International Journal of Fatigue 26 (2004) 899–910
www.elsevier.com/locate/ijfatigue
The tensile and low-cycle fatigue behavior of cold worked
316L stainless steel: influence of dynamic strain aging
Seong-Gu Hong, Soon-Bok Lee Department of Mechanical Engineering, Korea Advanced Institute of Science and Technology, 373-1 Guseong-dong, Yuseong-gu,
305-701 Daejeon, South Korea.
Received 18 March 2003; received in revised form 8 October 2003; accepted 5 December 2003
Abstract
v
Tensile and low-cycle fatigue (LCF) tests were carried out in air in a wide temperature range from 20 to 750 C and strain rates
of 1 104 1 102 =s to investigate the influence of strain rate on tensile and LCF properties of cold worked 316L stainless
steel, especially in the dynamic strain aging (DSA) regime. Dynamic strain aging caused a change in the tensile properties such as
v
strength and ductility in the temperature range from 250 to 600 C. This temperature range coincided well with the regime of
negative strain rate sensitivity. Cyclic stress response at all test conditions was characterized by an initial hardening during the
few cycles, followed by gradual softening until failure. Temperature and strain rate dependence on cyclic softening appears to
come from a change of the cyclic plastic deformation mechanism and the DSA effect. The regimes of DSA between tensile and
LCF loading conditions in terms of the negative strain rate sensitivity were consistent with each other. At the elevated temperature, a drastic reduction in fatigue resistance was observed, and this is attributed to the effects of oxidation, creep, and dynamic
strain aging and interactions among these factors. The regime of DSA accelerated a reduction in fatigue resistance by enhancing
crack initiation and propagation.
# 2004 Elsevier Ltd. All rights reserved.
Keywords: Low-cycle fatigue; 316L stainless steel; Dynamic strain aging; Cyclic softening; Cold work
1. Introduction
Type 316L stainless steel, which has been favored as
a structural material in several high temperature components such as the primary side of liquid metal cooled
fast breeder reactor (LMFBR), stands as a prominent
material for the next generation power facilities. This is
it offers a good combination of high temperature tensile and creep strength, corrosion resistance, coupled
with enhanced resistance to sensitization and associated
intergranular cracking. In LMFBR applications, the
components are exposed to severe conditions such as
v
high temperatures (ranging from 300 to 600 C) and
undergo temperature gradient induced cyclic thermal
stresses as a result of repeated start-up and shut-down.
Low-cycle fatigue (LCF) represents a predominant
Corresponding author: Tel.: +82-42-869-3029; fax: +82-42-8693210.
E-mail address: sblee@kaist.ac.kr (S.-B. Lee).
0142-1123/$ - see front matter # 2004 Elsevier Ltd. All rights reserved.
doi:10.1016/j.ijfatigue.2003.12.002
failure mode: specific attention on LCF is needed in the
design and life assessment of such components. In austenitic stainless steels [1–5], a reduction in fatigue
resistance with increasing temperature has been reported. This was ascribed to a change in the cyclic plastic
deformation mechanism, creep, and oxidation effects or
interactions among these factors. It has also been
v
noticed that in the temperature range of 300–600 C,
dynamic strain aging (DSA) induces a change in
material properties such as strength and ductility, and
thus a detailed research on influence of DSA is
required in this temperature region.
Dynamic strain aging is the phenomenon of interactions between diffusing solute atoms and mobile dislocations during plastic deformation. It depends on
deformation rate and temperature, which governs the
velocity of mobile dislocations and diffusing solute
atoms. At low temperatures and high strain rates, the
solute velocity is too slow compared with the dislocation velocity to cause a strain aging. In contrast, at
900
S.-G. Hong, S.-B. Lee / International Journal of Fatigue 26 (2004) 899–910
high temperatures and low strain rates, a solute atmosphere can move with the dislocations, and the effects of
DSA disappear. Therefore, DSA manifests itself in the
range of intermediate strain rates and temperatures [6].
Dynamic strain aging is manifested by a serrated flow
in the stress–strain curve, peaks or plateau in the variation of material strength with temperature, minima in
the variation of ductility with temperature, and negative strain rate sensitivity (SRS).
Mannan et al. [7] investigated DSA in 316 stainless
steel. At low temperatures, the diffusing velocity of the
interstitial solute atoms such as carbon and nitrogen is
fast enough to move and age the dislocations. However, at high temperatures, cyclic plastic deformationinduced vacancies enhance the diffusion of chromium
atoms, causing the aging of dislocations by the chromium atmospheres.
The most common alloys showing dynamic strain
aging are carbon steels, and is well established for
monotonic tensile loading. However, few systematic
investigations have been reported about the influence
of DSA on low-cycle fatigue behavior, and the results
are contradictory [1–5,8–10]. Abdel-Raouf et al. [9]
examined the elevated temperature fatigue deformation
characterization of Ferrovac E iron containing 0.007%
carbon and reported that DSA causes a reduction in
fatigue life. In contrast, in 2 1/4Cr–1Mo steel [10], an
increase in fatigue life was observed in the temperature
v
range of 427–600 C due to an interactive solid solution hardening resulting from cyclic stress enhanced
precipitation of Mo2C.
In this study, tensile and low-cycle fatigue tests were
v
carried out in the temperature range (20–750 C) and
strain rates of 1 104 1 102 =s for 17% cold
worked 316L stainless steel. The test results are analyzed and discussed from the following viewpoints: (1)
based on results of the tensile test, the influence of temperature and strain rate on tensile properties and tensile
fracture mechanism, focused on the DSA regime, was
investigated, and the temperature region where DSA
operates was determined. (2) A change in cyclic softening behavior with temperature and strain rate was analyzed in terms of cyclic plastic deformation mechanism
and DSA effect. (3) The DSA regimes for the tensile and
fatigue loading conditions were compared with respect
to the strain rate sensitivity. (4) The reasons for a
reduction in fatigue resistance at elevated temperatures
are discussed by considering the influence of dynamic
strain aging on crack initiation and propagation.
2. Experiment
2.1. Material and specimen
The material used in this study was 17% cold
worked (CW) 316L stainless steel having a chemical
composition (wt%): C 0.025, Si 0.41, Mn 1.41, P 0.025,
S 0.025, Ni 10.22, Cr 16.16, Mo 2.09, and N 0.043. The
as-received material was manufactured by the following
v
process: material was solution-treated at 1100 C for
40 min, then water-quenched, and finally a round bar
having a diameter of 16 mm was made through the
cold drawn process which introduced a 17% prestrain.
This treatment yielded an average grain size of 44.2
lm. The as-received material was fabricated into the
dog-bone type specimens having a gauge length of 36
mm and a gauge diameter of 8 mm in accordance with
ASTM standard E606-92. The specimen surface was
polished along the longitudinal direction, using emery
paper down to #2000 (13 lm), so as to remove surface defects such as machining marks.
2.2. Test equipment
A closed-loop servo-hydraulic test system with 5-ton
capacity was used to accomplish the tensile and fatigue
tests. A three-zone resistance type furnace which conv
trols temperature within a range of 1 C at steady
state was used for temperature control. A high temperature extensometer (MTS model no.: 632-13F-20,
gauge length: 25 mm) was used to control strain signal
and measure strain.
2.3. Test procedures
Tensile tests were conducted in laboratory air as a
constant cross-head speed of 0.2, 2, and 20 mm/min in
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the temperature range from 20 to 750 C. For displacement control (that is, cross-head control), it was
impossible to measure the exact value of strain along
the gauge length of a specimen during deformation.
Thus, the displacement signal was converted to a strain
signal by dividing it by the effective gauge length. An
extensometer which has a travel length of 10% was
used to determine the effective gauge length of the
specimen. Cross-head speeds of 0.2, 2, and 20 mm/min
correspond to strain rates of 1 104 , 1 103 , and
1 102 =s, respectively.
Low-cycle fatigue tests were carried out in laboratory air and under fully reversed total axial strain control mode using a triangular waveform at 400, 550,
v
600, and 650 C. A strain amplitude of 0.5% was used
for all the tests, and the strain rates were 1 104 ,
1 103 , and 1 102 =s. The displacement, load and
strain signals were measured for each cycle. One cycle
comprises of 200 data points. Fatigue life of the specimen was defined as a 30% load drop point of a stabilized peak load at half-life. Details of the experimental
procedure can be found elsewhere [11].
S.-G. Hong, S.-B. Lee / International Journal of Fatigue 26 (2004) 899–910
3. Results and discussion
3.1. Tensile test
3.1.1. Serrated flow in stress–strain curves
The segment of stress–strain curves at a cross-head
speed of 2 mm/min (_e 1 103 =s) is presented in
Fig. 1. As temperature increases, the induced stress for
the same strain value decreases, and Type A, B, D, E
serrations are observed in the stress–strain curve.
According to Mannan et al. [7], Type A, B, and C serrations have been observed for 316 stainless steel.
However, in the present study, Type D serration
occurred at low temperatures and A, B, E serrations at
high temperatures. Type A and B serrations occurred
at low strain and change over to Type E serration at
values of high strain. The condition of temperature and
strain rate for serration is shown in Fig. 2 and is a
function of temperature and strain rate. The regime of
serration moved to a higher temperature with an
increase in strain rate. Type D serration was first
v
observed at 300 C and e_ 1 104 =s, and Type A, B,
v
E serrations occurred at 700 C and e_ 1 102 =s.
Therefore, the regime of DSA, in terms of serrated flow
v
in stress–strain curves, was from 300 to 700 C and
depended on the strain rate.
3.1.2. Strain rate sensitivity and tensile properties
The effect of temperature and strain rate on 0.2%
yield strength and true ultimate tensile strength (UTS),
defined as the ‘(1 þ en )(engineering UTS)’ where en
represents an engineering strain at onset of necking, is
shown in Fig. 3(a) and (b). The results reveal a lack of
remarkable effect of strain rate on 0.2% yield strength.
However, the true UTS was significantly affected by
the strain rate. As shown in Fig. 3(b), the influence of
Fig. 1.
901
strain rate on true UTS can be categorized into three
different temperature regions: in region I, below
v
250 C, there was no significant effect of strain rate. In
v
region II, over 250–600 C, a negative strain rate versus stress response, which means that strain hardening
decreases as the strain rate increases, was observed. In
v
region III, above 600 C, a positive strain rate versus
stress response, which means that strain hardening
increases as the strain rate increases, was observed.
According to van den Beukel [12], dynamic strain aging
is manifested by negative SRS value, which is defined
as dr=dln_e. The value of SRS, defined as Dr=Dln_e, was
calculated at each temperature, and the result is presented in Fig. 3(c). The SRS value becomes negative in
v
the temperature range 250–600 C, and this temperature region is considered as the regime of DSA. In general, a negative SRS regime involves serration. In the
present study, it was valid at low temperatures but not
at high temperatures. At high temperatures, the regime
of serration was higher than the negative SRS regime.
When the influence of temperature on material
strength (0.2% yield strength and true UTS) is examined, material strength decreases with an increase
in temperature. There exists a temperature region
v
(250–600 C) where a reduction of strength with
increasing temperature is retarded or even slightly
increased (Fig. 3(a) and (b)). Such an increase of
v
strength from 250 to 600 C was attributed to DSA,
and is quite consistent with the negative SRS regime.
Dynamic strain hardening stress, defined as the difference between true UTS and 0.2% yield strength,
increases in the DSA regime since DSA effect is more
remarkable on UTS than on yield strength. The variation of dynamic strain hardening stress with strain
rate and temperature is shown in Fig. 3(d). A plateau
Segments of the stress–strain curves from tensile tests at e_ 1 103 =s.
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S.-G. Hong, S.-B. Lee / International Journal of Fatigue 26 (2004) 899–910
Fig. 2. Regimes of occurrence of serrated flow in the stress–strain curves.
or slightly increasing regions were observed from 250
v
to 600 C and correspond well with the negative SRS
regime. Dynamic strain hardening stress increased with
a decrease in strain rate (that is, the DSA effect
increased with a decrease in strain rate), and can be
explained that low strain rate reduces the velocity of
mobile dislocations. Thus, slow moving dislocations
can easily be aged by the solute atmosphere.
The minima in the variation of ductility with temv
perature was observed from 250 to 600 C and moves
to a higher temperature region with an increase in
strain rate (Fig. 3(e)). The temperature region of minimum ductility is consistent with the plateau in the variation of strength with temperature and the negative
SRS regime. However, it shows some differences with
the serration regime. The strain rate dependence on
v
ductility of the material was negligible below 300 C.
v
Above 300 C, ductility of the material increases with
decreasing strain rate, and gets more accelerated as
temperature increases.
The DSA of 17% CW 316L stainless steel occurred
v
in the temperature range 250–600 C and is manifested
by a plateau in the variation of strength and dynamic
strain hardening stress with temperature, a minima in
the variation of ductility with temperature, and a negative SRS regime. However, it shows marginal difference
in the occurrence of the serration regime.
3.1.3. Tensile fracture mechanism
The fracture mechanism for monotonic tensile loading conditions was studied by observing the tensile
fracture surfaces in a scanning electron microscopy
(SEM). Results reveal a change in tensile fracture
mechanism with temperature and strain rate (Fig. 4). A
classic cup and cone type fracture in the non-DSA
regime, which exhibits three zones (Fig. 4(a)): the inner
flat fibrous zone where fracture initiates, an intermediate radial zone, and an outer shear-lip zone where final
v
fracture occurs along a direction of 45 with the loading axis.
In the DSA regime, the fibrous zone size is reduced,
compared with that at other temperatures (Fig. 4(b)–(d)),
and it is consistent with the result of reduction in ductility in the DSA regime. In addition, the fibrous zone in
the DSA regime was twisted at final fracture.
3.2. Low-cycle fatigue test
3.2.1. Cyclic stress response
Cyclic softening behavior was observed through a
whole life except during the initial few cycles. This is
attributed to 17% tensile prestrain, which was introduced in the as-received material by cold-working.
Cyclic softening has been reported in materials, which
are manufactured by cold work process or have high
initial dislocation density [13–16]. In general, cyclic
behavior of material resulted from an interaction
between the generation of dislocations during cyclic
plastic deformation and the annihilation by recovery
processes. Cyclic softening of the cold-worked materials which have a high initial dislocation density, occurs
when the annihilation rate of dislocations is greater
than their generation rate, causing a net decrease in the
dislocation density, or when a rearrangement of the
dislocation structure takes place, resulting in an
increase in the mean free path of dislocations [13,16]. It
is suggested that this mechanism is responsible for the
observed softening. As shown in Fig. 5, the observed
cyclic softening behavior for all test conditions is classified into three regions: in region I, cyclic hardening is
observed and peak stress reaches a maximum value
S.-G. Hong, S.-B. Lee / International Journal of Fatigue 26 (2004) 899–910
903
Fig. 3. The variation of tensile properties of 17% CW 316 SS with temperature and strain rate; (a) 0.2% yield strength, (b) true ultimate tensile
strength, (c) strain rate sensitivity, (d) dynamic strain hardening stress, (e) elongation.
after the initial few cycles. In region II (about 90% of
fatigue life), cyclic softening occurs and continues up
until macrocrack initiation point where a steep decrement in peak stress occurs. In region III, the steep
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S.-G. Hong, S.-B. Lee / International Journal of Fatigue 26 (2004) 899–910
Fig. 4.
SEM micrographs of tensile fractured specimens at e_ 1 104 =s; (a) 200 C, (b) 300 C, (c) 400 C, (d) 550 C.
reduction in peak stress occurs because the effective
load-bearing area is reduced due to the propagation of
macrocracks. Static recovery, which is introduced during the thermal stabilizing process of a specimen before
tests [11], is thought to be the reason for the occurrence
of cyclic hardening in region I.
Cyclic stress response of 17% CW 316L stainless
steel is a function of temperature and can be categorized into two distinct regions. A negative strain rate
v
stress response in 400–600 C, transition behavior at
v
600 C, and positive strain rate stress response above
v
600 C. Such a tendency of cyclic stress response is
consistent with that in the monotonic tensile loading
conditions. The values of SRS were calculated at each
temperature using the stress–strain relations at half-life
(Fig. 6(a)) and compared with the tensile loading conditions (Fig. 3(c)). The results reveal a negative SRS
regime (DSA regime) between tensile and LCF loading
conditions coincides with each other in the temperature
v
range from 400 to 600 C, and SRS values under the
two loading conditions are similar. As shown in Fig. 5,
at each temperature, the SRS is maintained at a con-
v
v
v
v
stant value from the 20% of total life (N=Nf ¼ 0:2) to
the initiation of region III (N=Nf ¼ 0:9) for all test
conditions. This means that a material is stabilized,
when SRS is maintained at a constant value. The negative strain rate versus stress response in the DSA
regime appears to result from an increase in dislocation
density during deformation. An increase in the dislocation density with decreasing strain rate has been reported in the DSA regime for 316L(N) stainless steel [5].
During DSA, slow moving dislocations become easily
aged by the solute atmosphere. Thus, additional dislocations are generated to maintain the imposed deformation rate. This mechanism causes an increase in the
dislocation density. The matrix is hardened by an
increase in dislocation density, causing an increase in
flow stress needed to impose the same total strain during successive cycles.
Fig. 6(b) shows a change in plastic strain amplitude
Dep =2, at a total strain amplitude of 0.5%, with temperature and strain rate. Dep =2 increased in the DSA
v
regime 400–550 C with an increase in strain rate, and
v
the reverse was observed at 650 C. It is opposite to
S.-G. Hong, S.-B. Lee / International Journal of Fatigue 26 (2004) 899–910
Fig. 5.
v
v
905
v
v
Cyclic stress responses of 17% CW 316L SS at Det ¼ 0:5%; (a) 400 C, (b) 550 C, (c) 600 C, (d) 650 C.
strain rate dependence on peak stress (Fig. 6(a)) since
plastic strain amplitude is calculated from the difference between total strain amplitude and elastic strain
amplitude like Eq. (1):
Dep Det Dee Det Dr
¼
¼
2E
2
2
2
2
increases with a decrease in strain rate. This is consistent with the fact that fatigue resistance reduces with
decreasing strain rate. However, temperature dependence of plastic strain energy density is ambiguous and
exhibits no consistent tendency.
ð1Þ
where Dr=2 and E represent stress amplitude and elasv
tic modulus. A lower Dep =2 was induced at 400–550 C
due to an increase in strength by DSA, and a
v
maximum value is achieved at 650 C.
According to Morrow [17], the plastic strain energy
per cycle can be regarded as a composite measure of
the amount of fatigue damage per cycle since cyclic
plastic strain is related to the movement of dislocation,
and cyclic stress is related to the resistance to their
motion. The fatigue resistance of a metal can be characterized in terms of its capacity to absorb and dissipate plastic strain energy. Plastic strain energy density
per cycle (DWp) calculated at half-life is shown in
Fig. 6(c). DWp increases with decreasing strain rate at
each temperature; that is, the fatigue damage per cycle
3.2.2. Cyclic softening behavior
As mentioned in Section 3.2.1, cyclic softening was
observed at all test conditions and depended both on
temperature and strain rate. The parameter ‘‘softening
ratio’’ defined as Eq. (2) was introduced to quantify the
amount of cyclic softening during low-cycle fatigue.
Softening ratio ¼
rmax rmax jNf =2
rmax
ð2Þ
where rmax and rmax jNf =2 represent a maximum peak
stress and a peak stress at half-life, respectively.
The variation of softening ratio with temperature under
Det ¼ 0:5% and e_ 1 103 =s is presented in
Fig. 7(a). Softening ratio (that is, the amount of cyclic
softening) increases with an increase in temperature
v
v
in the range from 20 to 200 C and above 600 C.
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S.-G. Hong, S.-B. Lee / International Journal of Fatigue 26 (2004) 899–910
Fig. 6. The variation of LCF properties of 17% CW 316L SS at half-life with temperature and strain rate at Det ¼ 0:5%; (a) peak stress, (b) the
ratio between Dep and Det, (c) plastic strain energy density.
v
However, it decreases from 200 to 600 C where DSA
operates. The dependence of softening ratio on temperature appears to result from a change in dislocation
structure, which develops during LCF deformation. It
has been reported that the dislocation structure of austenitic stainless steels involving 316L(N) stainless steel
changes from cell structure at temperatures below 300
v
v
C to a planar one between 300 and 600 C (DSA
regime) and back to cell or subgrain structure beyond
v
600 C [1–5]. In the DSA regime, the mechanism of
interaction between mobile dislocations and the solute
atmosphere restricts dislocation cross slip and favors
planar slip deformation, causing an increase in flow
stress needed to impose the same total strain, and thus
a decrease in the amount of cyclic softening, compared
with other temperatures, occurs. In the non-DSA
v
v
regime from 20 to 200 C and above 600 C, cell or
subgrain structures develop during low-cycle fatigue
deformation, and its size increases with an increase in
temperature, causing an increase in the mean free path
of dislocations. This results in an increase in the
amount of softening with increasing temperature.
Details of the effect of temperature and strain amplitude on cyclic softening behavior of 17% CW 316L
stainless steel can be found elsewhere [18].
The variation of softening with strain rate in the
v
temperature range 400–650 C is shown in Fig. 7(b). In
v
the DSA regime (400–550 C), the reduced strain rate
introduced a decrease in the amount of softening. This
is attributed to an enhanced DSA effect at the lower
strain rate. However, in the non-DSA regime above
v
600 C, the opposite effect was observed since effect of
DSA is excluded and thermally assisted recovery is
v
enhanced at the lower strain rate. At 600 C, a transition behavior was observed.
S.-G. Hong, S.-B. Lee / International Journal of Fatigue 26 (2004) 899–910
Fig. 7.
907
The variation of softening ratio with temperature and strain rate at Det ¼ 0:5%; (a) temperature dependency, (b) strain rate dependency.
3.2.3. Low-cycle fatigue life
The influence of temperature and strain rate on fatigue life at Det ¼ 0:5% is shown in Fig. 8. As shown in
Fig. 8(a), a drastic reduction in fatigue life with
increasing temperature was observed at each strain
rate. According to the previous research [18], there is a
linear relation between the logarithm of fatigue life and
test temperature at a strain rate of 1 103 =s like
Eq. (3):
dlogNf ðTÞ
¼ C
dT
ð3Þ
where the slope C means a reduction rate of fatigue
resistance with increasing temperature and a higher
value of C implies a rapid reduction in fatigue resistance. Eq. (3) can also be applied to the results of strain
rates of 1 102 and 1 104 =s. The predicted fatigue
life using Eq. (3) is plotted as a solid line in Fig. 8(a).
The slope (C) is a function of strain rate and has a
value of 0.0021, 0.00152, and 0.00158 at each strain
rate.
A steep reduction in fatigue resistance with an
increase in temperature or decrease in strain rate can
be thought to come from the influence of oxidation
and creep, which becomes important at elevated temperatures. Driver et al. [19] have reported a reduction
in the fatigue resistance of type 316L(N) stainless steel
in laboratory air with an increase in temperature
compared with that in a vacuum. This was attributed
to oxidation enhanced Stage I transgranular crack
initiation. The oxidation effect becomes profound at
the low strain rate. Considering the homologous temv
perature, 400 C corresponds to about 0.36Tm (Tm:
melting temperature), and the creep effect becomes
important as temperature increases. A reduced strain
rate introduces an increase in tensile loading time per
cycle, where harmful damage is accumulated, and total
test time, which results in an increase in creep damage.
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S.-G. Hong, S.-B. Lee / International Journal of Fatigue 26 (2004) 899–910
Fig. 8. The reduction of fatigue resistance with temperature and
strain rate at Det ¼ 0:5%; (a) temperature dependency, (b) strain
rate dependency.
When the reduction of fatigue resistance observed in
the present study is explained by the effects of oxidation and creep, fatigue resistance at each temperature should reduce with decreasing strain rate, and the
reduction rate in fatigue resistance with decreasing
strain rate should be more severe at the higher temperatures. However, as shown in Fig. 8(b), the opposite
was observed; that is, the amount of reduction in fatigue resistance with reduced strain rate decreased with
an increase in temperature. This means that there is
another factor, which affects fatigue resistance besides
oxidation and creep, and is the DSA effect. In the
investigation on Alloy 800H [2], the effects of DSA on
crack initiation life and evolution of the microcrack
density with the number of cycles have been studied
and compared with results in vacuum, to exclude the
effect of environment on crack initiation. The results
reveal that DSA leads to an overall decrease in crack
initiation life, an appreciably higher crack density, and
reduction in fatigue life. In other research [3,4], it has
been noticed that DSA causes a faster reduction in fatigue resistance by way of rapid crack propagation
where transgranular fracture is dominant. The negative
strain rate stress response developed in the DSA regime
introduces a higher stress concentration at the crack tip
during cyclic deformation as the strain rate decreases,
causing an increase in crack growth rate. This leads to
a reduction in crack propagation life. It is well known
that crack propagation stage is dominant in LCF conditions of ductile materials, compared with crack
initiation stage, and thus the reduced crack propagation life implies a reduction in fatigue life. Therefore,
it can be inferred that the steep reduction in fatigue life
with reduced strain rate, in the temperature range from
v
400 to 550 C (DSA regime), compared with other
temperatures, is attributed to dynamic strain aging,
which accelerates the reduction of fatigue resistance. In
v
the non-DSA regime above 600 C, the DSA effect is
weakened or eliminated, and thus a slower reduction in
fatigue resistance with decreasing strain rate resulted.
The evidence of DSA-induced fatigue resistance
reduction is manifested in Fig. 8(a). When the reduction
of fatigue resistance with increasing temperature is
examined at each strain rate, a steep reduction in fatigue life is observed at the high strain rate of
1 102 =s, compared with other strain rates. It is contrary to expectation since oxidation and creep effects
will become more significant at low strain rate and
high temperature. From the tensile test results, the
regime of DSA is a function of strain rate and temperature, and moves to a higher temperature region as
strain rate increases. During LCF deformation, effect
v
of DSA is weakened or removed beyond 600 C at the
low strain rates below 1 103 =s. However, at the high
strain rate of 1 102 =s, the DSA effect is still valid
v
even at 650 C. Hence, a severe reduction in fatigue
resistance with an increase in temperature at the high
strain rate of 1 102 =s can be thought to result from
the DSA effect.
3.2.4. Failure mechanism in low-cycle fatigue loading
conditions
Fracture surfaces of LCF failed specimens were
examined using SEM to study the failure mechanism in
LCF loading conditions. Crack initiation occurred on
the external surface in the specimen and was enhanced
in the DSA regime, causing multiple crack initiation
sites on the border of the fracture surface (Fig. 9(a)).
v
At all test conditions (400–650 C and 1 104 1
102 =s), crack propagated inward in the transgranular
mode with the formation of fatigue striations (Fig. 9(b),
(d) and (e)). It has been reported in SA508 [3] that
dynamic strain aging would increase the degree of
S.-G. Hong, S.-B. Lee / International Journal of Fatigue 26 (2004) 899–910
909
Fig. 9. SEM micrographs of LCF failed specimens at Det ¼ 0:5%; (a) 550 C, e_ ¼ 1 102 =s, (b) higher magnification view of ‘‘A’’ in (a), (c)
v
650 C, e_ ¼ 1 104 =s, (d) higher magnification view of ‘‘A’’ in (c), (e) higher magnification view of ‘‘B’’ in (c).
v
inhomogeniety of deformation during LCF by solute
locking of moving dislocations so that DSA enhanced
the partitioning of cyclic strains into segregate regions.
This localized deformation leads to multiple crack
initiation sites as shown in Fig. 9(a). In the study on
316L(N) stainless steel [5], where a total strain amplitude of 0.6% was employed for LCF in air, crack
initiation is transgranular and occurs from slip bands
connected to the surface. Crack propagation is transgranular at higher strain rates (> 3 104 =s) in the
v
temperature range of 550–600 C, but at low strain
rates (< 3 104 =s), intergranular cracks are noticed
v
on the fracture surface. Often at 600 C at these strain
rates, crack propagation is assisted by oxidation. As
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S.-G. Hong, S.-B. Lee / International Journal of Fatigue 26 (2004) 899–910
shown in Fig. 9(d) and (e), the oxidation enhanced
v
crack propagation is observed at 650 C and strain rate
of 1 104 =s, but the crack propagation mode is transgranular with fatigue striations.
4. Conclusions
Based on the results and analyses on tensile and lowcycle fatigue tests, the following conclusions are made.
1. An increase in material strength and decrease in
ductility were induced by dynamic strain aging in
v
the temperature range from 250 to 600 C, and was
consistent with the regime of dynamic strain aging
manifested by a plateau in the variation of dynamic
strain hardening stress with temperature and a negative strain rate sensitivity. However, it shows some
difference with the regime of occurrence of the serrated flow in stress–strain curves.
2. In the regime of dynamic strain aging, the fibrous
zone size reduced, compared with that in other temperatures, and it is consistent with ductility
reduction in the DSA regime.
3. The negative SRS regimes (DSA regimes) between
tensile and LCF loading conditions coincide with
each other in the temperature range from 400 to
v
600 C, and SRS values under the two loading conditions are similar.
4. Softening ratio (the amount of cyclic softening during LCF) decreased in the DSA regime, compared
with that in other temperatures, and it was due to
the hardening by DSA.
5. A steep reduction in fatigue resistance with decreasing strain rate in the DSA regime, compared with
that in other temperatures, is attributed to the
enhanced crack initiation and propagation due to the
DSA effect.
6. Dynamic strain aging enhanced crack initiation,
causing multiple crack initiation sites, and the crack
propagation mode was transgranular with fatigue
v
striations at all test conditions (400–650 C and
1 104 1 102 =s).
Acknowledgements
This work has been supported by Computer Aided
Reliability Evaluation (CARE) National Research
Laboratory in Korea Advanced Institute of Science
and Technology (KAIST).
References
[1] Kanazawa K, Yamaguchi K, Nishijima S. Mapping of low cycle
fatigue mechanisms at elevated temperatures for an austenitic
stainless steel. ASTM STP 942, 1988. p. 519–30.
[2] Bressers J. In: Marriott JB, Merz M, Nihoul J, Ward J, editors.
High temperature alloys, their exploitable potential. Amsterdam:
Elsevier Applied Science; 1987. p. 385–410.
[3] Valsan M, Sastry DH, Bhanu Sankara Rao K, Mannan SL.
Effect of strain rate on the high-temperature low-cycle fatigue
properties of a nimonic PE-16 superalloy. Metall Trans 1995;
25A:159–71.
[4] Srinivasan VS, Sandhya R, Bhanu Sankara Rao K, Mannan SL,
Raghavan KS. Effect of temperature on the low cycle fatigue
behavior of nitrogen alloyed type 316L stainless steel. Int J Fatigue 1991;13(6):471–8.
[5] Srinivasan VS, Valsan M, Sandhya R, Bhanu Sankara Rao K,
Mannan SL, Sastry DH. High temperature time-dependent low
cycle fatigue behavior of a type 316L(N) stainless steel. Int J
Fatigue 1999;21:11–21.
[6] Rodriguez P. Encyclopedia of materials science and engineering,
vol. 1 (5). NY: Pergamon Press; 1988. p. 504–8.
[7] Mannan SL, Samuel KG, Rodriguez P. Trans Ind Inst Metals
1983;36:313.
[8] Kim DJ. The effects of dynamic strain aging on the low-cycle
fatigue behaviors at high temperature for the AISI 304L stainless
steel. Ph.D. thesis, Department of Material Science, KAIST,
1988.
[9] Abdel-Raouf H, Plumtree A, Topper TH. Effects of temperature
and deformation rate on cyclic strength and fracture of low carbon steel. ASTM STP 519, 1973. p. 28–57.
[10] Challenger KD, Miller AK, Brinkman CR. An explanation for
the effects of hold periods on the elevated temperature fatigue
behavior of 2 1/4Cr–1Mo steel. Trans ASME J Eng Mater Tech
1981;103:7–14.
[11] Hong SG, Lee SB. Development of a new LCF life prediction
model of 316L stainless steel at elevated temperature. Trans
KSME A 2002;26(3):521–7.
[12] van den Beukel A. On the mechanism of serrated yielding and
dynamic strain aging. Acta Metall 1980;28:965–9.
[13] Bhanu Sankara Rao K, Valsan M, Sandhya R, Mannan SL,
Rodriguez P. An assessment of cold work effects on straincontrolled low cycle fatigue behavior of type 304 stainless steel.
Metall Trans 1993;24A:913–24.
[14] Plumbridge WJ, Dalski ME, Castle PJ. High strain fatigue of a
type 316 stainless steel. Fract Eng Mater Struct 1980;3:177–88.
[15] Ganesh Sundara Raman S, Padmanabhan KA. Effect of prior
cold work on the room-temperature low cycle fatigue behavior
of AISI 304LN stainless steel. Int J Fatigue 1996;18(2):71–9.
[16] Sherman AM. Fatigue properties of high strength-low alloy
steels. Metall Trans 1975;6A:1035–40.
[17] Morrow J. ASTM STP 378, 1964. p. 45–84.
[18] Hong SG, Samson Y, Lee SB. The effect of temperature on lowcycle fatigue behavior of prior cold worked 316L stainless steel.
Int J Fatigue 2003;25(9-11):1293–300.
[19] Driver JH, Gorier C, Belrami C, Vidan P, Amzallag C. Influence
of temperature and environment on the fatigue mechanisms
of single-crystal and polycrystal 316L. ASTM STP 942, 1988.
p. 438–55.