Accepted Manuscript
Dislocation configurations through austenite grain misorientations
Arpan Das
PII: S0142-1123(14)00175-3DOI: http://dx.doi.org/10.1016/j.ijfatigue.2014.06.012Reference: JIJF 3402
To appear in: International Journal of Fatigue
Received Date: 10 March 2014Revised Date: 22 June 2014Accepted Date: 23 June 2014
Please cite this article as: Das, A., Dislocation configurations through austenite grain misorientations, InternationalJournal of Fatigue (2014), doi: http://dx.doi.org/10.1016/j.ijfatigue.2014.06.012
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Dislocation configurations through
austenite grain misorientations
Arpan Das
Fatigue & Fracture Group, Materials Science & Technology Division
CSIR - National Metallurgical Laboratory
(Council of Scientific & Industrial Research), Jamshedpur 831 007, India
E-Mail: [email protected], Tel.: +91-(0)657 2345192, Fax: +91-(0)657 2345213
Abstract
In the present investigation, many strain controlled low cycle fatigue experiments of
austenitic stainless steel were carried out at various total strain amplitudes under
ambient temperature where the strain rate was kept constant. Dislocation cell
developed due to strain cycling was measured through extensive analytical
transmission electron microscopic investigation and the deformed austenite grains'
misorientation was measured through extensive electron back scattered diffraction
experiments. A strong connection has been established with the dislocation
substructures' configurations, the deformed austenite grains' misorientation and the
extents of induced phase transformation occurs while cyclic plastic deformation of
metastable austenite at various total strain amplitudes. It has been investigated that
with the increase in strain amplitude, dislocation cells are getting more uniform. It has
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also been found that with the increase in strain amplitude, dislocation cell size
decreases drastically towards the higher strain amplitudes.
Keywords
Misorientation, Dislocation substructures, Low cycle fatigue, Deformation induced
martensite, Austenitic stainless steel.
Introduction
Cyclic plastic deformation behaviour with the corresponding microstructural
characteristics and morphologies of nuclear grade AISI 304LN austenitic stainless steel
at various total strain amplitude has already been explored and reported by the present
author elsewhere [1, 2]. On the basis of deformation characteristics of this present alloy
and the corresponding experimental evidences, it is reasonable to consider dislocation
reactions with the grain boundaries and the austenite grain's misorientation as one of
the controlling micro--mechanisms of cyclic hardening-softening behaviour. At the
position of peak tensile stress (Figure 1 in [1]), the mobile dislocations pile-up against
the cell wall (shown in Figure 4 in [2]) and on the other hand, when the compressive
stress is applied externally, the dislocation moves in the reverse manner and gets
accumulated in the cell wall on the opposite direction. In this circumstance, the
dislocations are arrested at the cell wall for a very short period. The progressive
accumulation of dislocation pile-up during strain cycling and the corresponding
increase in the resistance of grain boundaries to the dislocation pile-up might be
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another source of initial hardening in addition to the rapid increase in the dislocation
density.
It has been well established that most of the austenitic stainless steels are unstable
upon deformation and transform in to deformation induced martensites (DIM) (i.e., ε
(hcp) and/or to α/ (bcc)) [2-7]. Present author has already shown the formation,
nucleation micro-mechanisms and the martensitic transformation sequences, i.e., γ (fcc)
→ ε (hcp), γ (fcc) → ε (hcp) → α/ (bcc) and γ (fcc) → α/ (bcc) after both tensile and
fatigue deformation of AISI 304LN stainless steel at various loading conditions under
ambient temperature. It was also found that DIM can nucleate at many microstructural
locations, e.g., shear-band intersection, isolated shear-band, shear band-grain boundary
intersection, grain boundary triple points etc. [2-7]. The magnetic response of the
cyclically deformed austenite at various strain amplitudes has also been documented in
literature [5]. Such solid state phase transformation during plastic deformation imparts
a good combination of strength and toughness to the metastable austenitic stainless
steels.
It has already been investigated that cyclic deformation alters the microstructure and
causes the instabilities which influence the cyclic plastic deformation with different
strain amplitudes imposed [1, 2, 5]. Fatigue life strongly depended on the extent of
martensite present in the matrix, strain amplitude imposed to the alloy and the initial
grain size of the material. Fatigue properties deteriorated on martensite formation,
owing to more crack initiating sites becoming available [5]. It has already been
investigated by the present author that the crack density measured on the fatigue
fracture surfaces increases drastically with increasing strain amplitude and there is a
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strong connection of it with the extent of DIM while low cycle fatigue (LCF) deformation
of metastable austenite [5].
It has also been reported by the present author that austenite grains are more
responsible for determining the cyclic plastic behaviour of the present alloy and with
the increase in strain amplitude, the activity of the single slip system decreases and that
of the multiple slip system increases [8]. The connection between the strain amplitude
dependencies of the organisation of dislocation cell substructures with the formation of
α/ (bcc) martensite during cyclic plastic deformation of austenite has already been
investigated by the present author elsewhere [1].
In the present investigation, an attempt has been made to correlate the dislocation
substructural configurations (i.e., shape, size etc.) with the corresponding deformed
austenite grains' misorientation data after LCF deformation of austenitic stainless steel
tested at various total strain amplitude under room temperature. With regard to the
cyclic softening characteristics of the alloy, the development of dislocation cell sub-
structures and the role of grain boundaries need to be taken into account. As mentioned
above, a high density of dislocations has been preferentially formed in the area adjacent
to the grain boundaries at the very early stage of cyclic deformation. Dislocation pile-up
against the austenite grain boundary has been frequently observed when the specimens
are cyclically exhausted for more number of cycles (i.e., at the lower strain amplitudes).
Evidently, interactions between dislocations and grain boundaries played an important
role in the course of strain cycling.
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Experiments
In order to test the hypothesis, a large number of LCF experiments at various total
strain amplitudes (∆εt) ranging from ±0.50 to ±1.60% have been done on smooth
cylindrical solid specimens of 14.00 mm gauge length and 7.00 mm gauge diameter.
Nuclear grade AISI 304LN austenitic stainless steel was used for the present
investigation. Chemical composition (wt. %) of the material is: C 0.03, Mn 1.78, Si 0.65, S
0.02, P 0.034, Ni 8.17, Cr 18.73, Mo 0.26, Cu 0.29, N 0.08 and the balance Fe. The Md30
temperature of the material at which 50% of the austenite transforms to martensite at a
true strain of 0.30, calculated from the equation of Angel [9], is found to be 2.8°C. Initial
austenite grain size of the material was found to be 70 µm.
Tests were conducted under ambient condition until complete fracture of the specimens
in a servo-electric Instron machine of 100 kN load capacity (8862) in laboratory air
environment. A triangular waveform with a constant strain rate of 0.01 s-1 was used for
cyclic straining. An axial extensometer of 12.5 mm gauge length was kept attached to
the specimen surface for controlling the test parameters. For the quantification of α/
(bcc) martensite formed in the gauge section of fatigued specimens during strain
cycling, magnetic measurement technique was employed. The as-received condition did
not indicate the presence of martensite in the alloy.
Specimens for electron back scattered diffraction (EBSD) experiments have been
obtained by the transverse slicing of the LCF failed specimens leaving 2-3 mm distance
from the fracture end so as to avoid the regions of excessive plastic deformation. EBSD
technique has been extensively employed for the measurement of the grain boundary
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characteristics for all the specimens. The parameters that are kept constant for all the
EBSD experiments are: scan area (350 µm x 200 µm), magnification (500 X) and step
size (0.5 µm). Several frames (6-9) were scanned for each specimen.
Analytical Transmission Electron Microscopy (TEM) investigation has also been carried
out for all the fatigue failed specimens in order to understand the evolution of phase
transformation characteristics and mechanisms, formation of dislocation sub-structures
and its connectivity with the austenite grain boundary misorientation under two beam
conditions at an operating voltage of 200 kV. The TEM foils for all the samples were cut
approximately 2.5 mm away from the fatigued fracture surface so as to avoid excessive
plastic deformation. There was no dislocation cell present in the as received
microstructure. Dislocation cell size has been measured on several bright field images
by standard linear intercept method for all the fatigue fractured specimens.
Results and Discussion
Figure 1 explains that with the increase in strain amplitude, number of cycles to failure
decreases abruptly as expected but on the other hand, the peak tensile stress increases
drastically. The two dimensional grain boundary connectivity (i.e., austenite grain
boundary triple point density) decreases with the increase in strain amplitude up to
±1.0% but beyond that it increases drastically. Dislocation cell size decreases drastically
with the increase in strain amplitude from ±0.85 to ±1.40%. Dislocation cell size data
indicates considerable variability in the dislocation cell size from one region to another.
The dislocation evolution is inhomogeneous, which is due to the grain boundary effect.
However, this is also reflective of the fact that all the bright field images of dislocation
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cells have not been taken at exactly the same conditions of material-beam orientation.
Recently Mayama et al. [10] investigated the strain amplitude dependent organisation
of dislocation substructures in cyclically deformed 316L stainless steel in their elegant
study. Dislocations are developed step-by-step at lower stain amplitudes and easily to
create cells substructures due to the multiple slip system activity at higher strain
amplitudes [8].
Figure 2 (a) shows the as received microstructure of 304LN austenitic stainless steel
showing a large number of annealing twins spread throughout and polygonal austenite
grains. Microstructural investigations have been performed on fatigue failed specimens,
revealing inhomogeneous distribution of α/ (bcc) martensite shown in Figure 2 (b),
which is mainly attributed to the local stress distribution and hence crystallographic
variant selection of martensite. According to Ye et al. [11], the increase in slip band
density and activation of more number of slip systems during cycling have also been
well documented to bear a direct correlation with strain amplitude dependent cyclic
hardening behaviour.
Figure 3 shows the extent of α/ (bcc) martensite as a function of imposed strain
amplitude. It has been seen from Figure 3 that with the increase in strain amplitude, α/
(bcc) martensite content increases drastically. Even though there is scatter in the data
points, the trend is increasing in nature. The scatter is primarily due to the fact that α/
(bcc) martensite does not form uniformly throughout the gauge length of the fatigued
specimen.
Austenite to DIM transformation is believed to be triggered when the susceptible
austenitic stainless steels are deformed at temperatures below Md30. Present author
reported that a number of factors, e.g., steel chemistry, stress state, stress, strain, strain
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rate, grain size, initial crystallographic microtexture, and temperature of deformation
influence the formation of DIM [12]. The present author also demonstrated that a large
amount of published data relating the fraction of DIM to plastic strain, in fact, can be
described in terms of the pure thermodynamic effect of applied stress [13].
The formation of α/ (bcc) martensite during cyclic loading strongly affects the cyclic
mechanical behaviour of various types of steels. Figure 4 shows the misorientation
profile of the undeformed (as received material) as well as the deformed austenite
grains at various total strain amplitude. It is to be noted that there are two peaks of
maximum relative frequency present for all the specimens (i.e., approximately at 1.5
and 59.5°). In between, there is no such peak observed. There is a systematic variation
of the relative frequency of the austenite grain boundary misorientations with the
imposed strain amplitude.
Figure 5 shows that with the increase in total strain amplitude, the extent of medium
angle grain boundary (i.e., MAGB = 10 - 30°) and high angle grain boundary (i.e., HAGB ≥
30°) fraction decreases drastically, but on the other hand, low angle grain boundary (i.e.,
LAGB ≤ 10°) fraction increases drastically. Dislocation cell size is also varying in the
same manner as with MAGB and HAGB fraction with the strain amplitude range of ±0.85
to ±1.40%. It is mainly attributed that the MAGB and HAGB help dislocation cell to
reduce drastically in the strain amplitude range from ±0.85 to ±1.40%. It is understood
from the figure that more the fraction of LAGB, the cell size is less and vice-versa.
Dislocation structures are found near the grain boundaries, and the dislocation
structure is low energy walls or cells inside the grain.
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According to Watanabe et al. [14], HAGB can slide more easily than low sigma
coincidence boundaries, as observed for Σ3 and Σ13 coincidence boundaries, which
showed significant slide hardening in aluminium. In general, the dissociation of lattice
dislocation at coincidence boundaries becomes easier with increasing the Σ value and
the easiest at high-energy random boundaries (conventionally with Σ value larger than
29). In other words, random boundaries can more easily absorb lattice dislocations and
produce sliding than low Σ coincidence boundaries and the LAGB.
According to Humphreys et al. [15], during recovery, the stored energy of the material is
lowered by dislocation movement and this recovery process is not a single
microstructural variation but a series of events. The observed dislocation configurations
are strongly dependent on the applied strain amplitude, as well as on the stacking fault
energy (SFE) of the material [16, 17]. Different morphologies of dislocation sub-
structures (i.e., cells, tangled dislocation, wall and channels) formed in the cyclically
deformed alloy at various total strain amplitudes which have been discussed by the
present author reported elsewhere [1, 2]. According to Kuhlmann-Wildsdorf et al. [18],
the dislocation cell substructure is the last transformation of the dislocation
arrangements generated in fatigue.
It has been clearly understood that the size, shape and morphologies of dislocation cells
are different with the strain amplitude imposed. This is mainly attributed to the
variation of strain accumulation during cyclic plastic deformation at various strain
amplitudes. According to Huang et al. [19], higher strain amplitude will induce multiple
slip systems to facilitate the formation of dislocation cell sub-structures during LCF
deformation. The poorly developed cells might be the result of dislocations from
different slip systems interacting and trapping each other at intersections. Ma and Laird
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[20] have also pointed out that the dislocation structure changes from cells to loop
patches when the amplitude of loading changes from high to low. Long back, Feltner and
Laird [21] demonstrated different types of dislocation arrangements as a function of the
number of cycles to failure and the slip character of the material for cyclically deformed
polycrystalline fcc alloys. The value of Nf depend on the loading amplitude. According to
Laird et al. [22], higher strain amplitude induces multiple slip systems to facilitate the
formation of dislocation cell. The details of slip system activity during cyclic plasticity of
deformed austenite have been reported by present author elsewhere [8].
The observed dislocation configurations (Figure 6) are strongly dependent on the
employed strain amplitude, plastic strain accumulation as well as on the SFE of the
metastable alloy. Figure 6 (a-d) represents different morphologies and configurations of
dislocation substructures formed in the as received as well as in the cyclically deformed
austenite at various total strain amplitude. From this figure, it has been clearly
understood that the size, shape, configurations and morphologies of dislocation cells are
different with the nominal variation of imposed strain amplitude. Majority of
dislocations are actually distributed in the cell walls which are frequently referred as
sub cell boundaries. This is mainly attributed to the strain accumulation during the
cyclic plastic deformation at different strain amplitude and the induced phase
transformation. The common type of heterogeneous dislocation distribution in the
three dimensional cell structures develops under multiple slip conditions. Figure 7
shows the variation of dislocation cell size as a function of equivalent stress experienced
by the material. As the strain amplitude increases, dislocation cell size decreases
drastically at higher strain amplitudes range from ±0.85 to ±1.40% for the present
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material. This is mainly attributed to the austenite grain boundary misorientations and
the solid state martensitic transformation of the present alloy.
It has already been reported by the present author that with the increase in strain
amplitude, DIM formation increases which has been reported elsewhere [1, 2].
According to Raman et al. [17], dislocation tangles form during primary hardening and
the cell sub-structures are formed during secondary hardening. Owing to the
cumulative strain, the dislocations arrange themselves into relatively more stable
configurations (i.e., cell). The amount of DIM is significant by the time the material
attains the peak tensile stress. Hence, DIM forms much earlier than the dislocation cell
formed. Beyond ±0.85% strain amplitude, the dislocation cell size is decreased, but the
accumulation of strain increases with increasing strain amplitude, which generates
more DIM. Higher fraction of DIM triggers dislocation cells to grow further. According to
Huang et al. [19], the dislocation structures evolve with a strain amplitude decrease
from high to low, the low energy structure of walls, labyrinth walls, cells and
misorientation cells, which were formed at the higher strain amplitude, transfer to
dislocation structure of scattering walls, loop patches and the cell structures.
According to Koneva et al. [26], saturation value of dislocation cell size in most of the
polycrystalline fcc metals and alloys deformed at room temperature is slightly larger
than 0.10 µm and typically lies in the range of 0.20-0.60 µm. In austenitic stainless
steels, various kind of dislocation sub-structures have been observed, depending on the
nominal variation of strain amplitude imposed and the plastic strain accumulation
inside the material.
Figure 7 represents a clear comparison of dislocation cell size as a function of
equivalent stress of different fcc metals under various loading conditions at different
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temperatures. According to Challenger and Moteff [23], at 650°C, dislocation cell size
remains almost constant at various stress levels, while at 816°C, it decreases with
increase in stress for 304 austenitic stainless steel. They also performed fatigue tests for
316 austenitic stainless steel with temperature variation (i.e., 430°C, 650°C and 816°C)
and measured the dislocation cells size. At 430°C, cell size decreases with applied stress;
at intermediate and at higher temperatures (i.e., 650°C and 816°C respectively),
dislocation cell size decreases with equivalent stress drastically. A distinct transition
from rough dislocation cells at 430°C and 650°C to regular sub grains at 816°C have
been investigated by them. They also demonstrated that the dislocation cell size is the
characteristics of a given temperature and strain rate, and it has a direct influence on
the strain hardening capabilities of that material.
Ei-Madhoun et al. [25] investigated that stress and strain response and development of
dislocation substructures in deformed polycrystalline aluminium revealing that
saturation stress is linearly related to the inverse of dislocation cell size. Zhang et al.
[24] carried out many cyclic plastic deformation experiments on polycrystalline copper
at ambient temperature under cyclic shear. They investigated that during the transient
stage of dislocation cell formation, the stress amplitude is inversely proportional to the
dislocation cell size. This correlation is identical to that between the saturated stress
magnitude and the corresponding dislocation cell size. This is in agreement with the
present investigation.
From Figure 7, it is to be noted that the vast majority of the data (i.e., cell size) falls well
within 1.0 µm and it has been investigated that the saturation value is approximately
0.40 µm. A decrease in dislocation cell size corresponds to an increase in the average
scalar dislocation density. An excess dislocation density is rapidly accumulated in cell
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walls, which results in cell misorientation. At the same time, the amount of dislocation
accumulation inside the cells remains low. With increase of strain amplitude, dislocation
density increases, cell size decreases, cell structure becomes better developed and cell
walls become more sharply defined. Deformation in any polycrystalline materials is
inherently inhomogeneous, which produces interaction stresses between grains and
dislocation structures on different slip systems within the grains, and these stresses can
promote the softening. According to Bassim et al. [27, 28], in the case of unidirectional
deformation, minimization of stored dislocation strain energy is the driving force, which
is indicated by the clearly much reduced dislocation density in the cell structure as
compared with the wall structure, and even more persuasively by the fact that the cells
nearly conform to the theoretically derived cell structure of minimum energy. With the
increase in strain, the cell wall width can decrease, increase or even remain constant in
different materials [21]. According to Huang et al. [19], the back force is larger than the
interior of the grain with a decrease in strain amplitude. According to them, the
dislocation development is faster at the grain boundaries and the twin boundaries due
to strain localization.
As strain amplitude increases, DIM fraction increases rapidly (Figure 3). Beyond the
strain amplitude of ±0.85%, dislocation cells are formed prominently. With the increase
in strain amplitude, dislocation cell size decreases abruptly (Figure 7). It has also been
found that the cell wall width increases drastically with the increase in strain amplitude.
According to Breedis [16], the deformation substructure in austenite influences the
subsequent nucleation of martensite on cooling to below the MS temperature. The
interaction of DIM with dislocation cells; shear bands with dislocation cells are shown at
different strain amplitude in Figure 8 (a-d). Figure 8 (a) shows the dislocation jungles at
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strain amplitude ±0.85%. Dislocation cells are observed at strain amplitude ±1.00%,
±1.20% and ±1.40% (Figure 8 (b-d)). Hence, the extent of DIM and their heterogeneous
distribution caused different dislocation cell configurations and characteristics. Hence,
there is a strong connection between DIM and dislocation cell configurations.
Conclusions
It appears that there is a systematic correlation of the grain boundary connectivity,
austenite grain misorientations and their types with the dislocation cell sub-structures'
configurations with the strain amplitude variation under cyclic plastic deformation of
metastable austenitic stainless steel. The present investigation also concludes that as
the strain amplitude increases, the dislocation cell sizes are getting more uniform in size
and there is an interaction between grain boundaries and DIM. It has also been found
that with the increase in strain amplitude, dislocation cell size decreases abruptly in the
range of higher strain amplitudes. It has also been investigated that with the increase in
strain amplitude, MAGB and HAGB fraction decreases drastically and on the other hand
LAGB fraction increases. DIM, grain boundary connectivity and their configurations are
the significant contributory factors in forming the dislocation cell substructures.
Acknowledgements
The author acknowledges the support and encouragement of Dr. S. Srikanth, Director,
CSIR-NML. The author also expresses his gratitude to Dr. S. Sivaprasad, Principal
Scientist for his assistance with the mechanical characterization. I also express my
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gratitude to Dr. M. Ghosh, Senior Scientist of CSIR--NML, for helping in carrying out TEM
characterisation.
References
[1] Das A, Sivaprasad S, Chakraborti PC, Tarafder S. Connection between deformation-
induced dislocation substructures with martensite formation. Phil Mag Let 2011;
91(10):664-5.
[2] Das A, Sivaprasad S, Chakraborti PC, Tarafder S. Morphologies and characteristics
of deformation induced martensite during low cycle fatigue behaviour of
austenitic stainless steel. Mat Sci and Engg A 2011; 528:7909-14.
[3] Das A, Sivaprasad S, Ghosh M, Chakraborti PC, Tarafder S. Morphologies and
characteristics of deformation induced martensite during tensile deformation of
304 LN stainless steel. Mat Sci and Engg A 2008; 486(1--2):283--286.
[4] Das A, Tarafder S. Experimental investigation on martensitic transformation and
fracture morphologies of austenitic stainless steel. Int J of Plast 2009;
25(11):2222--2247.
[5] Das A. Magnetic properties of cyclically deformed austenite. J of Mag and Mag Mat
2014; 361:232-242.
[6] Das A, Sivaprasad S, Ghosh M, Chakraborti PC, Tarafder S. Correspondence of
fracture surface features with mechanical properties in 304LN stainless steel. Mat
Sci and Engg A 2008; 496:98-105.
[7] Das A. Martensite - Void Interaction. Scripta Mat 2013; 68(7):514-517.
[8] Das A. Slip system activity during cyclic plasticity. Met and Mat Trans A 2014;
45(7):2927-2930.
16
[9] Angel T. Formation of martensite in austenitic stainless steels. J of the Iron and
Steel Inst 1954; 177(5):165-74.
[10] Mayama T, Sasaki K, Kuroda M. Quantitative evaluations for strain amplitude
dependent organization of dislocation structures due to cyclic plasticity in
austenitic stainless steel 316L. Acta Mater 2008; 56(12):2735-2743.
[11] Ye D, Xu Y, Xiao L, Cha H. Effects of low-cycle fatigue on static mechanical
properties, microstructures and fracture behaviour of 304 stainless steel. Mat Sci
and Engg A 2010; 527(16-17):4092-102.
[12] Das A, Tarafder S, Chakraborti PC. Estimation of deformation induced martensite
in austenitic stainless steels. Mat Sci and Engg A 2011; 529:9--20.
[13] Das A, Chakraborti PC, Tarafder S, Bhadeshia HKDH. Analysis of deformation
induced martensitic transformation in stainless steel. Mat Sci and Tech 2011;
27(1):366--370.
[14] Watanabe T, Tsurekawa S, Kobayashi S, Yamaura S. Structure dependent grain
boundary deformation and fracture at high temperatures. Mat Sci and Engg A
2005; 410-411:140-7.
[15] Humphreys JF, Hatherly M. Recrystallization and Related Annealing Phenomena,
Pergamon Press 1995.
[16] Breedis JF. Influence of dislocation substructure on the martensitic transformation
in stainless steel. Acta Met 1965; 13(3):239-50.
[17] Raman SGS, Padmanabhan KA. Influence of martensite formation and grain size on
the room temperature low cycle fatigue behaviour of AISI 304LN austenitic
stainless steel. Mat Sci and Tech 1994; 10:614-20.
[18] Kuhlmann-Wilsdorf D, Laird C. Dislocation behaviour in fatigue. Mat Sci and Engg
1977; 27:137-156.
17
[19] Huang HL. A study of dislocation evolution in polycrystalline copper during low
cycle fatigue at low strain amplitudes. Mat Sci and Engg A 2003; 342:38-43.
[20] Ma BT, Laird C. Dislocation structures of copper single crystals for fatigue tests
under variable amplitudes. Mat Sci and Engg A 1988; 102:247-58.
[21] Feltner CE, Laird C. Factors influencing the dislocation structures in fatigued
metals. Trans of TMS-AIME 1968; 242:1253-57.
[22] Laird C, Charsley P, Mughrabi H. Low energy dislocation structure produced by
cyclic deformation. Mat Sci and Engg 1986; 81:433-450.
[23] Challenger KD, Motef J. A correlation between strain hardening parameters and
dislocation substructure in austenitic stainless steels. Scripta Mat 1972; 6:155.
[24] Zhang J, Jiang Y. An experimental study of the formation of typical dislocation
patterns in polycrystalline copper under cyclic shear. Int J of Plast 2007; 55:1831-
42.
[25] Ei-Madhoun Y, Mohamed A, Bassim MN. Cyclic stress-strain response and
dislocation structures in polycrystalline aluminum. Mat Sci and Engg A 2003;
359:220-7.
[26] Koneva NA, Starenchenko VA, Lychagin DV, Trishkina LI, Popava NA, Kozlov EV.
Formation of dislocation cell structure in face centred cubic metallic solid
solutions. Mat Sci and Engg A 2008; 483-484:179-83.
[27] Bassim MN, Kuhlmann-Wilsdorf D. Stresses of hexagonal screw dislocation arrays.
IV. Cell aggregates. Phys Status Solidi A 1973; 17(2):379-393.
[28] Bassim MN, Kuhlmann-Wilsdorf D. Stresses of hexagonal screw dislocation arrays.
V. Short range stresses and their contribution to latent hardening. Phys Status
Solidi A 1973; 19(1):335-346.
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List of Figures
Figure 1: Peak tensile stress, number of cycles to failure, grain boundary triple point
(GBTP) density and dislocation cell size as a function of strain amplitude imposed
during LCF deformation of metastable AISI 304LN stainless steel at room temperature.
AR - As Received.
Figure 2: (a) As received microstructure showing polygonal austenite grains and a large
number of annealing twins and (b) Cyclically deformed austenite showing the extent of
α/ (bcc) martensite formed in it and the dense shear bands.
Figure 3: The extent of α/ (bcc) martensite as a function of strain amplitude -
ferritescope measurement [1, 2].
Figure 4: Relative frequencies of austenite grain boundary misorientation profile during
LCF deformation of AISI 304LN stainless steel at various strain amplitude tested at
room temperature. In between there is no sharp peak.
Figure 5: Extents of undeformed and deformed austenite grain boundaries (i.e., LAGB,
MAGB and HAGB) fraction and dislocation cell size as a function of strain amplitude
imposed during LCF deformation of AISI 304LN stainless steel at room temperature. AR
- As Received.
Figure 6: Dislocation substructures observed by TEM bright field images: (a) as received
condition, (b) strain amplitude = ±0.85%, (c) strain amplitude = ±1.20% and (d) strain
amplitude = ±1.40%.
19
Figure 7: Dislocation cell size as a function of equivalent stress magnitude of some fcc
materials under various loading conditions at different temperatures. A - AISI 304 SS -
LCF at 650°C, B - AISI 304 SS - LCF at 816°C, C - AISI 316 SS - LCF at 430°C, D - AISI 316
SS - LCF at 650°C, E - AISI 316 SS - LCF at 816°C, F - Polycrystalline copper (textured
with grain size of 75 µm) - cyclic torsion at ambient, G - Polycrystalline copper (textured
with grain size of 75 µm) - tension-compression at ambient, H - Polycrystalline copper
(textured with grain size of 75 µm) - 90° out of phase non-proportional loading at
ambient, I - Polycrystalline copper (textured with grain size of 75 µm) - transient stage
at ambient, J - Polycrystalline copper (textured with grain size of 10 µm) - tension-
compression at ambient, K - Polycrystalline copper (textured with grain size of 10 µm) -
cyclic torsion at ambient, L - Polycrystalline copper (textured free with grain size of 75
µm) - tension--compression at ambient, M - Polycrystalline aluminium - LCF at ambient
and N - Present study. A - E (Challenger and Moteff [23]); F - L (Zhang and Jiang [24]); M
(Ei-Madhoun et al. [25]) [1].
Figure 8: Dislocation substructures and martensite/shear band interactions at strain
amplitude of (a) ±0.85%, (b) ±1.00%, (c) ±1.20% and (d) ±1.40%. A strong connection
of DIM with dislocation substructures is noted.
MATERIALS SCIENCE & TECHNOLOGY DIVISION
CSIR - National Metallurgical Laboratory
(Council of Scientific & Industrial Research)
Jamshedpur 831 007, India
Highlights
1. LCF experiments were done at various strain amplitudes under ambient.
2. Metastable austenitic stainless steel was chosen.
3. Dislocation cell size developed due to strain cycling was measured by TEM.
4. Austenite grain misorientation was measured by EBSD.
5. Dislocation cell size, austenite grains' misorientation and extent of DIM were
correlated.
Graphical Abstract
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