress increment shows a slight decrease. In Figure 2b,

the two components of A in Eq. 5 are shown separately.

The five stages of aging that have been observed in alloy systems

such as Nb-0 [6,35], Fe-C [34,45,46], and Fe-N [38,51,52] alloys and

probably occur in Ta [12,18], V [5], and Mo [8] as well are:

Stage I: This stage is observed clearly only at low temperatures since

locking occurs by stress induced ordering of interstitials in the strain

fields of dislocations and occurs within the time span of approximately

one solute atom jump. The strength of pinning agrees reasonably well

with the Schoeck-Seeger model.

Stage II: In this stage, k reaches a maximum and remains constant. The

lower yield stress reaches a maximum and the LUders strain increases very

rapidly. The rationale for this stage is that the formation of Cottrell

atmospheres takes place during aging and upon reloading dislocations become

unpinned from their atmospheres.

Stage III: Further increases in the lower yield stress are due to an

increase of the AoH parameter. In this stage, the Luders strain remains


3 .300

b //

S 10 100 '1000 10000


b b

0o 0
^ / p9-0-0---0-----.0---0--c
C 2 Oxyoen, ppm

o iN *0600

Fr 10 100 1000 10000
Aging Time, Minutes

Figure 2. An example of the stages of the yield return in
Hb-O alloys [8]; (a) the increase in Ac with time,
(b) the components of 0A and their dependence on time.

nearly constant as the yield stress increases and enhanced strengthening

occurs. The principal rationale for this hardening is that dislocations

have been aged to the extent that they tend to remain immobile or pinned

upon subsequent reloading. Thus, new or additional dislocations

are created and the yield stress continues to increase.

Stage IV: During this stage, solute continues to be accomodated in the

strain field of dislocations but no longer effects an increase in Ao.

Hence, Ac remains approximately constant.

Stage V: As more and more interstitial solute segragates to dislocations,

a condition of averagingg" is satisfied and precipitates may form;

hence, the loss of a coherent strain field or the robbing of solute near

dislocations and a mild decrease in the hardening component.

The above stages of static strain aging appear to hold true for

most of the body-centered cubic metals containing interstitial oxygen,

carbon or nitrogen. However, in the case of face-centered cubic metals

containing interstitial impurities no complete investigations of the

behavior of the return of the lower yield stress have been carried out.

1.1.6 Static Strain Aging in Nickel and FCC Alloys

Among the fcc commercial alloys, nickel containing carbon is

probably the most significant where an interstitial (carbon) is

deliberately added to improve mechanical properties. Other than pure

nickel, only fcc multicomponent alloys such as the austenitic stainless

steels [37,53,54] contain carbon for similar reasons and exhibit

mechanical properties similar to nickel-carbon alloys.

There exists some experimental evidence related to the static strain

aging of nickel-carbon alloys. In particular, two short notes were

published by Macherauch et al. [55] and by Macherauch and V6hringer [56]

regarding static strain aging in Ni-0.05 w/o C after restraining

slightly beyond the initial yield plateau. Their data were plotted

by this author and an approximate t/3 or t1/4 time dependence of the

lower yield stress return was exhibited. They did not speculate on

the kinetics; however, they determined an activation energy of

10.22 kcal/mole in agreement with the activation energy for diffusion

of hydrogen in nickel (see, for example, Boniszewski and Smith [57]).

The only other investigation relating to nickel-carbon alloy is

due to Sukhovarov [58] and Sukhovarov et al. [59]. Using compression

specimens deformed at room temperature and aged between 433 and 493K

for various times, they deduced with apparent difficulty (because

serrated flow occurred) that the average activation energy was

30.7 kcal/mole, somewhat lower than the carbon migration energy in

nickel. This author plotted their lower yield stress data and noted

that 6o/AOmax varied approximately as t0.3. They conclude (incorrectly,

it is believed) that the Cottrell-Bilby model explains the rise in

Ao and that probably the formation of precipitates eventually occurs;

they never observed this aspect.

In addition, since serrated flow occurred in the investigation of

Sukhovarov [58], it is probable that the data were scattered because

the lower yield stress was not as clearly defined as in the present

investigation where yield point measurements were made under conditions

precluding serrated flow. On this basis, their data should be used


Hydrogenated nickel exhibits strain aging behavior [57,60-65] when

deformed below room temperature. Much effort has been directed toward

understanding fracture, ductility, and other embrittlement related

phenomena attributable to hydrogen. Also, serrated flow [57, 60-65] is

exhibited in hydrogenated nickel between approximately 130 and 2250K
-4 -1
[65] at a nominal strain rate of 10- s The kinetics of static

strain aging in hydrogenated nickel were very briefly investigated by

Boniszewski and Smith [57] and they concluded that the Cottrell-Bilby

model can account for static strain aging of charged specimens. However,

they did not speculate on the exact kinetics that the experimental

data may have followed.

Marek and Hochman [66] have demonstrated the existence of static

strain aging effects in AISI 316 alloy and related it to the approximate

activation energy for diffusion of interstitial carbon in austenite.

However, the effect was most marked in the micro-yield region (0.01% proof

stress) with no effect on flow stress after yielding, UTS, or elongation,

which is indicative of a low interstitial/dislocation interaction energy.

1.2 Dynamic Strain Aging

Dynamic strain aging (DSA) is a feature exhibited in most commercial

metals and alloys [1]. In general, aspects of DSA have created very

little interest in the past, probably since in steel it exhibits its

most significant effects over a temperature range around 4500K where

steel is not normally worked. Other bcc metals such as titanium,

tantalum, niobium and vanadium exhibit DSA over a temperature range

where these metals are most needed [1].

As with steel, nickel containing carbon exhibits its effects at

relatively low temperatures (300 to 5000K, roughly). Early investigations

by Sukhovarov and Kharlova [67] confirmed that dynamic strain aging

occurs in nickel when alloyed with small amounts of carbon. In a

subsequent investigation Popov and Sukhovarov [68] indicated that the

apparent activation energies associated with the appearance and

disappearance of serrated flow are 20 and 33 kcal/mole, respectively.

No conclusion regarding the very low activation energy for the onset of

serrated flow was ventured. On the other hand, Nakada and Key [51]

have indicated that the onset activation energy in Ni-C alloys is

152 kcal/mole and that for the disappearance of serrations is 264

kcal/mole. It is a well established experimental fact that the acti-

vation energy associated with the diffusion of carbon in nickel is

approximately 35 kcal/mole (see Table 2). Generally, in interstitial

alloys the activation energy for the onset of serrated flow is

associated with the diffusion of impurity atoms and is made on the

basis that when the velocity of dislocations is approximately equal to

that of the velocity of the diffusing impurity atoms, a drag or pinning

of dislocations occurs. Thus, the pinning as observed through the

serrated flow phenomenon is assumed to be controlled by the diffusion

of impurity atoms just as in the static strain aging experiment. By

plotting log E versus 1/T, where c is the strain rate at temperature T

where serrated flow is first observed, an activation energy may be

deduced. The values of 152 and 20 kcal/mole for the onset of serrated

flow determined by the preceding authors are much too small to be related

to the diffusion of carbon in nickel. Popov and Sukhovarov [68] made

no conclusions regarding this apparent anomaly. Nakada and Keh,

however, ventured that pipe diffusion of carbon along dislocation cores

controls serrated flow in nickel.



The Diffusivity of Carbon in Nickel

Do Q Technique
cm2/sec kcal/mole

0.048 34.8 Elastic and magnetic aftereffect [69]

0.13 34.5 Radioactive tracer [70]

0.1 33.0 Radioactive tracer [71]

-- 38.5 Thermogravimetric [72]

-- 39.7 Thenogravimetric [72]

-- 32.3 Magnetic aftereffect [73]

Regarding the activation energy associated with the disappearance

of serrations, Kinoshita et al. [74] proposed that this value may

represent the sum of the activation energies for the diffusion of solute

plus the binding energy of solute atoms to dislocations. On this basis,

Nakada and Keh [51] have deduced a binding energy of 11.0 kcal/mole

(0.5 eV) for a carbon atom to a dislocation in nickel assuming that

serrated flow in Ni-C alloys is caused by carbon directly. Popov and

Sukhovarov [68] attributed their value of 333 kcal/mole (1.4 eV) for

the disappearance of serrations to a combination of creep processes

coupled with Cottrell atmosphere formation.

Other than the above nickel-carbon studies and the experimental

work on hydrogenated nickel [57,60-65] which shows strain aging, no other

investigations of the effect of interstitials on the stress-strain and

work hardening behaviors in pure face-centered cubic metals have been

conducted. However, some face-centered cubic ferrous alloys appear to

possess mechanical properties similar to those of nickel-carbon alloys.

In an investigation by Jenkins and Smith [54] complications due to

substitutional alloying elements such as Cr occurred. Nevertheless,

AISI 330 stainless steel (Fe-15Cr-33Ni-0.4C), exhibits similar dynamic

strain aging trends. A calculation of the energies for the onset and

disappearance of serrations revealed that 26.6 and 62.0 kcal/mole are the

onset and termination activation energies. They indicated that the

onset activation energy is very close to that for vacancy migration in

Fe-30 Ni. As the Portevin-Le Chatelier effect is absent for low carbon

content, they conclude that vacancies alone are not responsible and that

carbon-vacancy pairs account for the observed activation energy. Mention

was not made of the exact mechanism for the pinning during serrated flow.

A similar argument appears to apply in the case of Nickel 200.

Other strain aging effects occur in face-centered cubic alloys

but these arise mainly from the diffusion of substitutional solutes

and are outside the scope of this dissertation.

1.3 Work Hardening in Metals and Alloys

In studying the work hardening of metals and alloys it is desirable

to determine the mechanisms that control the rise in flow stress.

In general, this involves relating the macroscopic behavior to changes

in the microscopic structural features of the metal. For example,

observations of slip line lengths or dislocation structures can

supplement an explanation of work hardening.

From the macroscopic point of view, polycrystalline stress-strain

curves have been shown to be generally piece-wise continuous [75-81].

For example, Figure 3 shows schematically that a polycrystalline face-

centered cubic metal may deform so as to show discontinuities in its

stress-strain behavior. Zankl [75] and others [76-81] have shown that

these stages can be correlated very well with deformation processes.

According to his experimental work [75] the identifiable stages are

related to the following processes:

1. The Transition Stage. This extends from zero plastic strain

to approximately 0.1%. In this stage, multiple slip starts first

in the largest grains and then spreads into neighboring grains.

2. Stage I. This begins when all grains have begun to deform

with slip still involving multiple slip systems. Stage I in

polycrystals is thus basically different from the easy glide

Stage I of face-centered cubic crystals. It ends at approximately

E = 1.0%.

3. Stage II. Here slip tends to occur predominantly on a single

(primary)system but with interaction from secondary systems. The

deformation is accordingly analogous to that in Stage II of a fcc

single crystal. Large grains may break down into several regions [80]

with different primary systems. This stage extends to about

e p5.0% in pure fcc metals such as copper and nickel.

4. Stage III. As in fcc single crystal deformation, this stage

is controlled largely by dynamic recovery and has been associated

with cross slip [82].

In general, polycrystalline stress-strain curves appear to be

continuous in shape and the stages difficult to identify on such curves.

This is in marked contrast to single crystal stress-strain curves which

often exhibit well-defined stages. A sensitive empirical method [83]

for detecting polycrystalline stage behavior has been developed as a

logical projection of previous empirical analyses [84-92]. This is

based upon the assumption that each stage of stress-strain behavior can

be reasonably described by a modified Swift [92] equation:

e = Eo+com (6)

where a is the true stress, E the true plastic strain, m the work

hardening exponent and o. and c are constants.

One may solve very simply for the parameters in Eq. 6 by plotting

log 0 versus log o where 0 = a A schematic example of such a plot

is shown in Figure 4 for a typical face-centered cubic metal such as

copper or nickel. Any straight line on this type of plot has an

equation of the type:

log 0 = (1-m) log o log cm (7)



Figure 3. Schematic example of stage
fcc metals.

behavior in polycrystalline


LOG 0-

Figure 4. Schematic description of the work hardening behavior
in a metal using a log C-log a plot.

Thus, the value of m typically characterizes the power law

relationship of Eq. 6. An m equal to one is a linear stress-strain

curve and a log 0 versus log o plot would show a line with a zero

slope; a parabolic stress-strain curve would show a (l-m)-value of

-1.0 (i.e., m=2) and so on. High m values correspond to curves with a

great deal of curvature, i.e., very rapidly decreasing work hardening

rate with increasing stress as in Stage III when dynamic recovery processes

reduce the work hardening rate very rapidly with continued deformation.

All other parameters held constant, a high value of m for a single

stress-strain curve would, in general, imply that the material has low

ductility, even though the material might possess a reasonably high

ultimate strength. However, it should be noted that a three or four

stage stress-strain behavior could well lead to a combination of both high

strength and high ductility depending upon the m values of the various

stages and the extent of a particular stage during deformation.

It should be noted that the stage behavior observed by using Eqs. 6

or 7 is within limits independent of the empirical power law equation

used. For example, an analysis based upon the Crussard and Jaoul

method [87-89] using a log a log E diagram shows that discontinuities

occur at the same places on the stress-strain curve as determined by

using a log 0 log a diagram.

In the current investigation, interstitial solute concentration

was the principal alloy variable known to affect stress-strain behavior.

Interstitial elements tend to significantly increase the strength of a

metal while generally decreasing its ductility. This in turn affects

the stage behavior of the parent metal.

Another factor influencing the stage behavior of metals and

alloys is the stacking fault energy. This intrinsic property of a

metal or alloy determines the separation distance of the two partial

dislocations of an extended dislocation [93].

For high SFE metals like nickel [22,23,94] or aluminum [93] the

separation distance is small and mobile dislocations may be assumed

to approach total dislocations. In low SFE metals and alloys such as

Ni-40Co and Ni-60Co [94] this distance becomes appreciable and the

partial may be separated by a wide band of stacking fault. By

effectively varying the geometry of the dislocation by lowering its

stacking fault energy, the deformation behavior might be expected to

change as well. Thus, it would appear that in a low stacking fault

energy metal a dislocation effectively loses a degree of freedom of

movement by virtue of its assuming a planar character. This should

reduce the ability of the material to undergo dynamic recovery involving

either cross-slip or climb.

One may also view a stacking fault as a building block for a

deformation twin. In general, twin boundary energies are lower in lower

stacking fault energy metals. This observation is in agreement with the

fact that twinning plays a greater role in the deformation behavior of

metals and alloys [95,96] with low stacking fault energies. Twinning

was not observed in the present investigation.

1.4 Anelastic Phenomena in Nickel and FCC Ferrous Alloys

Considerable research has been performed for many years on the

effects of interstitial solutes on the strain aging of body-centered

cubic metals and alloys [2]. In addition, much emphasis has shifted

toward examining a variety of internal friction effects (such as the'

Snoek effect, the cold-work peak and dislocation damping behavior)

that can be correlated closely to strain aging phenomena [97].

In contrast, however, there exists but a dearth of strain aging

and internal friction studies of interstitial solid solutions of face-

centered cubic metals. The reason for this lack of interest probably

stems from the fact that in terms of the mechanical behavior of these

alloys, the interstitials apparently cause a less dramatic effect on

the mechanical properties. In addition, fcc metals are not expected to

exhibit a Snoek peak because of site symmetry [30,32,98,99].

Very different specific mechanisms for the observed relaxation

peaks in fcc metals and alloys have been suggested by various authors.

Adler and Radeloff [99] reviewed the types of defects which could

account for internal friction in fcc metals and alloys:

(1) interstitial-solute clusters which have noncubic strain

fields [30,98];

(2) interstitial-substitutional solute clusters in which the

interstitial reorients preferentially under stress if it is a

near neighbor to an immobile substitutional solute [98,99];

(3) interstitial-vacancy complexes of different types, i.e., Wu

and Wang [100] suggested a defect consisting of one interstitial

occupying a vacancy with a nearest neighbor interstitial in its

normal site.

Within each category listed above are many possible specific

combinations which could in principle cause a relaxation effect.

There is evidence for relaxation due to carbon pairs in both Ni

and the fcc allotrope of Co, and for oxygen pairs in Ag [101]. The

existence of an internal friction peak associated with dissolved

C in Ni was first reported by KB and Tsien [102] and further inves-

tigated by KQ, Tsien and Misek [103]. The peak is quite small and

occurs at 5230K for a frequency of 1 Hz. By application of a saturating

magnetic field, the existence of the peak was shown to be unrelated to

the ferromagnetic nature of the sample. Further, the peak was found to

decline in strength as the carbon precipitated from solution. KB and

Tsien pointed out that unassociated C atoms, located at the body-centered

position of the fcc structure and in the equivalent positions midway

along the cube edges, could not be responsible for the relaxation, since

the symmetry of such defects is cubic. After Tsien [104] found in later

work that the relaxation strength varied essentially as the square of

the carbon content in solution, it became evident from mass-action

considerations that carbon pairs were the responsible defect.

Diamond and Uert [69] also investigated the diffusion of carbon

in nickel utilizing elastic aftereffect measurements above and below

the Curie temperature and noted no discontinuity in carbon diffusivity.

Hence, magnetic transformation effects do not affect carbon diffusivity.

In addition, they concluded that the elastic aftereffect is due

principally to interstitial diffusion of C-C pairs in agreement with

the results of the previously mentioned group of Chinese investigators

[102-105]. The simplest interstitial pair configuration in their view

consists of two atoms occupying nearest neighbor octahedral sites which

are the largest for fcc metals, e.g., the two sites 00 and 00. The

stress induced reorientation of such pairs (designated 110 pairs) which

gives rise to anelastic effects, is the result of one of the atoms

jumping into an unoccupied nearest neighbor interstitial site. A summary

of Ni-C diffusion data is presented in Table 2. It should be noted

that all methods agree reasonably well with an activation energy of

35 kcal/mole for the migration energy of carbon in nickel. No

distinction is made in the last four references cited in the table

concerning the nature of the mobile species, e.g., dicarbon complexes.

C.T. Tsien [104] considered the effect of impurities on an internal

friction peak in a carburized 18.5 w/o manganese steel. The principal

experimental observation was that the internal friction peak height

varies linearly with carbon content. It was proposed that the addition

of Mn to Fe-C alloys may reduce the opportunity of forming carbon-

carbon atom pairs and there may be greater probability of forming

Mn-C pairs instead. Thus, in high manganese steels the internal

friction peak is not attributed to rotation of the carbon atom pairs but

due to that of carbon-manganese pairs. As a result, the height of the

internal friction peak was observed to be directly proportional to the

carbon content.

This brings us to the point of the possible mechanism involved in

Nickel 200 which has 0.18 w/o Mn and smaller amounts of iron and copper.

Presuming a pair mechanism (in order to obtain a tetragonal defect which

would account for an internal friction peak) one might presume a possi-

bility of C-C atom pairs or C-Mn atom pairs causing the static strain

aging effects. However, in the investigations previously sited, not one

was conducted using a deformed metal and this factor may be an important

consideration. Hence, a third and possibly fourth speciesmay be involved,

namely, the vacancy and the self-interstitial. It has been demonstrated

by Seitz [106] that in fcc metals the predominant defect produced

during plastic deformation is the vacancy. Table 3 shows the data of

a number of authors [107] who have described the vacancy concentration

during plastic deformation for a variety of fcc metals and alloys as

Cy = Ken (8)

where cv = vacancy concentration (atom fraction)

e = true strain

K = proportionality constant


The principal experimental technique used to determine c involves

monitoring resistivity changes during deformation and subsequent

annealing of the specimens [106,108]. The identification of the

defects annealing out during each recovery stage has been the subject

of extensive previous work in nickel. Studies have been made of the

resistivity recovery spectrum following neutron irradiation, electron

irradiation and quenching from high temperatures [109,110]. In addition,

changes in magnetic properties have also been studied in conjunction

with the resistivity recovery process [73,111,112].

On this basis [108] cold rolled nickel behaves as

c = 2.1 x 10-'4 (9)

In addition, it has been estimated that approximately eight times as

many vacancies as self-interstitials are generated.

Point defects are generally believed to be generated during plastic

deformation by two basic mechanisms [31,113]: (a) nonconservative

motion of jogs on screw (or mixed character) dislocations and

(b) recombination of dislocations containing edge type components.

The second mechanism is due to annihilation of edge dislocations and may


Estimates of the Rate of Vacancy Production
During Plastic Deformation

Material K Type of Deformation

Cu 1.9 x 10-4 tensile elongation

Ni 2.1 x 10-4 cold rolling
Ni 12 x 10- shock forming

NaC1 1 x 10- compression

70-30 Brass 5.9 x 10-4 tensile elongation

Au 2.9 x 10-4 tensile elongation

Al 0.2 x 10-4 tensile elongation

* From Ulitchny and Gibala [107]

not be important relative to the first mechanism in the production

of vacancies until Stage III deformation occurs. The relation between

vacancy concentration and strain is also probably only valid to strains

of the order of 10%. It has been noted that the vacancy concentration

tends to approach a saturation value at large strains [108]. Thus,

vacancies may play an important role in deformed and aged nickel.

As concerns the manganese impurity in Nickel 200, a qualitative

argument may be made against the existence of Mn-C complexes as opposed

to C-V complexes. The gram atomic volume of Mn [114] is 7.39 cm while

that for Ni [115] is 6.59 cm3. Thus, a Mn atom is only approximately

4% oversize on a nickel lattice site. In an interstitial octahedral

lattice site, a carbon atom is about 14% oversize. Relative to binding

to a vacancy, the carbon atom should provide a stronger compressional

center of dilatation than the manganese atom even though in the former

case an octahedral interstitial site is occupied and in the latter a

normal lattice site is occupied. Hence, the binding energy of a vacancy

to a carbon atom should be higher than that of a manganese atom to a

carbon atom or a vacancy. In addition, Nickel 200 contains approximately

three times less manganese than carbon on an atom fraction basis.

The value for the diffusion energy of a vacancy through a lattice

is very sensitive to the presence of impurities which tend to slow down

a freely migrating vacancy because of binding to impurity atoms [116,117].

In pure nickel, the vacancy migration energy is between 0.8 and 0.9 eV

and for impure nickel [117,118] (99.9% pure plus an unspecified amount of

carbon) is approximately 1.1 eV. Thus, one may deduce that the approximate

binding energy of carbon to vacancies is between 0.2 and 0.3 eV. In

short, the carbon-vacancy interaction could cause the formation of a

defect complex which might cause strain aging in Nickel 200. It is

not clear whether or not dicarbon-vacancy defects may be completely

ruled out as a possible complex causing strain aging.

An excellent discussion concerning internal friction and strain

aging to carburized ferrous austenite by Ulitchny and Gibala [107]

suggests that the relaxation phenomena in these alloys are attributable

to the rotation of carbon-vacancy pairs. Their conclusion is based

upon experiments in which vacancies were produced in stainless steels

by (1) quenching, (2) deformation and (3) irradiation. These processes

have in common that they (a) increase the vacancy concentration in

the metals, (b) increase the observed peak heights of the bound pair

peak and (c) increase the peak heights in proportion to the relative

numbers of vacancies they are anticipated to produce. These alloys

possess mechanical properties quite similar to nickel-carbon alloys.

In addition, the diffusivities of both carbon atoms and vacancies in

the austenitic stainless steels are similar to those in nickel-carbon

alloys [107] suggesting that a similar mechanism in the present investi-

gation of Nickel 200 should not be ruled out.



2.1 Materials

Nickel 200 bars of 0.75 inch diameter were obtained through a

local supplier and from the International Nickel Company. These two

heats had slightly different compositions (Table 4); the mechanical

properties as a result were somewhat altered, albeit small. All static

strain aging experiments were conducted utilizing Nickel 200b. The

standard tensile tests were conducted using Nickel 200a. This procedure

was followed in order to minimize possible scatter of the data,

particularly in the static strain aging experiments. In addition,

Nickel 270 was also purchased; its composition also appears in Table 4.

Nickel 200 contains approximately 0.18 w/o Mn and 0.10 w/o C.

The highest equilibrium solubility limit [119]for the nickel-carbon

system is 0.27 w/o carbon. The room temperature solubility limit of

carbon in nickel is only 0.02 w/o [120]. Also, nickel carbides are

unstable in Ni-C alloys [119,121]. The development of visible graphite

in nickel during cooling is generally agreed to occur very slowly.

The present specimens, which were furnace cooled from the annealing

temperature (10730K) at a rate of approximately 2.0/s, are believed

to have retained all of the carbon in excess of the equilibrium solubility

in solution since graphite was not observed either by optical or

transmission electron microscopy of the annealed specimens.



Alloy Compositions

Ni 270






Cu <0.001

All other impurities less than 100 ppm each.

Ni 200a







Ni 200b







2.2 Experimental Techniques

2.2.1 Swaging

The 0.75 inch (19.1mm) bar stock of all the materials were cold

swaged in a Model 3F Fenn rotary swaging machine to a diameter of

0.25 inch (6.4mm). Intermittent annealing was not necessary. The

resulting swaged bars were machined into threaded-end specimens with a

nominal reduced section of 0.8 inch (20mm) and a gage diameter of

0.15 inch (3.8mn).

2.2.2 Annealing

Annealing was accomplished in a Vacuum Industries Minivac furnace

assembly. This unit utilizes a resistance heated tantalum element.
Pressures as low as 10 millitorr can be maintained. No cold trap

was used. All Nickel 200 specimens were annealed for 30 minutes at

800C (10730K); Nickel 270 specimens were annealed for 32 minutes at

5950C (868K). These treatments resulted in specimens with a mean grain

intercept of approximately 22pM. Annealing twin boundary intercepts

were not counted in obtaining this result.

2.2.3 Specimen Profile Measurements

A Jones and Lamson Optical Comparator capable of measuring to

0.0001 inch (2.5pM) in the vertical and 0.001 inch (25pM) in the

horizontal directions was used to measure the profile of as-annealed

specimens. For all tests a specimen gage length was assumed to be

identical to its reduced section and was determined to within 0.005 inch

(0.13mm) with experience.

2.2.4 Tensile Testing

All tensile testing was performed on two Instron machines (Model TT-C

and Model FDL of 10,000 and 20,000 pound capacities, respectively). The

standard crosshead speed was 0.02 inch/minute resulting in a nominal

specimen strain rate of 4.2 x 10-4 s Three additional strain rates

were also employed with the Nickel 200 specimens. All tests were conducted

between 77 and 9000K. Above 2970K tests were carried out in a capsule

using commercial purity argon gas. At no time was oxidation visible on the

specimen surfaces. Below ambient temperature, liquid nitrogen (770K),

dry ice-acetone (1960K), or ice-water baths were employed.

Load-time curves were processed to yield true stress versus true

plastic strain curves as well as the slope of these curves as a function

of stress or strain. The stage behavior of the specimens was analyzed [83].

In brief, the procedure involves the plotting of log E versus log o

and identifying portions of the curves through which straight lines may

be passed. Each linear interval is assumed to represent a stage.

This assumption was tested against results of Zankl [75], Schwink and

Vorbrugg [77] and others [76,78-81] and a good correlation was obtained

between stages determined in this manner and the method of Zankl and

others using different plotting and metallographic procedures. To each

linear region a different set of parameters (m,c, o) may be deduced

corresponding to the Swift [92] equation

E'Eo+Cm (10)

2.2.5 Static Aging Experiments

Annealed tensile specimens were prestrained approximately 5% to

a stress level of 38.5 ksi (265 MPa) at 2730K and immediately unloaded,

removed from the testing jig and immersed in a silicon oil bath at

473, 448, 428 and 4080K or in a boiling water (distilled) bath for

times as long as 2 x 106 seconds (approximately two weeks). Control

of the constant temperature baths was held to within approximately

one-half degree. Upon completion of the aging treatment, the specimen

was removed and quickly quenched into cold water and tested immediately.

The reloading temperature for the static strain aging tests was 273K

as in the prestrain. It was recognized early in the investigation that

at room temperature and a strain rate of 4.2 x 104 s- (the rate

corresponding to a crosshead speed of 0.02 inch/minute, standard at

this laboratory), discontinuous flow occurred in Nickel 200. To

alleviate this feature it was decided to conduct the prestrains in

ice-water baths at a rate of 4.2 x 10-3 s This had the additional

benefit of producing a stable and reproducible lower yield stress

plateau and, in addition, the ice-water bath assured that specimens

were at the test temperature after removal from the high temperature

aging baths.

Quite apparent from the start was the fact that specimen alignment

offered a problem. Upon removal and replacement of a specimen in the

tensile jig, it was evident that exact repositioning was difficult and

uncertain. Therefore, on restraining a specimen after aging it, a

small bending moment normally develops which may cause yielding to

occur nonunifonnly across the gage section of the specimen. The result

is that the upper yield point as observed on the machine chart was

usually absent. As a consequence all data reported are lower yield

stresses. Compilation of the data included interpreting the Luders

extension as the chart displacement which occurred at constant load

(Figure 5). This method proved to provide the most consistent set of

data and tended to alleviate apparent alignment or reloading displacements.

Occasionally (perhaps 10% of the time), a yield point was observed and

the data fell consistently in line with other LUders extension data

recorded by the former method.

The latent hardening achieved during long term and high temperature

aging treatments was computed by linearly extrapolating the post-

yielding curve back to the reloading line as demonstrated in Figure 5.

This method also proved consistent. However, in most cases of short

aging times or low temperature aging, the extrapolated stress fell

below the prestrain value. This is physically unreasonable and is

attributed to alignment effects which generally eliminate the yield

point as mentioned previously and may cause the extrapolations to come

back to the reloading curve somewhat low. However, the resulting data,

again, proved to be consistent.

As an example of typical experimental results, Figure 6 shows a

series of load-time curves which were obtained after restraining a

series of specimens at 2730K to 5%, aging at various times at 408K,

and restraining at 2730K. Figures 6b and 6d show examples of partial

yield points that were obtained in a few cases during the present inves-


Aged 2.92 x 104s/4480


Ext.- -157 N

O 3s

Figure 5. Illustrating the method used to determine the aging
parameters EL, AC= oLY o and AOH = OExt o .
The dashed loading line indicates the approximate
loading line which would have been observed in the
absence of misalignment of the test specimen.

^ --157 N


/ s/ 960 s

157 N


1.54 x 104s 1.74 x I's

157 N


5.78 x 10 s


Figure 6. Selected load-time curves obtained after restraining
a series of specimens 5% at 2730K, aging at 4080K
for the times indicated, and restraining at 2730K.



3.1 The Behavior of the Lower Yield Stress Increase, At

Figure 7 illustrates the dependence of the lower yield stress

increment, Ao, on time and temperature over five decades of time and

at five temperatures. The principal features of Figure 7 are

(1) The curves appear to approach a common value at small times.

It would seem that the data obtained at the higher aging temperatures

and for very short aging times are influenced by the time required to

heat the specimen to temperature. This was confirmed by assuming a

fixed heat-up time and displacing the curve (at each temperature) to

shorter aging times. Thus, a heating time of approximately 40 to 60

seconds straightens out the start of the higher temperature curves to

approximately the same linear dependence as exhibited by the 3730K curve.

In addition, a simple heat transfer calculation indicated that a time

constant of approximately 40 to 60 seconds should describe the specimen

heat-up time. The dashed lines in Figure 7 show the approximate

corrections necessary to account for heating the specimens in the baths.

(2) Each Ao curve shows a roughly linear increase with log t for

times before reaching the maximum in Ac.

(3) All curves show a well-defined peak whose height increases

slightly, the lower the aging temperature.



25 -

20 _/448


I 2 3 4 5 6
10 10 10 10 10 10
.Time (sec)

Figure 7. The time and temperature dependence of the return of
the lower yield stress in Nickel 200. Specimens were
prestrained to a stress level of 265 MPa. The dashed
curves are approximate corrected curves which account
for specimen heat-up in the aging baths.



Least Squares Parameters for Static Strain Aging Data
Assuming Ao Is a Function of In t

t (sec)

480 t 1.43 x 106

240 t 6.62 x 104

60 t 2.16 x 104

30 t 3.4 x 103

not fitted

Slope Intercept r
(MPA/In sec)

2.088 -4.922 0.996

2.647 -3.356 0.998

2.602 1.618 0.995

2.318 6.689 0.998


(4) Ao decreases significantly after max is passed. Just

after the peak, the decrease in Ao is almost linear with log t.

However, at 473"K, the highest temperature investigated, the data at

very long aging times show that the lower yield stress decreases to

a constant value of approximately 17 MPa.

Table 5 lists values obtained by the method of least squares for

the slopes and intercepts (at t=1s) of the aging curves in Figure 7.

A Ao versus In t relationship is assumed to hold for times before

Ac reaches a maximum.

In addition, Figure 8 shows a set of Am curves which were normalized

to their respective maximum peak heights, Ammax. An approximate value

of 28.5 MPa for Aomax was assumed for the 3730K curve since, for the

aging times investigated, a peak was not attained. This diagram has

the effect of making the aging curves more nearly parallel.

Figure 9 is a plot of the increase in lower yield stress versus

t /7 using data from the 373, 408, 428 and 4480K aging curves of

Figure 7. Least squares analysis of log Ao versus log t curves indicated

that Ao varies approximately as the 0.14 and 0.15 power of time. See

Table 6 for the complete results. This represents an approximate time

dependence of t/7

It is important to note that this author prepared a specimen which

had been aged for approximately 6 hours at 5250K (specimen taken from

the undeformed threaded end of a deformed tensile specimen) and observed

it carefully in a transmission electron microscope. No evidence of

precipitates or free graphite in the grains, at dislocations or at grain

boundaries was observed. Thus, it is highly unlikely that precipitation

of carbon occurs during aging between 373 and 4730K in Nickel 200.

Figure 8. Normalized aging curves for Nickel 200. The 3730K
curve was normalized to an assumed maximum of
28.5 MPa. The dashed curves reflect approximate
corrections for the heat-up time of the specimens.

o I 2 3 4 5 6



0- 0


o 4480K
// 4280K
5- 4080K

o I 2 3
t (min7)

Figure 9. Illustrating the approximate t1/7 power law relation
governing the aging of Nickel 200 at temperatures below
4480K. Data for the 4480K (shown) and 473K cases
do not fit this relation well.


Least Squares Parameters for Static Strain Aging Data
Assuming a Log Ao-Log t Linear Relationship

Temperature t (min Slope Intercept r

373 1.0 < 23,752 0.145 0.804 0.993

408 1
428 2 t 710 0.137 1.109 0.997

448a 2t 100 0.099 1.232 0.998

473b 10< t 5 30 0.055 1.307 1.000

a Three points

bTwo points

3.2 The Liders Extension, EL

Figure 10 illustrates the dependence of the LUders extension on

log t. The behavior is similar to that exhibited in Figure 7 for

curves of Ao versus log t. However, note that the peak in the Luders

extension occurs (at a given temperature) at an earlier time than the

corresponding peak in Ao. This may be indicative of the onset of

hardening and is consistent with results obtained in bcc metals. Note

also that at long times the extensions tend to return to their short

time values; that is, they do not remain constant after reaching their

maximum values. An interesting point is that the highest temperature

curve (4730K) reaches an apparent minimum at approximately 1500 minutes.

3.3 The Hardening Component, AoH

Figure 11 is a plot of AoH, the hardening component of the increase

in lower yield stress. This parameter was deduced as noted in Figure 5

by using the equation AoH = OExt oo. Figure 11 shows that this

hardening component appears only at discrete times. Note also that

AOH peaks at approximately the same time as Ao and decreases to a value

higher than that observed at very short aging times. It should be noted

that the values of AoH cannot be taken as exact due to alignment problems

which affected the choice of OExt (see Figure 5).

3.4 Activation Energies

In order to establish a mechanism for static strain aging, the

apparent activation energies associated with particular time dependent

aging events were deduced. The activation energy for the return of the

lower yield stress where in the interval an approximate logarithm of

I- 1Iq -I I I I I ,lll -- I I ;11H I I I 1111 I i I 1 111 I] -
448 .' --4 3730

0.8 --

3 0.2

0.0 I I
10 102 103 104 10 106
Time (sec)

Figure 10. The dependence of the LUders strain on time and temperature in Nickel 200.

I I I11 1 1 1 M4 I III1 I I IllIr I'll I I I-

Nickel 200
o 408
o 428
A 448
O 473


0 0
. . ..... I ,

. . .. I


105 106

101 102 103 104
Time (sec)

Figure 11. The approximate behavior of the secondary hardening component of the
lower yield stress increase (AoH= OExt ao).

I I I i 1 M I1 I I I'll I IiP r I . .

time behavior is exhibited was calculated on the basis of the respective

times to achieve a stress increase of 15.0 and 20.0 MPa. On this basis

the activation energy for the return of the lower yield stress in Nickel

200 is 25.2 3.2 and 26.4 2.7 kcal/mole.

The activation energy for the development of a 0.6% LUders

extension is similar, 24.3 kcal/mole. In addition, the shift of the

peaks of Ao versus log t is consistent with an activation energy of

approximately 22 kcal/mole. The downward trend of the aging curves behaves

in a manner corresponding to an activation energy of 29.0 kcal/mole on

the basis of the method of cuts at Ao= 23 MPa.

3.5 The Dependence of Ao and EL on Prestrain

To more fully characterize static strain aging, a series of

specimens were prestrained various amounts and then aged at 4480K for

a fixed time of 6000 seconds, the approximate time required to achieve

the maximum Ao at this temperature when the prestrain was 5% (see

Figure 7).

Figure 12 illustrates the dependence of Ao and EL on the amount

of deformation at a constant aging time and temperature. The Ac curve

shows that this parameter increases with prestrain, exhibits a broad

maximum and then decreases slowly. It is interesting that Stage III

of the work hardening behavior (see Section 3.8), as determined from

log O versus log o plots, begins at approximately 18% true strain at

2730K. Stage III is normally associated with dynamic recovery and in

view of the broad peak and subsequent decrease in Ao, it is possible

that Ao is reflecting the dynamic recovery.

1 30- 6

b 1.2

25 -

2 .8



0 5 10 15 20 25 30
Engr. Strain (%)

Figure 12. The dependence of Ao and EL on prestrain. Nickel 200 specimens were
prestrained at 2730K, aged for 6000 seconds at 4480K, and restrained
at 273'K.

The Liders extension increases continuously with prestrain as

indicated in Figure 12. The LUders strain is determined not only

by the size of the lower yield stress but also by the magnitude of

the work hardening rate. The latter decreases continuously with strain

and tends to make EL increase with strain. The fact that EL continues

to increase with E in Figure 12 is probably due to this cause. This

is similar to the case of Type A Luders bands which exhibit an increase

in LUders strain during plastic deformation [122].

3.6 Comparison of Nickel 270 and Nickel 200 Static Strain Aging

Several experimental observations indicate that the higher purity

Nickel 270 does not contain sufficient carbon to give rise to measurable

dynamic strain aging phenomena. Specifically, the Portevin-Le Chatelier

effect was not observed in this metal. Also, even at the highest

temperatures investigated, yield points or yield plateaus were not

observed in annealed material. Thus, Nickel 270 may be nearly repre-

sentative of pure nickel in terms of its mechanical properties.

To test if static strain aging is weakly exhibited in Nickel 270

a specimen was prestrained 5% at 273K and aged for 1200 seconds at

4730K. The resulting curve shown in Figure 13a appears to indicate some

aging since a short yield plateau is exhibited. However, the lower

yield stress increase for this specimen was only 4.76 MPa which is small

compared to the 23.00 MPa value for the commercial purity Nickel 200

specimen (Figure 13b) which was achieved in only one-half this aging time.

It is also possible that a significant portion of the 4.76 MPa yield

effect that has been observed in face-centered cubic metals [17,123].

Thus, one may be reasonably assured that strain aging phenomena in

Nickel 270 are generally weak.

Nickel 270

1200 s/4730


Nickel 200

-127 N 600s/473.



5% Pre-strain

Figure 13. (a) The yield return of a Nickel 270 specimen aged for a time to achieve a
maximum in Ao for Nickel 200. (b) Yield return for a Nickel 200 specimen
aged only one-half as long.

3.7 The Stress-Strain Behaviors

The details of the basic mechanical behavior of the high purity

nickel, Nickel 270, are shown in Figures 14-16. Those for commercial

purity Nickel 200 are shown in Figures 17-22. Nickel 200 unlike

Nickel 270 exhibits serrations and yield points.

As indicated in Figure 21, the Portevin-Le Chatelier effect was

observed over four orders of magnitude of strain rate in Nickel 200.

The figure also shows the approximate temperature intervals over which

Types A, B and C serrated flow were observed. Type C serrations were

sudden load drops appearing at regular intervals on the load-time

curve. The serrated flow intervals correspond closely to those observed

by Nakada and Keh [51] in nickel-carbon alloys indicating that the

presence of the manganese in Nickel 200 does not appreciably affect the

dynamic aging effects. The data of the present investigation are not

extensive enough to calculate the apparent activation energy for the

onset of serrations in Nickel 200 with accuracy. However, the data

appear to be consistent with the apparent activation energy 152 kcal/mole

calculated by Nakada and Keh for the somewhat purer alloys [51]. In

addition, the apparent activation energy for the disappearance of

serrations calculated by Nakada and Keh [51], 264 kcal/mole, also

appears to be reasonable for Nickel 200.

Figure 14 shows representative stress-strain curves for Nickel 270

obtained at several temperatures. The 0.2% offset flow stress and the

ultimate tensile strengths of Nickel 270 are plotted in Figure 15 as

functions of the temperature. Note that the ultimate stress decreases

monotonically with temperature without any undue irregularity. This

type of stress-temperature variation is characteristic of a metal which

E 12

0 77*K



0 0.1 0.2 0.3 0.4 0.5

Figure 14. True stress-true plastic strain curves for Nickel 270 4.2 x
Figure 14. True stress-true plastic strain curves for Nickel 270 (e = 4.2 x I0-4 s1).

0 100 200 300 400 500 600 700 800 900
T (K)

Figure 15. Variation of the 0.2% yield stress and the ultimate
tensile strength with temperature in Nickel 270
(E = 4.2 x 10-4 s-l.

o Uniform

60 0

50 -

40 -




0 100 200 300 400 500 600


700 800 900

Figure 16. Variation of the uniform and total elongation with
temperature in Nickel 270 (E = 4.2 x 10-4 s1).

does not exhibit pronounced dynamic strain aging. The total and

uniform elongations are illustrated in Figure 16 and show no anomalies.

A minimum in ductility was observed at 8350K and surface cracking was

noted on the specimens. Cracking also appeared on the specimen tested

at 8500K. The Nickel 270 specimens were annealed at 8680K and this

temperature was the upper testing limit. The elongations are reasonably

constant over a wide range of temperature (approximately 200 to 6500K).

A representative sample of Nickel 200 stress-strain curves are

shown in Figure 17. It should be noted that the curves at 300 and 5250K

were serrated. Only average stress-strain behavior can be shown in

these cases. Note that at 525K the curve shows an anomalously high

ultimate strength. The enhanced strengthening during work hardening

is best illustrated in Figure 18 which shows dependence of the 0.2% flow

stress and the ultimate stress on the temperature.

Because of the scale of the drawing in Figure 18 the 0.2% offset

stress appears to decrease monotonically with temperature. However,

a plot of the 0.2% offset stress for Nickel 200 at two strain rates with

the stress axis expanded as in Figure 19 shows that there is a small

yield stress plateau between approximately 300 and 4750K. This is

generally characteristic of dynamic strain aging in bcc and hcp metals.

The stresses were not normalized with respect to the elastic modulus, as

is customary, because nickel exhibits a large magnetostriction [124,125] and the

choice of modulus is uncertain below 626K, the Curie point [126].

This plateau is weakly exhibited compared to that of titanium [127-129],

for example. In comparison to the 0.2% stresses observed in Nickel 270,

Nickel 200 exhibits a larger temperature dependence (compare Figures 15

and 18).

NM 12-

S 10- 77. K
S 8 2K 192"K

S----------- 700"K

0 0.1 0.2 0.3 0.4

Figure 17. True stress-true plastic strain curves for Nickel 200 (c = 4.2 x 10-4 s-1).

Nickel 200

* 4.2 x 10-5 sec-1
-4 -1
* 4.2 x 104 sec1
0 4.2 x 10-3 sec-
S4.2 x 10-2 sec-


es \ \

400 -

0 0.2 I
I. 0-
* I I I I. -

200 400

600 800

Figure 18. The temperature dependence of the 0.2% yield stress
and ultimate tensile strength of Nickel 200.




_ I I 1 I I 1






I I *
Nickel 200
4.2 x 10-4 s"
o 4.2 x 10'3 s"

S 1 I I

200 400

~eo ~800

Figure 19. The temperature and strain rate dependence of the
0.2% yield stress in Nickel 200 on an expanded
stress axis.


An important feature to note in Figure 18 is that anomalous

strengthening is exhibited between 300 and 6000K. In Figure 20 the

strain dependence of the flow stress is shown for 5, 11, 19, and 31%

plastic strain. This figure demonstrates that only above approximately

11 to 19% plastic strain does the anomalous strengthening become

significant. That is, the strengthening effect is not due to anomalous

work hardening as Sukhovarov and Kharlova [67] previously suggested.

Note that unlike the behavior of the ultimate tensile strengths plotted

in Figure 18, the stress levels attained at 31% strain exhibit a rate

dependent shift in the peaks. These peaks have a rate dependence

corresponding to an activation energy of approximately 38 kcal/mole.

Figure 21 illustrates the temperature dependence of the uniform

and total elongations in Nickel 200 at several strain rates. The total

elongation at a strain rate of 4.2 x 10- s-l shows a mild ductility

minimum (blue-brittle effect) between approximately 300 and 450K. This

minimum is not well-defined. However, a well-defined but small reduction

in area minimum does occur in Nickel 200 that is strain rate dependent

as shown in Figure 22. This minimum was noted to have shifted in

accordance with an apparent activation energy of approximately 25 kcal/mole.

A reduction in area minimum is not always manifested in dynamic strain

aging [130]. It is interesting to note, however, that 25 kcal/mole is

approximately the activation energy for vacancy migration in nickel [109,118].

Not only is the loss in reduction in area not marked, but the lowest

value recorded is still above 75%. Adjunct observations of the fracture

surfaces under a low-power microscope did not reveal any striking

difference in fracture mode at the reduction in area minimum.


600- N.-- .

S500 -


300 5 %

200 -

4.2 x o10 s'
10 o 4. x 103 s-
4.2 x 10 s

0 100 200 300 400 500 600 700 800 900

Figure 20. Variation of the stresses at 5, 11, 19, and 30% plastic strain with
temperature in Nickel 200.

Nickel 200
4.2 x 10 sec'
4.2 x I0 sec' -
o 4.2 x 10-3 sec'-
s80 4.2 x 10-2 sec-

: ./ ,V/ Total

G e- '- o
\Q N. A'--- ,
G '. "

B .C. \" Uniform -

20 L A iABC
1 --- -------1 --- --I---'-

0 200 400 600 800

Figure 21. Variation of the uniform and total elongations with
temperature in Nickel 200. Also shown are the
approximate temperature ranges over which serrations
were observed at the respective strain rates.



96- Ni 270

94 ^ 4.2 x 1064s-

92 A


s88 Ni 200

8s .. ..... .. 0. / 42 x 1O- .s-

\. *\ 4.2 x lO s-1

\'\o I\
80 -

78 \ -


740 20 40 -O
0 200 400 500 SOO !000

Figure 22. Variation of the reduction in area
for Nickel 200 and Nickel 270.

with temperature

Note that the higher purity Nickel 270 shows more ductility at

all temperatures than does Nickel 200. Although a mild reduction in

area minimum does occur in Nickel 270, it is spread over a wide

interval between approximately 425 and 750K and is not as pronounced

as that in Nickel 200. Even the lowest reduction in area observed in

this investigation is considerably higher than that in most commercially

available bcc metals and alloys.

The strain rate sensitivity in Nickel 200 and Nickel 270 was not

investigated because the temperature interval of serrated flow in

Nickel 200 was some 3000K wide and it was believed that strain rate

changes conducted during discontinuous plastic flow would prove in-

conclusive. It was noted, however, that at room temperature during

moderately heterogeneous plastic flow, changes in rate resulted in the

appearance of flow stress transients in Nickel 200 and a steady state

strain rate sensitivity very close to zero was observed.

In summary, whereas the higher purity Nickel 270 shows no anomalies

in ultimate strength and elongation with increasing temperature, Nickel 200

between 300 and 6000K shows anomalous strengthening, serrated flow, a

small yield stress plateau and a mild elongation minimum.

3.8 The Work Hardening Behavior of Nickel 270 and of Nickel 200

Figures 23 and 24 show a cross-section of the log 0 versus log o

curves for Nickel 270 and Nickel 200, respectively, deformed at a

nominal strain rate of 4.2 x 10-4 s.1 These curves satisfactorily

represent the general trend of work hardening at all strain rates

investigated. No attempt has been made to draw in the straight lines

representing the stage behavior in order to reduce the complexity of

the figures.

o o
.. Oo



CD 9.0


850K 800*K

S (Stress in N/n)

7.6 7-8 8 0 8.2
LOG 0-

8.4 8.6 8.8
8.4 8.6 8.8

Figure 23. The log 0-log a curves of Nickel 270 (e = 4.2 x 10-4 s- ).


4 ,
* 4

0 4

0a *192-K
S 550'K

.700 K

* 0.




= 4.2x i04 s"7

\ee sooo" t **. e o

0 o0
0 0o

(Stress in MPa)

0o 0 0 o


525AK .
A On

o '650*K


* 850*K

7.8 8.0 82 8-4 8-6
LOG (o

8.8 9.0

Figure 24. The log 0-log a curves of Nickel 200 (- = 4.2 x 10-4 s ).


The work hardening behavior of Nickel 270 and of Nickel 200 are

similar to those found earlier in pure face-centered cubic metals. The

stages appear in the same manner as Zankl [75], Schwink and Vorbrugg [77],

and others [76,78-81] have shown to be the case for pure face-centered

cubic metals. It should be noted that hexagonal close packed metals

such as zirconium and titanium [131] as well as body-centered cubic

metals such as iron and niobium [132] show much higher m values (order

of 7 to 40) than the fcc alloys presently being investigated. That is,

all log o versus log a plots for these metals show much steeper slopes.

Also, these metals tend to show only one or two stages of deformation

behavior indicating that the deformation in these metals is possibly

controlled by a different set of deformation phenomena.

Figures 25 and 26 show the m values obtained for Nickel 270 and

Nickel 200 by measuring the slopes of log 0 versus log a plots at

different temperatures. The error of each particular mil value is

approximately 0.1 as measured by the plausible maximum and minimum

slopes that might conceivably characterize a particular work hardening

stage. It should also be noted that Stages I and III are difficult

to characterize in many cases. The parameter minI is plotted in the

figures only to show trends in the third stage as a function of temperature.

They are not accurately defined since in Stage III log O-log o plots

are not linear but curved. However, Stage II is generally uniquely

defined by a straight line on log O-log a curves.

Nickel 270 and Nickel 200 possess mil values which are close to 1.5

as shown in Figures 25 and 26. Note that in Nickel 200 mlI remains

constant over a wide range of temperatures and is strain rate independent.

8 ml

E 4 m


0 100 200 300 400 500 600 700 800 90S
T (

Figure 25. The variation of mi[ and milj with temperature in Nickel 270 (E = 4.2 x 10-4 s-1.

7- "' 4.2 x 10-5-'1
7 4.2 X 10-48-1
Nickel 200 o 4.2 x Io-33

0 f0
0 -
0 0 -- o
4o _- o-- o m0

2 *n

---" ^- ^

0 200 400 600 800 1000


Figure 26. The variation of m11 and mlII with temperature in Nickel 200
(E = 4.2 x 10-4 s-1).

It is important to emphasize that the apparent work hardening peak

observed in Nickel 200 (Figures18 and 20) is not a true work hardening

peak in the sense that Nickel 200 stress-strain curves show anomalous

increases in work hardening over a uniquely defined temperature interval.

The peak is associated with an increase in the uniform elongation

around the temperature of the ultimate stress peak. Figure 27, a plot

of (o5%-0.5%) versus temperature, readily shows this behavior. This

parameter (where o5% and o0.5% are the true stresses at 5 and 0.5%

true plastic strain, respectively) confirms that there is little

temperature dependence of the work hardening. The constancy of mil in

Figure 26 also substantiates this observation. Figures 28 and 29

show very clearly that the onset of Stage III is delayed over exactly

the same temperature interval as develops the peak in ultimate stress.

These facts imply that there is no prominent work hardening peak in

nickel as earlier authors have indicated [67]. There is, however, a

significant delay in the onset of Stage III or of dynamic recovery.

In Stage III deformation in both a polycrystalline or single

crystal fcc metal or alloy, the stress-strain curve shows a high

curvature, i.e., work hardening decreases very rapidly with continued

deformation. Thus, Stage III m values tend to be high. Figures 25 and

26 show that mIII increases with increasing temperature. This is

consistent with the concept that dynamic recovery becomes more important

with increasing temperature.

Now consider the extent or length of Stage II which is not illus-

trated well in the log 0-log o diagrams. In purer metals such as

Nickel 270 it is observed that Stage II, which is governed primarily by

slip on single slip systems within each grain [75], is lengthened by

0 100 200 300 400 500 600 700 800 900
T (K)

Figure 27. The variation of the work hardening parameter (5%-' 0.5%) with temperature
(E = 4.2 x 10-4 s-).

the presence of interstitial carbon as in Nickel 200. Figure 28

compares the approximate strain at which Stage II deformation begins,

E2, and that at which Stage III begins (and Stage II ends), E3, for

Nickel 200 and Nickel 270. The principal effect of interstitial

carbon in the DSA interval in this respect is to postpone Stage III

dynamic recovery processes to a later time in the deformation history.

Also the general level of the Nickel 200 curve is higher than that

for Nickel 270. It is not clear at this time how the carbon accounts

for the general delay in Stage III.

Nickel 200 exhibits an e3 maximum over the dynamic strain aging

interval. This maximum is closely related to the maximum exhibited

in the uniform elongation (Figure 21) and accounts for the strengthening

observed over this temperature interval. This maximum in e3 is strain

rate dependent as shown in Figure 29. Not only do the peaks shift to

higher temperatures with increasing strain rate but the peak heights

decrease as well. This is similar to the behavior of the flow stresses

described in Figures 19 and 21. The shift in peaks follows an approximate

activation energy of 37 kcal/mole, similar to that obtained for the

flow stress peaks.

50- 3 NICKEL 270
o C2

1 3 NICKEL 200

(All strains ore true plastic strains)
40 -


I s? ,. .= Z z--= ..... __ _o" o ....
0 100 200 300 400 500 600 700 800 900

Figure 28. The variation of E2 and E3 (the approximate strains
at which Stages II and III, respectively, begin) with
temperature for Nickel 270 and Nickel 200 deformed at
a strain rate of 4.2 x 10-4 s-1.

I I i I I- I I I I

* 42 x
* 4.2 x
o 4.2 x
> 4.2 x


id5 s'





0 200 400

600 800

Figure 29. The dependence of e3 on temperature and strain rate
in Nickel 200.



4.1 Rationale for Static Strain Aging in Nickel 200

The static strain aging data for Nickel 200 aged between 373 and

4730K after a 5% prestrain indicate:

1. The activation energy for the return of the lower yield stress

for times before the maximum in Ao is approximately 253 kcal/mole.

2. The kinetics of the rise in the lower yield stress can be

described by a In t time law.

According to the previous discussion three principal causes for

strain aging in other metals have been identified: (a) Cottrell pinning,

(b) Suzuki locking and (c) Snoek pinning.

It would appear on the outset that aging in Nickel 200 is such that

Cottrell pinning must be ruled out as a candidate mechanism. First,

the kinetics for aging as described by Cottrell and Bilby [9] are t2/3

In Nickel 200 the rise in Ao with time is much slower, approximately t /7

this is very different from the Cottrell-Bilby time law. Secondly, the

activation energy deduced to govern the early stages of pinning in

Nickel 200 is not consistent with the migration energy of carbon, the

principal impurity in Nickel 200, as the Cottrell-Bilby model would

suggest. The migration energy of carbon in nickel is approximately

35 kcal/mole (see Table 2) and thus significantly different from the

observed activation energy. Accordingly, the random migration distance

for carbon at all temperatures investigated during the time required to

cause maximum strengthening is only about 8b which is a very short

diffusion distance compared to distances involved in the formation of

a Cottrell atmosphere.

Suzuki pinning is not believed to be responsible for the observed

strain aging behavior in Nickel 200 primarily because the stacking fault

energy of pure nickel is very high. Hence, the probability of forming

stable faults in nickel is very low. It is believed that carbon

impurities do not lower the stacking fault energy enough to have a

significant effect with regard to this mechanism.

Snoek pinning by itself cannot account for the rise in Ao unless

anisotropic defects exist in the material in sufficient numbers and

have a sufficiently high interaction energy with dislocations. As will

be demonstrated the carbon-vacancy pair can satisfy these requirements.

However, Snoek pinning should occur in times, very short at the temperatures

investigated, compared to the times observed for the rise in Ao.

The formation of ordered carbon-vacancy pair atmospheres near

dislocations can account for the observed static strain aging behavior.

The very long times required to reach the maximum in Ao can be rationalized

in terms of a stress assisted migration of vacancies toward dislocations

where enhanced trapping by carbon occurs. The result is an accumulation

of ordered carbon-vacancy pairs near the dislocation and the concommitant

growth of an ordered carbon-vacancy pair atmosphere. In short, the

proposed model utilizes aspects of both the Cottrell-Bilby and the

Schoeck-Seeger theories of static strain aging.

In summary the following stages are envisioned: (a) At the end of

deformation vacancies diffuse to and become trapped in large measure

by adjacent carbon atoms. (b) The carbon-vacancy pairs thus formed

undergo Snoek ordering within the stress fields of the dislocations.

(c) Then a slow migration of vacancies toward the dislocation develops

an atmosphere of carbon-vacancy pairs near the dislocation.

(d) Eventually the vacancy concentration decreases because of losses

to sinks. This reduces the pair concentration and causes an eventual

loss of strengthening. These stages will now be considered in detail.

4.2 The Mechanism of Static Strain Aging Exhibited in Nickel 200

4.2.1 The Distribution of Vacancies, Carbon Atoms and Dislocations
After Plastic Deformation

After the prestrain of 5%, the vacancy concentration in nickel

is approximately 10-5 atomic fraction as given by Eq. 9. The

particular heat of Nickel 200 used in this investigation contains

approximately 0.1 w/o or 0.5 a/o carbon. Thus, at 5% strain the carbon

to vacancy ratio is approximately 500 to 1. If the carbon and the

vacancies are distributed at random throughout the nickel lattice, then

the mean carbon atom spacing will be approximately 6b; the mean vacancy

spacing will be approximately 50b.

After plastic deformation to 5% strain, the dislocation density [133,134]

is estimated to be between 3 and 8 x 109 cm-2. Assuming an equidistant

array of straight dislocation lines as a first approximation to the

dislocation configuration, then their mean spacing is approximately

450 to 730b.

4.2.2 Vacancy Trapping by Carbon Atoms

In the absence of other defects such as grain boundaries or

dislocations, vacancies in a metal are attracted to oversized impurity

atoms by a hydrostatic pressure gradient. The strain energy released

when a vacancy is moved from an infinite distance away from an impurity

to the impurity is known as the binding energy of the vacancy-impurity

pair. This phenomenon is commonly called trapping and it is generally

accepted that impurity atoms can trap vacancies in metals. For example,

it has been shown that the presence of carbon in austenitic stainless

steel forestalls radiation induced swelling in certain critical nuclear

reactor parts. This indicates that void formation is hindered by the

presence of carbon due to trapping of vacancies [135]. Trapping of

carbon by vacancies in irradiated mild steel has been demonstrated as

well [136].

4.2.3 The Concentration of Carbon-Vacancy Pairs

After deformation in Nickel 200 and before aging of a specimen,

it is reasonable to assume that vacancies become trapped by carbon

atoms initially. This probably occurs very rapidly even at 3730K, the

lowest aging temperature used, since the vacancies do not have to travel

far to become trapped by carbon atoms. Vacancies in nickel are very

mobile as evidenced by their approximate 0.8 to 0.9 eV migration energy

and assuming that an equilibrium is established between the carbon-vacancy

pairs and free vacancies, their concentrations may be calculated on

the basis that the binding energy of the carbon-vacancy pair can be

estimated to be 0.2 to 0.3 eV per pair from diffusion data (Section 1.4).

For the moment, it is assumed that the dislocations in the metal do not

affect the concentrations. The concentrations of single (free) vacancies

(c ) and bound vacancies in carbon-vacancy pairs (c ) vary with carbon

concentration (Cc), temperature T, and carbon-vacancy binding energy (B)

according to the equations [116]

ccv = Z cf eB/kT (11)


cv cf + (12)
v= CC V

where cv is the vacancy concentration generated during plastic defor-

mation, cc is the free carbon concentration and Z is 6, the nearest

neighbor coordination number for a carbon atom in an octahedral site

of the fcc lattice.

Substituting Eq. 12 into Eq. 11 gives

cc = Z(cc-Ccv)(cvcv )eB/kT (13)

For c = 5 x 103 and c = 105 one may write on the basis that
c v

Ccv = Zcc(c -ccv)eB/kT (14)

Substituting a value for B equal to 0.3 eV and a temperature of

408K into Eq. 14 gives the carbon-vacancy pair concentration of

9.93 x 10-6, very close to the total vacancy concentration of 10-5.
The free vacancy concentration from Eq. 12 is then 6.92 x 108 atom

fraction. Thus, almost all (99.3%) vacancies are bound to carbon atoms

based upon the choice of B.

In the absence of carbon atoms the concentration of vacancies at

thermal equilibrium is given by

cv e (15)

where Qf is the vacancy formation energy. At 4080K and a Qf of

approximately 39 kcal/mole (1.7 eV) [118] the vacancy concentration

is approximately 1.8 x 1021 atom fraction. This is many orders of

magnitude smaller than the free vacancy concentration calculated above

and suggests that the deformed metal is not in its lowest energy state

with regard to numbers of vacancies. Thus, there exists a distinct

tendency for vacancies to migrate to sinks. Because of the high binding

to carbon atoms, vacancies spend a large portion of their migration time

bound to carbon atoms and, hence, the process of annealing to dislocations

and grain boundaries is sluggish and takes a considerable time.

Consider next the influence of a dislocation on the carbon-vacancy

pair concentration. It has been demonstrated that the carbon-vacancy

pair may be visualized as an elastic dipole which can reorient in the

strain field of a dislocation. In Eqs. 11-14, the binding energy, B,

is completely general. That is, in the specific case which includes

Snoek ordering, B becomes position dependent. Specifically,

B = B' + u(r) (16)

where B' is the binding energy of a carbon-vacancy pair and u(r) is

the position dependent interaction energy of a carbon-vacancy pair with

a dislocation. Thus, the effect of u(r), which varies as r- is to

enhance the binding of vacancies to carbon atoms near the dislocation.

This will be explained in detail presently.

4.2.4 Theory of Schoeck Locking by Carbon-Vacancy Pairs

In the case of nickel-carbon alloy, the carbon-vacancy pair acts

as a dipole (Section 1.4). Appendix A presents a calculation of the

approximate interaction energy between an assumed carbon-vacancy

dipole and a screw dislocation in a fcc lattice. The interaction

calculation related to an edge dislocation is more complicated [137].

However, as indicated first by Nabarro [138] and later expanded upon

by Cochardt et al. [137], the strength of interaction between a dis-

location and impurity atoms may be assumed to be very similar whether

the dislocation is in the screw or edge orientation.

Let us consider the case of locking of a screw dislocation by an

ordered atmosphere of dipoles in the fcc lattice. The reasoning is

analogous to that for the calculation made by Schoeck and Seeger [28]

for carbon (or nitrogen) in alpha iron. However, the interaction

energy for a dipole in a fcc lattice as shown in Figure 35 in Appendix A

is somewhat different. The carbon atoms in nickel are assumed to occupy

octahedral interstitial sites, e.g., the body-centered position of the

fcc unit cell. The six nearest neighbors are face-centered atoms. A

dipole is formed when a vacancy is situated on a face-centered position

as schematically shown in Figure 30. Thus, the carbon-vacancy dipole

may be oriented in one of three possible 100 type directions. These

positions are denoted by an orientation number 1, 2 or 3. The dislocation

line is assumed to lie along one of four 101 type directions.

It is shown in Appendix A that for the specific case of a screw

dislocation, each of the three possible orientations of the carbon-

vacancy dipole interacts differently with the dislocation. This gives

rise to three possible interaction energies and one of the three

orientations is most stable. That is, a carbon-vacancy dipole in,

say, the 1-orientation can flip to another orientation (say, 2) and

the result is that the interaction energy between that particular dipole

Figure 30. A schematic illustration of the assumed configuration
of the carbon-vacancy pair and the three possible
independent orientations that it may assume.

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