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Materials Science and Engineering A 479 (2008) 229–235
Effect of cryogenic treatment on distribution of residual stress
in case carburized En 353 steel
A. Bensely a,∗ , S. Venkatesh a , D. Mohan Lal a , G. Nagarajan a ,
A. Rajadurai b , Krzysztof Junik c
a Department of Mechanical Engineering, College of Engineering, Anna University, Guindy, Chennai 600025, India
Department of Production Engineering, Madras Institute of Technology, Anna University, Chromepet, Chennai 600044, India
c Institute of Materials Science and Applied Mechanics, Department of Mechanics, Wroclaw University of Technology, Poland
b
Received 6 July 2006; received in revised form 14 June 2007; accepted 16 July 2007
Abstract
The effect of cryogenic treatment on the distribution of residual stress in the case carburized steel (En 353) was studied using X-ray diffraction
technique. Two types of cryogenic treatment: shallow cryogenic treatment (193 K) and deep cryogenic treatment (77 K) were adopted, as a
supplement to conventional heat treatment. The amount of retained austenite in conventionally heat-treated, shallow cryogenically treated and
deep cryogenically treated samples was found to be 28%, 22% and 14%, respectively. The conventionally heat-treated, shallow cryogenically
treated and deep cryogenically treated samples in untempered condition had a surface residual stress of −125 MPa, −115 MPa and −235 MPa,
respectively. After tempering the conventionally heat-treated, shallow cryogenically treated and deep cryogenically treated samples had a surface
residual stress of −150 MPa, −80 MPa and −80 MPa, respectively. A comparative study of the three treatments revealed that there was an increase
in the compressive residual stress in steel that was subjected to cryogenic treatment prior to tempering. The experimental investigation revealed
that deep cryogenically treated steel when subjected to tempering has undergone a reduction in compressive residual stress. Such stress relieving
behaviour was mainly due to the increased precipitation of fine carbides in specimens subjected to DCT with tempering.
© 2007 Elsevier B.V. All rights reserved.
Keywords: Case carburized steel; Cryogenic treatment; Grinding; Residual stress; Retained austenite
1. Introduction
Scientists and engineers have been successful in improving
the impact and fatigue properties of metals by intentionally
introducing residual stresses into the surface of these materials. Residual stress exists in an elastic solid body in the absence
of, or in addition to, the stresses caused by virtue of load or
temperature or both. Such stresses can arise from deformation
during cold working, in welding from weld metal shrinkage,
and in changes in volume due to thermal expansion. In recent
years, there has been a considerable interest in the properties and
development of compressive residual stress. Knowledge of the
nature of residual stresses is of paramount importance since they
can lead to creep failure, fatigue and stress corrosion cracking in
sensitive materials. Mack Alder and Olsson [1] have observed
∗
Corresponding author. Tel.: +91 44 22203262.
E-mail address: benzlee5@yahoo.com (A. Bensely).
0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved.
doi:10.1016/j.msea.2007.07.035
that after case hardening, the gear tooth would experience compressive residual stresses, which are beneficial in maintaining
an appreciable endurance limit. But these beneficial stresses are
counter acted by detrimental tensile residual stresses in the core.
The explanation of this phenomenon is based on the uneven volumetric expansion in the core and at the surface layer during
phase transformation. Hence, it is an important consideration in
the heat treatment of steel.
Various methods are used to improve the behaviour of steel.
Carburizing is one of the prevalent methods used for this purpose. It is a thermo chemical diffusion process used to produce a
hard wear resistant case. Due to carburization, the percentage of
carbon in the case will increase to the carbon potential applied.
Carburizing is beneficial to the component but the increase in
carbon content leads to retention of austenite after hardening,
which is not desirable. In some steels, those with higher carbon
content and alloy steels, the martensite finish temperature (Mf ) is
below 273 K, which means that at the end of the heat treatment,
there is as much as 5–15% of austenite remaining [2].
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A. Bensely et al. / Materials Science and Engineering A 479 (2008) 229–235
The amount of retained austenite exhibits significant effects
on the magnitude of compressive stresses formed and ultimately,
on dimensional stability. Some of the factors affecting retained
austenite formation include chemical composition, quenching
temperature, quenching cooling rates, austenitizing temperature,
grain size and tempering. The effects of retained austenite on the
properties of steel includes decrease of tensile strength and yield
strength, reduction of maximum attainable surface compressive
stresses and this is proportional to the amount of retained austenite present. Also, strength and compressive residual stresses are
reduced by the presence of retained austenite, thereby reducing
fatigue resistance [3]. Tempering process reduces the percentage of retained austenite. But it may not always be the effective
way to reduce the amount of retained austenite, because of the
marked decrease of hardness, mechanical strength, and wear
resistance that partially cancels the positive effect of retained
austenite transformation.
For this reasons a more effective method could be cryogenic
treatment. The transformation of austenite to martensite would
influence the residual stress, which in turn would affect the
behaviour of the material. The material under investigation is
En 353, which finds applications in heavy-duty gears, shafts,
pinions, rocker arms, camshafts, gudgeon pins, push rod leads,
levers, bushes and other small arm components. The objective
of the present study is to examine the distribution of residual
stress in En 353 due to cryogenic treatment.
2. Standpoints in research
A lot of research has been going on over the past few decades,
in the field of cryogenics. Ioan Alexandru and Vasile Bulancea
[4] have concluded that the residual stresses play a very important role both during and after the cryogenic cooling. These
stresses constitute the main cause for the continuation of the
transformation of retained austenite to martensite, within the
cryogenic field. It is surprising to see the fact that after cryogenic
cooling the value of the residual stresses becomes even lower
compared with classical quenching, either in water or in oil, to
the ambient temperature. When quenched steel is cryogenically
treated, the retained austenite will transform to martensite. Then
the size of the component will have only a little expansion and
the stability of the component will increase. Also, it has been
observed that the structure of cryogenically cooled materials
proved to be much more uniform and dense. In addition, cryogenic cooling has induced the occurrence of very fine carbides
with dimensions less than 1 ␮m, which occupy the micro voids
and contribute to the increase of both density and coherence
with the metal. Molinari et al. mentioned that carbide precipitation occurs with a higher activation energy thus leading to
a higher nucleation rate and in turn to finer dimensions and a
more homogeneous distribution. A new phenomenon referred as
tempered martensite detwining was observed. Deep cryogenic
treatment can strongly reduce the wear rate of the hot work
tool steel. This result was interpreted on the basis of increased
toughness, because in the presence of delamination the ability
of materials to oppose crack propagation can really increase
the mechanical stability on the wear surface and the load bearing capacity. Therefore, even if the deep cryogenic treatment
does not influence the hardness, it increases both toughness
and wear resistance [5]. Barron conducted preliminary tests
to determine the effect of cryogenic treatment on lathe tools,
end mills, zone punches and concluded that an increase in tool
life from 50% to more than 200% was observed for the tools
which had been soaked in liquid nitrogen for 12 h [6]. Fanju
Meng et al. have studied the effect of cryogenic treatment on
the wear behaviour of Fe–12Cr–Mo–V–1.4C tool steel. The
results have shown a dramatic increase in wear resistance especially at high sliding speeds [7]. Microstructural analysis of the
sample after cryogenic treatment has shown fine carbide precipitates of size in the range of 10 nm, which are characterized
as ␩-carbide. This formation of ␩-carbides helps to improve
the wear resistance. Mohan Lal et al. [8] studied the effect of
cryogenic treatment on T1 type-high speed material and found
that soaking at 203 K can attain the maximum hardness of 67
HRC.
The effect of cryogenic treatment on the residual stress
behaviour of En 353 prior to tempering and after tempering is
not yet reported. Hence the present study deals with the measurement of residual stress distribution, which in turn helps to
know whether cryogenic treatment helps in improving the life
of the component.
3. Experimental
3.1. Sample preparation
Commercially available 30 mm diameter bar stock of En 353
raw material was procured. In order to confirm the composition of the material procured, elemental analysis was carried
out using optical emission spectroscope (OES), the results of
which are shown in Table 1. The raw sample conforms to the
chemical requirements of En 353 specifications, as far as the
constituents are concerned. It was also found that weight percentage of carbon has increased from 0.16 to 0.75 in the case
Table 1
Chemical composition of the untreated En 353 steel (weight %)
Sample description
Carbon
Silicon
Manganese Phosphorus Sulphur
Nominal composition
0.12–0.18
0.1–0.35
0.6–1.0
Raw material
0.16
0.21
0.68
Case composition after conventional heat treatment
Uncertainty
0.75
0.20
0.67
±0.011
±0.013
±0.013
≤0.035
0.022
Chromium
Nickel
Molybdenum
0.0025–0.05
0.75–1.25
1.0–1.5
0.08–0.15
0.028
0.76
1.19
0.10
0.021
0.022
±0.011
±0.002
0.76
1.19
0.09
±0.007
±0.006
±0.020
A. Bensely et al. / Materials Science and Engineering A 479 (2008) 229–235
after case carburizing, which shows that the carbon potential
employed for carburizing is 0.75. After confirming the composition the 30 mm diameter bar stock was machined down to
25 mm diameter rod and finally cut down into six equal pieces
of length 15.2 mm. En 353 is widely used in making crown
wheel and pinion, which undergoes carburizing treatment. This
component requires a case depth of 1 mm after final grinding
operation before it is put into application. In order to ensure a
similar case depth of 1 mm, a case depth of 1.2 mm was achieved
during carburizing process. This in turn can be made to reach
1 mm after the various treatments by performing final surface
grinding operation to remove the excess 0.2 mm before residual stress measurements. Thus six specimens each of 25 mm
diameter and 15.2 mm length were obtained along with a blank
specimen (25 mm diameter × 15.2 mm length) of En 353. They
were subjected to liquid carburizing at 1183 K until the blank
specimen reached a case depth of 1.2 mm. Confirmation was
made by cutting and examining the blank specimen under optical
microscope. Then the six samples were air-cooled, followed by
hardening at 1093 K for 30 min. All the six specimens were then
quenched from 1093 K in oil at 313 K. Immediately after quench
hardening one sample was segregated from the six and it was
designated as conventionally heat-treated untempered specimen
(CHTUT). One of the remaining five samples was subjected
to tempering at 423 K for 1.5 h and was designated as conventionally heat-treated tempered specimen (CHTT). Meanwhile
the remaining four samples were segregated into two groups of
two specimens each. Both the samples of one group, which had
already undergone quench hardening, were subjected to shallow cryogenic treatment. In shallow cryogenic treatment the two
samples were kept in a mechanical freezer at 193 K for 5 h. After
5 h both were removed from freezer and exposed to room temperature. One of the two shallow cryogenically treated samples
was kept as such and it was designated as shallow cryogenically treated untempered specimen (SCTUT), while the other
one was immediately tempered at 423 K for 1.5 h and it was
designated as shallow cryogenically treated tempered specimen
(SCTT). Similarly the two specimens of the other group, which
had already undergone quench hardening process, were immediately subjected to deep cryogenic treatment. In deep cryogenic
treatment the two specimens were cooled from room temperature to 77 K in 3 h, soaked at that temperature for 24 h and heated
back to room temperature in 6 h. These very low temperatures
were achieved using computer controls in a well-insulated treatment chamber with liquid nitrogen (LN2 ) as working fluid. After
reaching the room temperature, one of the samples was kept as
such and it was designated as deep cryogenically treated untempered specimen (DCTUT) and the other specimen was subjected
to tempering at 423 K at 1.5 h and it was designated as deep cryogenically treated tempered specimen (DCTT). After completing
the treatments all the six specimens were subjected to conventional surface grinding so that a case depth of 1 mm similar to
that in the actual component was attained. Surface grinding was
carried out at 2800 revolutions per minute with a feed of 50 ␮m
using a CUMI make green silicon carbide abrasive wheel having
13 mm thickness and 60-grit. Cutting oil was used as the coolant
during grinding.
231
3.2. X-ray diffraction method
X-ray diffraction (XRD) measurements are based on the
change in the interplanar spacing (strain) by virtue of load or
temperature or both. Diffraction effects are produced when a
beam of X-rays of specific wavelength passes through the threedimensional array of atoms, which constitutes the crystal [9].
Each atom scatters a fraction of the incident beams, and if the
required conditions are fulfilled then the scattered waves reinforce to give a diffracted beam. These conditions are governed
by the Bragg equation:
nλ = 2dhkl sin θ
(1)
where n is the integer denoting the order of diffraction, λ the
wavelength of the incident beam, dhkl the interplanar spacing of
the hkl planes, and θ is the angle of incidence on the hkl planes.
Metallurgical applications of XRD technique include identification and quantitative analysis of crystalline chemical compounds
(phase analysis and quantification, e.g. retained austenite determination), residual macro- and micro-stress analysis.
In the present study, X-ray diffraction measurements
were performed on XStress 3000 diffractometer (Stresstech
Oy/Finalnd). Xstress 3000 X-Ray analyzer is based on solidstate linear sensor technique (MOS, Dual 512 pixels) and
goniometer in modified Ψ geometry (symmetrical side inclination). XRD analysis was performed at room temperature using
´˚ radiation for residual stress measurements
Cr K␣ (λ = 2.2909 A)
and retained austenite content also was determined (by using
vanadium filters). Accuracy obtainable using a diffractometer is
in the region of 0.5% for the range above 1 to 2 volumes percentage of austenite. Samples of 25 mm diameter and 15 mm
height were used for the measurements. Volume fraction of
retained austenite was determined by X-ray phase analysis using
{2 0 0, 2 1 1} peaks of martensite and {2 0 0, 2 0 0} peaks of
austenite. The main objective of the work was to determine
residual stress distribution both prior to and after conventional
heat treatment, shallow cryogenic treatment and deep cryogenic
treatment conditions. Residual stresses can be divided into two
general categories, macro-stresses, where the strain is uniform
over relatively large distances and micro-stresses produced by
non-uniform strain over short distances, typically a few hundred
˚ Both types of stresses can be measured by X-ray diffraction
A.
techniques. In the current study, macro-stresses were measured.
The basis of stress measurement by X-ray diffraction is the accurate measurement of changes in interplanar d spacing caused by
the residual stress. When macro-stresses are present the lattice
plane spacing in the crystals (grains) change corresponding to the
residual stress and the elastic constants of the material. This produces a shift in the position of the corresponding diffraction, i.e. a
change in Bragg angle θ. The determination of interplanar spacing depends on the accurate measurement of the corresponding
θ.
The data were acquired in several Ψ angles (in range of
0–45◦ ) and the stresses were calculated by the classical sin2 Ψ
method. The peak position was calculated by the cross correlation method. The parameters selected for experimentation
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A. Bensely et al. / Materials Science and Engineering A 479 (2008) 229–235
Table 2
Experimental inputs for residual stress analysis
Miller indices hkl
Diffraction angle (◦ )
Angular resolution (◦ )
Poisson’s ratio
Exposure time (s)
Penetration depth (␮m)
Young’s Modulus (GPa)
211
156.4
0.029
0.3
4
4
200
are listed in Table 2. Due to the limited penetration (4 ␮m for
chromium radiation in steel) of X-ray only surface stresses could
be measured. Hence stresses at shallow points in the samples
were determined by repeating the measurements after removing
(electro polishing) layers of known thickness. When performing
layer removal for residual stress depth profiling it is important to consider any redistribution or relaxation in the residual
stress in the exposed surface. A generalized solution proposed by
Sikarskie based on the original solutions of Moore and Evans
was used [10,11]. For electro polishing, the electrolyte used,
consisted of 80 cm3 perchloric acid (70%), 700 cm3 ethanol,
100 cm3 butoxyethanol and 120 cm3 -distilled water. The jet polishing was performed at temperatures ranging from 253 K to
263 K at a voltage between 20 V and 30 V and at a current
of about 600–900 mA. The specimens were jet polished using
the electrolyte at 283 K. The peaks on the XRD patterns were
indexed with the X-ray polycrystalline powder diffraction files
(International Center for Diffraction Data (ICDD)).
4. Results and discussion
4.1. Retained austenite
The retained austenite present in the specimens of En 353
subjected to conventional heat treatment, shallow cryogenic
treatment and deep cryogenic treatment prior to and after tempering was measured using X-ray diffractometer with vanadium
filters. The results are shown in Table 3.
After conventional quench hardening of carburized specimens (i.e. CHTUT), it was found that there was 28.1% retained
austenite. This, on tempering (i.e. CHTT) does not yield any
reduction in retained austenite content because the temperature
employed for tempering is just 423 K, which cannot decompose the available austenite. The decomposition of austenite
occurs from 513 K to 593 K [12]. Hence there is no change
in retained austenite prior to and after tempering in case of
Table 3
Retained austenite content by X-ray measurements
Type of treatment
Retained austenite (%)
CHTUT
CHTT
SCTUT
SCTT
DCTUT
DCTT
28.1
28.5
22.0
22.8
14.9
14.3
±
±
±
±
±
±
3.5
6.1
7.6
5.9
5.8
4.1
conventional heat treatment. The conventional quench hardened
specimen (i.e. CHTUT) when kept in a mechanical freezer at
193 K has resulted in a reduction of retained austenite from
28.1% to 22% (i.e. SCTUT). As the temperature applied for tempering is not sufficient to decompose the austenite into ferrite
and cementite, 22.8% remains even after tempering (i.e. SCTT).
This clearly indicates that shallow cryogenic treatment does not
completely transform the retained austenite into martensite but
it has reduced the percentage of retained austenite available in
comparison with conventional heat treatment. When the conventionally quench hardened carburized specimens (i.e. CHTUT)
were subjected to deep cryogenic treatment, a 50% reduction
of available retained austenite (28.1% austenite in CHUT) was
observed in DCTUT specimens. Increasing carbon content of
steel increases the potential for retained austenite on quenching. This is because Ms temperature decreases with increasing
carbon content. The effect of carbon content can be calculated
using the Steven and Haynes equation for steel containing up to
0.5% carbon [13].
Ms (◦ C) = 561 − 474C − 33Mn − 17Cr
−17Ni − 21Mo
Mf = Ms − 215
(2)
(3)
where carbon, manganese, nickel, chromium and molybdenum
are concentrations of these elements in percent. For carbon concentration more than 0.5% corrections must be made. Since the
carbon concentration in the surface of En 353 after carburizing is 0.75% correction curve was used. It was found that 20 K
has to be added to the Ms temperatures derived using Steven
and Haynes equation. 100% martensite transformation can be
expected to happen at a temperature of 226 K for a carbon content of 0.75% in case. The present study clearly indicates that
100% removal of retained austenite was not obtained even after
cooling the samples to a very low temperature of 77 K. Thus,
some amount of retained austenite exists in deep cryogenically
treated samples and shallow cryogenically treated samples also.
4.2. Residual stress distribution
Fig. 1 shows the distribution of residual stress in samples
tested prior to tempering. A uniform difference in values of
residual stress is observed corresponding to the measured depth
from the surface. The gradient of residual stress is the same, irrespective of which treatment is employed but the range of values
varies with respect to the treatment. This can be attributed to
the transformation of retained austenite to martensite. A reduction of around 6% retained austenite from CHTUT samples
does not induce compressive residual stress in SCTUT samples, whereas a tremendous increase in compressive residual
stress is observed in DCTUT samples and is due to 50% transformation of available retained austenite (28.1% in CHTUT
samples) to martensite. The process of subjecting samples to
shallow cryogenic treatment is totally different from deep cryogenic treatment apart from the level of temperature reduction.
A. Bensely et al. / Materials Science and Engineering A 479 (2008) 229–235
Fig. 1. Distribution of residual stress in En 353 steel specimens prior to tempering.
In shallow cryogenic treatment the CHTUT samples available
at room temperature are suddenly exposed to 193 K, similarly it
is immediately taken out of the mechanical freezer and exposed
to room temperature. In the case of deep cryogenic treatment
process the CHTUT samples were slowly taken to 77 K from
room temperature and slowly brought back to room temperature. Hence whatever stress developed during the process has not
been relieved. This mechanism need to be validated by carrying
out a detailed investigation on the effects of lattice parameters
‘c’ and ‘a’ with respect to the rate of cooling to 193 K and 77 K
and rate of warming to room temperature. Lower the temperature of cryogenic treatment, greater will be the transformation of
austenite to martensite and hence, greater will be the compressive residual stress induced. This compressive residual stress is
beneficial with respect to wear and fatigue [3]. This was evident in the earlier research on wear resistance of cryogenically
treated En 353 steel [14].
Cryogenic treatment cannot represent a final heat treatment
since subsequent tempering is absolutely necessary to achieve
fine carbide precipitation. Tempering reduces residual stresses,
increases ductility, toughness and ensures dimensional stability.
Fig. 2 shows the distribution of residual stress in conventional
heat treatment, shallow cryogenic treatment and deep cryogenic
233
treatment after tempering. The untempered conventionally heat
treated and shallow cryogenically treated samples, which when
subjected to tempering (CHTT, SCTT) seemed to retain the
residual stress after tempering; while those subjected to deep
cryogenic treatment followed by tempering (DCTT) showed a
remarkable stress relief. This indicates that the extent to which
the metal is cooled below room temperature plays an important
role. The distribution trend is not uniform and the stress band is
also altered. The substantial relief of compressive residual stress
in the DCTT specimens may be due to the occurrence of finer
precipitates throughout the matrix and the loss of tetragonality of
martensite. The mechanism behind the large change in residual
stress can be due to atomic level change during the cool down,
soak and ramp up process of deep cryogenic treatment as well
as the immediate tempering process. The cool down process of
deep cryogenic treatment causes the transformation of retained
austenite to freshly formed martensite, which has different lattice parameter (higher c/a ratio) than the original martensite
[15]. Owing to volume contraction in the DCT process, the crystalline lattice tends to decrease. However, due to the difficulty
in diffusion of carbon atoms at low temperature, precipitation of
ultra fine carbides will not take place at a very low temperature
of 77 K. But increased holding time at cryogenic temperature
involves localized diffusion of carbon leading to cluster formation. These clusters act as nuclei for the formation of ultra fine
carbides on subsequent warm up and tempering. These changes
in lattice parameter are confirmed by a recent in situ neutron
diffraction study, which indicated that the lattice parameters ‘a’
and ‘c’ of the martensite behave differently during the cooling and warming-up processes [16]. The lattice parameter ‘a’
changes with the temperature almost linearly, following almost
the same curve during the cooling and warming-up process, indicating a pure thermal elastic effect. The lattice parameter c, on
the other hand, first decreases with the cooling temperature, but it
does not follow the same slope while warming-up, and increases
only very slightly during the warming-up process, indicating it is
not only a pure thermal effect. It is inferred from the above result
that carbon atoms segregation did occur during the cold treatment process. Because the carbon atoms predominantly occupy
the octahedral or tetrahedral sites in the martensitic lattice, the
segregation of carbon atoms from the octahedral or tetrahedral
site to the defect regions mainly affects the ‘c’ lattice parameter. The capacity of carbon atoms to diffuse increases as the
temperature rises back to room temperature. During this the
carbon atoms move a short distance to segregate on the twin
crystal surface or on other defects, form ultra fine carbides
˚ [17] leading to greater relief of residual
of diameter 26–60 A
stress.
4.3. Grinding effects
Fig. 2. Distribution of residual stress in En 353 steel specimens after tempering.
Figs. 3–5 show the CHTUT, SCTUT and DCTUT samples
after grinding, respectively, whereas Figs. 6–8 show the CHTT,
SCTT and DCTT samples after grinding, respectively. On comparison, grinding cracks are visible in all the specimens prior to
tempering. As the samples are exposed to low temperature, the
conversion of retained austenite to martensite has manifested
234
A. Bensely et al. / Materials Science and Engineering A 479 (2008) 229–235
Fig. 3. Optical photograph showing the influence of grinding on CHTUT specimen.
Fig. 5. Optical photograph showing the influence of grinding on SCTUT specimen.
as cracks during the subsequent grinding operation. The intensity of cracks visible is different for all the samples and can
be related to the amount of untempered martensite available,
which is highly brittle in nature. The presence of grinding cracks
in SCTUT and DCTUT is more when compared to CHTUT.
This could be attributed to the conversion of retained austenite to martensite at low temperatures. The cracks visible in the
tempered specimen can be related to the fine carbide precipitation. As cryogenic treatment facilitates fine carbide precipitation
during tempering, no visible cracks are observed in SCTT and
DCTT specimen. This clearly indicates that the toughness of
the matrix holding the carbides has improved. Low temperature
conditioning of martensite accelerated the precipitation of fine
carbides during tempering. This resulted in large relaxation in
comparison with CHTT. Since there is little carbide precipitation in CHTT and due to partial transformation of austenite to
martensite, brittleness is more in CHTT samples. In SCTT and
DCTT samples, uniform relaxation has occurred in every fine
domain of the matrix resulting in reduced brittleness compared
to CHTT samples. During subsequent tempering of conventionally heat-treated samples, no appreciable change in residual
stress is noticed. This clearly shows that there is not much fine
carbide precipitation and hence a few cracks are still visible even
after tempering. The number of cracks available after tempering
in CHTT samples is larger than in the case of tempered SCTT
Fig. 4. Optical photograph showing the influence of grinding on CHTT specimen.
Fig. 6. Optical photograph showing the influence of grinding on SCTT specimen.
A. Bensely et al. / Materials Science and Engineering A 479 (2008) 229–235
235
to the specimen subjected to DCT. The decrease in the temperature increases the lattice defects and thermodynamic instability
of the martensite, which drives the carbon and alloying elements to nearby defects. These clusters act as nuclei for the
formation of fine carbides on subsequent tempering. The intensity of carbide precipitation depends on the extent to which the
specimens are cooled below room temperature. The large relaxation of residual stress in DCTT specimens indirectly reflects
about the amount of carbide precipitation. Based on relaxation
phenomenon it is concluded that DCTT exhibit the maximum
carbide precipitation than SCTT and CHTT specimens. This
has reflected in an earlier study on the improved wear resistance
obtained for En 353 after deep cryogenic treatment with tempering [14]. The study clearly shows that En 353 when subjected
to deep cryogenic process has attained the maximum compressive stress (−235 MPa) before tempering. This aids the carbide
precipitation during the subsequent tempering process resulting
in large relaxation (−80 MPa).
Acknowledgements
Fig. 7. Optical photograph showing the influence of grinding on DCTUT specimen.
The authors wish to thank the support of Ministry of Science and Computerization, Poland, towards residual stress
measurements. The authors also gratefully acknowledge M/s.
Chennai Metco Private Limited, for extending their metallurgical facilities for the successful completion of the work.
The authors gratefully acknowledge the financial support from
Department of Science and Technology, India, under FIST program (SR/FST/ETI-053/2002) for setting up a state of the art
cryogenic treatment facility, which enabled for the successful
completion of the project.
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Fig. 8. Optical photograph showing the influence of grinding on DCTT specimen.
and DCTT samples, but DCTT samples does not show any sign
of cracking after grinding operation.
5. Conclusion
The study confirms that cryogenic treatment (i.e. shallow
cryogenic treatment and deep cryogenic treatment) should be
followed by tempering. Untempered specimens of shallow cryogenic treatment and deep cryogenic treatment have surface
cracks, which is not beneficial to the life of the component.
The retention of residual stress in samples subjected to conventional heat treatment and shallow cryogenic treatment, followed
by tempering can lead to better fatigue properties, as compared
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