J. Mater. Sci. Technol., 2010, 26(12), 1107-1113.
Mechanical Properties and Temper Resistance of Deformation
Induced Ferrite in a Low Carbon Steel
Luhan Hao, Namin Xiao, Chengwu Zheng and Dianzhong Li†
Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang
110016, China
[Manuscript received August 11, 2009, in revised form March 2, 2010]
The microstructures and mechanical properties of deformation induced ferrite (DIF) in the low carbon steel
Q235 under different deformation temperatures have been investigated systematically. Through deformation
induced ferrite transformation (DIFT), ferrite grain can be refined to 3 μm and accounts for above 85% of the
overall fraction. Yield strength of DIF (>500 MPa) is increased by up to 100% compared with the conventional
low carbon steel. Comparison of microstructure and mechanical properties in the Q235 steel with DIF and
tempered DIF microstructure illustrates that the strengthening mechanism of DIF microstructure is the combination of grain boundary strengthening and carbon supersaturated strengthening. Electron back-scattered
diffraction (EBSD) analysis and high magnification scanning electron microscopy (SEM) observation denote
that high-angle grain boundary among ultrafine ferrite grain and the transformation product of retain austenite membrane along ferrite boundaries are responsible for the stability of ferrite grain size during tempering
process. Transmission electron microscopy (TEM) analysis demonstrates that the transformation product of
retained austenite membrane between ferrite grain boundaries is cementite.
KEY WORDS: Low carbon steel; Deformation induced ferrite transformation; Grain size;
Mechanical properties; Temper resistance
1. Introduction
Low carbon steels are one of the most widely used
structural materials in industry. However, their relatively low strength because of the low carbon content
and ferrite structure limits them to be used in a wider
range. It is well-known that refining grain size is an effective way to strengthen metallic material, as clearly
clarified by Hall-Petch relationship. Without adding
too much alloy elements, grain refinement strengthening is not only economical but also makes steels easily
recycled. Besides, in comparison to other strengthening mechanisms, grain refinement strengthening is
the only way to attain high strength without sacrificing toughness. During last decades, considerable
research works have been implemented in producing
† Corresponding author. Prof., Ph.D.; Tel.: +86 24 23971281;
Fax: +86 24 23891320; E-mail address: dzli@imr.ac.cn (D.Z.
Li).
steels with ferrite grain size much smaller than that
can be achieved by conventional controlled rolling[1–6] .
There are mainly two categories to fabricate ultrafine ferrite grains. One is called severe plastic deformation (SPD)[1] process such as equal channel angular pressing (ECAP)[2] , accumulative roll bonding (ARB)[3] and high pressure torsion (HPT). The
common ground of SPD is that significant deformation is needed to make ferrite grain size refined
and the process is not suitable for mass production.
The other category is advanced thermo-mechanical
control process (advanced TMCP) which is based
on conventional thermo-mechanical control process
and relatively convenient to put into effect in industries. Deformation induced ferrite transformation (DIFT, some researchers call it as strain induced
transformation[4–6] ) is a sort of advanced TMCP. It
has been proved to be an effective way to refine ferrite grains. In this approach, the refinement of ferrite
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L.H. Hao et al.: J. Mater. Sci. Technol., 2010, 26(12), 1107–1113
Fig. 1 Schematic diagrams of experiment schemes: (a) deformation at 830 and 850◦ C, (b) deformation at 780
and 800◦ C, (c) tempering at different temperatures for 1 h
is achieved by the un-saturated nucleation and limited
growth[5] . Heavy deformation above critical temperature Ar3 introduces highly deformed regions such as
deformation bands at austenite grain interiors, which
become active sites for ferrite nucleation; the ferrite
grain which nucleates and begins to grow will quickly
impinge on adjacent ferrite grains which have also nucleated in its vicinity. The accelerating cooling after deformation is important to inhibit the possibility of ferrite coarsening and reverse transformation to
austenite.
Previous work has been done on microstructures of
DIFT, effect of deformation parameters (temperature,
strain rate, strain, cooling rate etc.) on DIFT and carbon distribution of deformation induced ferrite (DIF)
in low carbon steels[7–9] . However, the mechanical
property of the low carbon steels containing DIF microstructure was usually denoted by micro-hardness
due to the lack of bulk steels[10,11] . Hence, the detailed mechanical data and strengthening mechanism
of DIF microstructure can not be clearly identified.
Besides, as a non-equilibrium transformation product,
the stability of DIF microstructure under different reheating temperature is also a crucial issue needed to
be discussed. In the present work, the detailed mechanical properties (using miniature dog-bone specimens) and microstructures of DIF in the low carbon
steel Q235 via different temperature deformations are
investigated to study the more appropriate schedules
and the strengthening mechanism of DIFT. The temper resistance of DIF microstructure and mechanical
properties of tempered DIF in the Q235 steel are also
discussed in this paper.
2. Experimental
The low carbon steel used in this experiment was
a Q235 rough slab from Anshan Iron and Steel Group
Cooperation, China. The chemical composition is
shown in Table 1. The austenite to ferrite equilibrium
transformation temperature was 845◦ C (Ae3 ) as calculated by Thermo-Calc software. After normalizing
at 920◦ C for 1 h to get uniform original microstructure, the rough slab was machined into a number of
cylindrical specimens with the diameter of 8 mm and
Table 1 Chemical composition of the Q235 steel
used in the study (%, mass fraction)
C
0.14
Si
0.24
Mn
0.64
P
0.011
S
0.005
Al
0.05
Fe
Bal.
Fig. 2 Specimens before deformation (a), after deformation (b) and for tensile test (c)
the height of 15 mm for thermal-mechanical simulation experiments on Gleeble 3500 machine (USA).
The route to get DIF microstructure and the tempering route afterwards are shown schematically in
Fig. 1. After austenitizing at 950◦ C for 5 min, the
specimens were cooled at the cooling rate of 15◦ C/s
to 780 and 800◦ C, and 1◦ C/s to 830 and 850◦ C, respectively. Then uniaxial compression was applied
to the specimens at the true strain rate of 20/s with
the height reduction of 80% (true strain 1.6). The
shapes of specimens before and after deformation are
shown in Fig. 2. Finally, the deformed specimens were
quenched by water immediately after compression to
conserve high-temperature microstructure. It is well
known that the higher the cooling rate, the lower
the ferrite transformation start temperature will be.
The above-mentioned different cooling rates between
specimens cooled to 780, 800, 830 and 850◦ C were
meant to change the critical temperature Ar3 , and
thus avoided the proeutectoid ferrite transformation.
After thermal-mechanical compression, the specimens
were machined as the specification for tensile tests
(Fig. 2(c)). The tensile specimen s dimension was designed with a gauge width of 2 mm and a gauge length
of 5 mm. Mechanical properties were measured in Instron 5848 micro-tester (USA) with the strain rate of
6×10−3 s−1 . Temper stability of DIF microstructure
was also investigated for specimens experienced deformation induced ferrite transformation above Ae3 .
L.H. Hao et al.: J. Mater. Sci. Technol., 2010, 26(12), 1107–1113
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Fig. 3 Microstructure of DIF in the Q235 steel after compressive deformation at temperature 780◦ C (a) and
850◦ C (b) with strain rate of 20 s−1 and strain of 80%
As Fig. 1(c) shows, the specimens were tempered
at 200, 400 and 600◦ C for 2 h followed by furnace
cooling, respectively. DIF specimens were sectioned
through the center parallel to the compression axis
for microstructure observations. After the center sections were ground, polished, and etched with a 4% nital solution, the microstructures were observed by optical microscopy (OM, ZEISS AXIOVERT 200MAT,
Germany), scanning electron microscopy (SEM, Shimadzu Superscan SS-550, Japan) and transmission
electron microscopy (TEM, FEI Tecnai T20, Japan).
The electron back-scattered diffraction (EBSD, Hitachi S-3400N, Japan) scanning was carried out at the
center of the center section, and in a 50 μm×100 μm
area at a step size of 0.4 μm on a hexagonal grid.
Quantitative analysis of phase fraction and average
grain size was evaluated on SISC-IAS analysis software.
3. Results
3.1 Microstructure
Figure 3 shows the optical microstructures of asquenched DIF specimens after compressive deformation at temperature 780 and 850◦ C, respectively. It
can be seen that the two microstructures are mainly
composed of equiaxed ferrite grains. Ferrite is the major phase with the average grain size of around 3 μm
and accounts for above 85% of the overall fraction,
as shown in the white areas. The microstructures in
Fig. 3(a) and (b) are nearly the same except a slight
difference in grain size, which is undistinguishable.
Figure 4 represents the SEM microstructures of
DIF specimens tempered at different temperatures for
1 h. It is hard to make out the difference in grain
size from these photos. Quantitative analysis indicates that tempering temperature has little effect on
ferrite grain size. Large quantity of cementite particles present at ferrite grain boundaries and threepoint junctions when the tempering temperature is
600◦ C.
3.2 Mechanical properties
Table 2 shows the mechanical properties of DIF
specimens under different deformation temperatures.
It can be seen that the yield strength of the steel is
around 500 MPa, which is increased by up to 100% as
compared to that of conventional Q235 steel. The tensile strength is also increased to 550 MPa and the elongation is still in the appropriate range around 35%.
The mechanical data demonstrate that DIFT is an
effective way to enhance the strength.
The tempered DIF specimen s mechanical properties are shown in Table 3. Compared with the untempered one, the ductility is increased while the strength
is decreased. With the increase of tempering temperature, the disparity in strength and ductility between untempered specimen and tempered specimen
becomes larger.
DIFT has the potential to refine ferrite grain size
to a larger extent than conventional TMCP. For low
carbon steel, the ferrite grain size in conventional
TMCP is around 10–20 μm and the yield strength is
in the range of 200–300 MPa[12] . In this work, about
3 μm ferrite grain size can be obtained through appropriate DIFT process. The mechanism of DIFT
is to increase ferrite nucleation positions and ferrite
transformation temperature through heavy deformation above Ar3 . It is usually considered that DIFT
occurs during the deformation process[4] and the high
nucleation rate is the key factor for the refinement of
ferrite grains.
4. Discussion
4.1 Strengthening effect of DIFT
The data of strength in Table 2 show a difference of
40 MPa in yield strength between deformation at 780
and 850◦ C. It is because that the lower deformation
temperature provides larger driving force for ferrite
transformation. And the low temperature makes deformation energy easy to be stored. Those will lead
to the higher nucleation rate for the low temperature
DIF under the same other deformation parameters,
which means that the low temperature induced ferrite
is a little finer than that of high temperature (Fig. 3).
According to Hall-Petch relationship, the finer the av-
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Fig. 4 Microstructure changes of DIF in the Q235 steel after tempering at different temperatures for 1 h: (a) untempered, (b) 200◦ C, (c) 400◦ C, (d) 600◦ C
Table 2 Mechanical properties of the Q235 steel with the microstructure of proeutectoid ferrite and DIF treated
under different deformation temperatures with the strain rate of 20 s−1 and strain of 80%
Deformation
temperature
Without deformtion
780◦ C
800◦ C
830◦ C
850◦ C
Yield strength,
σy /MPa
245.6
528.5
503.1
515.0
482.0
Tensile strength,
σb /MPa
438.2
571.9
551.2
556.8
542.7
erage grain size, the higher the yield strength will
be. Hence, from the thermodynamic point of view, it
is reasonable that 780◦ C deformation induced ferrite
strength is higher than that of 850◦ C. The strengthening of steels is usually accompanied with ductility loss.
For the specimens deformed at 780◦ C, uniform elongation is reduced to 11.5% compared with 15%–18%
in the relatively high temperature of 850◦ C (Table
2). It is mainly attributed to the more densely distributed grain boundaries that make dislocation and
slip bands hard to move and thus inhibit the plastic deformation. It can be concluded that 800–830◦ C
is the more appropriate temperature range to induce
fine ferrite grains, at which an excellent combination
of strength and ductility can be obtained.
In addition, the presented mechanical properties
data of DIF specimens indicate the decrease of significant work hardening. This is reflected by the increase of yield ratio, which alters from 0.56 of the
proeutecoid ferrite to 0.9 of DIF. Except the influence
of fine ferrite grain size, residual dislocations and internal stresses brought by heavy deformation and accelerating cooling rate are partially responsible for the
decrease in strain hardening. This result corresponds
Uniform elongation,
δu /%
22.5
11.5
17.2
17.8
14.9
Total elongation,
δu /%
37.0
29.2
36.2
40.3
36.2
Yield ratio,
σy /σb
0.56
0.92
0.91
0.92
0.89
reasonably well with the previous work by Miller[13] in
which refinement of ferrite leads to an enhancement in
yield strength with little work hardening. According
to Hart and Considere s criterion—the onset of plastic instability in tension occurs when σ≥dσ/dε, where
σ is the flow stress and dσ/dε is the work hardening
rate, and a higher work hardening rate is essential to
enhance uniform elongation in steels. It seems that
the ways to reduce yield ratio and increase work hardening behavior of ultrafine ferrite, such as introducing
some hard second phase to the microstructure[11,14]
and fabricating a layered composite microstructure[15]
will exert favorable influence on ductility.
4.2 Temper resistance and strengthening mechanism
of DIFT
The possibility of remedying the ductility loss of
DIF specimens due to residual stress caused by deformation and quenching process and the stability of
high temperature DIF specimens was investigated by
the tempering experiment. Table 3 demonstrates that
the ductilities of tempered specimens are indeed elevated but at the expense of strength. Low temper-
L.H. Hao et al.: J. Mater. Sci. Technol., 2010, 26(12), 1107–1113
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◦
Table 3 Mechanical properties of the Q235 steel after DIFT at 850 C and tempered at different
temperatures for 1 h
Tempering
temperature
Untempered
200◦ C
400◦ C
600◦ C
Yield strength,
σy /MPa
482.0
446.0
408.0
380.0
Tensile strength,
σb /MPa
542.7
515.0
506.0
474.0
ing temperature 200◦ C is effective to eliminate intrastress in DIF specimen which can make uniform elongation enhance 5% and total elongation increases
to the value of proeutectoid ferrite (around 43%).
With the enhancement of tempering temperature, the
strength gets further decreased while the total elongation remains around 42%. The average grain sizes of
the tempered specimens maintain at about 3 μm. Detailed quantitative analysis data of average grain size
after tempered at 200, 400 and 600◦ C are 2.81, 2.83
and 2.85 μm, respectively. In comparison to 2.84 μm
of untempered ferrite, average ferrite grain size can
be considered as invariable during tempering process
for 1 h. Nearly equivalent ferrite grain size and an
apparent difference in strength denote that grain refinement is not the only strengthening mechanism for
the DIF specimens.
Previous work[16] has proved that DIF is a kind of
carbon supersaturated ferrite by using electron probe
microanalyzer (EPMA) line analysis of carbon distribution in proeutecoid ferrite and DIF. The supersaturated carbon in ferrite tends to diffuse to grain
boundaries or precipitate within grain interior during
long soaking time at the moderate tempering temperature. It demonstrates that carbon supersaturation
in untempered DIF microstructure is responsible for
the extra strength increment compared with the tempered ones. It can be concluded that Hall-Petch relationship fails to reflect the exact relationship between
the average ferrite grain size and strength of DIF because of the effect of complicated carbon distribution
in DIF microstructure. During 200–400◦ C tempering process, no significant change in grain boundaries
occurs but the amount of dispersed carbide at grain
interior increases, as shown in Fig. 5(a) and (b). Figure 5(c) and (d) shows abundant globular carbides
precipitate along grain boundaries or at some threepoint junctions at tempering temperature of 600◦ C.
Yield ratio drops from 0.9 to 0.8, denoting the working hardening effect of dispersed cementite. Some
researchers[17] deduced that precipitation strengthening through pinning effect of dispersive carbide to dislocations can further strengthen the steel with fine
ferrite microstructure. The fact does not go in this
way as shown in Table 3. The result demonstrates
that, as far as DIF microstructure is concerned, the
effect of precipitation strengthening is not as good as
carbon supersaturated strengthening.
DIFT is usually viewed as non-equilibrium phase
transformation in respect that at the deformation
Uniform elongation,
δu /%
14.9
20.0
18.8
22.0
Total elongation,
δt /%
36.2
43.0
42.0
43.0
Yield ratio,
σy /σb
0.89
0.87
0.81
0.80
temperature, austenite which has lower free-Gibbs
energy is the stable phase in deformation-free sample. Non-equilibrium microstructure tends to transform into equilibrium status during reheating process
by rearrangement of atoms and grain growth. It has
been proved that grain growth does not occur during
the tempering process. EBSD measurement of untempered specimen was performed in order to further
understand DIFT and explain the stability of ferrite
grain size. EBSD boundary misorientation map of
DIF microstructure at the deformation temperature
850◦ C is shown in Fig. 6. The bold black lines denote
high angle grain boundaries (HAGBs) of which misorientation angles are higher than 15 deg., while the
narrow gray lines denote low angle grain boundaries
(LAGBs) of which misorientation angles are 2–15 deg.
The boundary misorientation maps reveal that ultrafine ferrite grains are mostly surrounded by HAGBs
and little substructure exists in the grain interior.
Humphreys et al.[18] suggested that the microstructure should be stable against recrystallization and
coarsen uniformly, provided that it contained HAGBs
greater than about 65%. Hurley et al.[19] also suggested that the degree of misorientation between the
ferrite grains limits grain coalescence and growth of
ultra-fine ferrite. It seems that the stability of ferrite
grain size during tempering process can be roughly
explained by the HAGBs after DIFT. However, considering that recrystallized grains which have HAGBs
in the same way can still grow at elevated temperature, the explanation using HAGBs seems untenable.
Compared DIF with ferrite-pearlite microstructure,
as shown in Fig. 7, a huge discrepancy in grain boundary can be found. Ferrite-pearlite microstructure exhibits simple narrow boundaries, whereas most DIF
grain boundaries have another phase occupied, making boundary shown as bright white (as ellipse shown
in Fig. 7(b)). This phase is deemed as the cooling
product of residual austenite membrane (RAM) which
is left between fast nucleated ferrite grains because of
the high ferrite formation rate. TEM analysis denotes that the RAM is cementite, as shown in Fig. 8.
Due to the mass of carbon remaining in the austenite solution, a significant amount of cementites were
generated along ferrite grain boundary during the accelerating cooling process. They acted as the obstacle
of grain boundary movement by applying pinning effect to boundaries. It is suggested that the stability
of ferrite grain size during tempering process is the
cooperation of HAGBs and pinning effect of cemen-
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Fig. 5 Evolution of dispersed carbides during tempering process: (a) 200◦ C, (b) 400◦ C, (c) and (d) 600◦ C
Fig. 6 Boundary misorientation maps (a) and distribution of misorientaion angle of adjacent ferrite grains (b)
after DIFT at 850◦ C, with strain rate of 20 s−1 and strain of 80%
Fig. 7 Comparison of grain boundary among proeutectoid ferrite (a) and DIF (b) of the steel Q235
L.H. Hao et al.: J. Mater. Sci. Technol., 2010, 26(12), 1107–1113
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Acknowledgement
This work was financially supported by the National
Natural Science Foundation of China (NSFC) under Grant
No. 50871109.
REFERENCES
Fig. 8 Cementite on the DIF grain boundary
tite along grain boundaries.
5. Conclusions
(1) Above 500 MPa of yield strength and 550 MPa
of tensile strength can be obtained through appropriate deformation induced ferrite transformation parameters for low carbon steels. The ductility of DIF is
a little decreased compared with proeutectoid ferrite
but still in the reasonable range. 800–830◦ C is the
favorite temperature range to form high performance
DIF for the low carbon steel Q235.
(2) With the increase of the tempering temperature, the strength of the steel gets decreased step
by step, while the grain size of DIF remains in the
same grade. The strengthening mechanism of DIFT
is the coupling of grain boundary strengthening and
carbon supersaturated strengthening. And for DIF
microstructure, carbon supersaturated strengthening
is better than precipitation strengthening of carbide.
(3) DIF grains are mostly surrounded by HAGBs
and little substructure exists in the grain interior. The
cementite between adjacent grains and HAGBs are
the keys to keep non-equilibrium DIF stable in grain
size.
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