FIERF/FIA FINAL REPORT: Ultra-Fine Grain Processing of ...
Transcript of FIERF/FIA FINAL REPORT: Ultra-Fine Grain Processing of ...
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FIERF/FIA FINAL REPORT: Ultra-Fine Grain Processing of
Steel Billet – Phase I - Effect of Initial Microstructure on the Microstructural Evolution and Tensile Properties of 4140 Steel
Thomas Kozmel, Ryan Cassel, and Sammy Tin
Illinois Institute of Technology 10 W 32nd St.
Chicago, IL 60616 [email protected]
Phone: 1-312-567-3780 Fax: 1-312-567-7230
April 27th, 2015
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Abstract:
4140, a high strength steel, was heat treated from its initial pearlitic and ferritic
microstructure to produce three discretely different microstructures: pearlitic and ferritic,
bainitic, and martensitic microstructures. These microstructures were then warm rolled at
temperatures between 866 K (593 °C) and 977 K (704 °C) to produce varying levels of
dynamically recrystallized microstructures. Once a height reduction of 80% had been reached,
samples were machined into dog-bone tensile samples. Tensile tests revealed yield strengths
between 600 and 1160 MPa, which were shown to be a function of the processing parameters.
Electron backscatter diffraction (EBSD) was used to analyze the deformed microstructures.
Constitutive analytical-empirical models were developed such that the yield stress of the material
could be directly predicted from the initial microstructure and its processing parameters.
Keywords: Steel; EBSD; Microstructure formation mechanism; Mechanical properties testing;
Modeling
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1. Introduction:
4140 is an alloy steel can be processed and heat treated to readily obtain yield stresses as
high as 1650 MPa while still retaining good ductility. Its properties make it suitable for
applications where high strength and good toughness are required, but where the service
conditions only expose the material to relatively low to moderate temperatures. The carbon
content of 4140 gives it excellent hardenability and wear resistance. Overall, the properties of
this steel make it suitable for use in connecting rods, crankshafts, industrial tooling, high
pressure tubing, and other applications [1]. Although 4140 steel exhibits a well-balanced
combination of properties, the structural properties of this alloy may potentially be further
enhanced, however, by taking advantage of recent developments in grain size refinement
technologies [2-3]. For structural materials, ultra-fine grain (<100 nm) processing has been
shown to enhance both strength and toughness simultaneously. Compared to other common
strengthening mechanisms, such as work hardening [4], strengthening via grain refinement
typically does not accompany a corresponding reduction in ductility. Various studies have shown
that a variety of severe plastic deformation (SPD) techniques can be utilized to form physically
large bulk structures with homogeneous ultra-fine grained microstructures. Commonly used SPD
techniques include equal channel angular pressing (ECAP) [5-7], high pressure torsion (HPT) [8-
9], accumulative roll bonding [10], accumulative angular drawing [11], multi-axial forging
(MAF) [12-15], and even conventional rolling [16-17]. Grain refinement and the formation of
sub-micron grain sizes has been shown to be achievable by all of these methods provided that
sufficiently high plastic strains are achieved [18-19]. In addition to plastic strain, the deformation
temperature is also a key factor, as it controls both the recrystallization and grain growth
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kinetics. Recent studies have shown that warm rolling of steel at temperatures below those
typically used in conventional hot rolling can be used to induce the formation of sub-micron
grains via recrystallization and restrict their subsequent growth [16-17].
Typically, deformation temperatures used for thermal – mechanical processing of steels
are selected such that they are well above the corresponding austenitizing temperature (AC3).
Under these temperature conditions, the microstructure of the material consists of a single FCC
phase and both recrystallization and grain growth occur rapidly as the thermal energy provided to
the system enhances the kinetics of the system. As a result, forming and retaining
microstructures comprised of sub-micron grains during thermo-mechanical deformation becomes
challenging. For most commercial steel alloys, the recrystallization and grain growth behavior
under these conditions have been well characterized [20-22]. When compared to thermal –
mechanical deformation of steels below the austenitizing temperature, where the microstructure
predominaetely consists of ferrite and carbides, the transformation kinetics are greatly reduced
and recrystallization and grain growth are retarded [23]. However, in order to induce grain
refinement during deformation at these temperatures, comparatively higher levels of plastic
strain are required due to the higher stacking fault energy associated with the BCC crystal
structure of the ferrite. The higher stacking fault energies prevent dislocations induced during
deformation from dissociating into stacking faults that restrict dislocation climb and contributes
to enhanced dislocation mobility that makes BCC metals more likely to accommodate
deformation by recovery [24-25]. Hence, significantly higher overall levels of plastic strain are
required in order to induce continuous dynamic recrystallization (CDRX) when deforming steels
below their respective austenitizing temperature [26-29]. It should also be noted that the presence
of carbides within the ferritic structure may also effect the recrystallization and grain growth
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kinetics as finely distributed carbides, serve to pin both dislocations and mobile grain
boundaries.
Throughout many alloy systems during the initial stages of continuous dynamic
recrystallization, the plastic strain induces the formation of cells that are bounded by dense
dislocation walls (DDWs). The interior of these cells contain relatively few dislocations and
their size is dependent on the kinetics and mobility of the dislocations. These cells or “subgrain”
structures form as the material attempts to rearrange the accumulated dislocations into
configurations that possess the lowest total energy. As these structures form, the initial relative
misorientation between adjacent dislocation walls is typically less than 1 degree. However, as the
plastic strain levels increase, the dense dislocation walls begin to evolve and the misorientations
within the sub-grain boundaries gradually increase such that they may eventually rotate to form
high angle boundaries and become continuously dynamically recrystallized. The mechanisms
associated with dynamic recrystallization may be iterative as dislocation walls and tangles may
continue to form within the newly formed grains with continued straining and serve to further
refine the system [26, 30-32].
Several recent studies have correlated the effects of initial microstructure with the
resulting microstructure of a given set of processing parameters. For example, with respect to
4140 steel, it was found that quenching and tempering a martensitic structure resulted in a finer
grain size than quenching and tempering an equivalent bainitic microstructure. However, upon
warm rolling at 773 K (500 °C), the initially bainitic microstructure exhibited a higher degree of
grain refinement than that of the martensite [33]. Effects of initial microstructure have also been
observed in Ti-6Al-4V [34], where the starting grain structure and morphology affected the
surface finish and flow behavior during ECAP. Hence, it is apparent that the starting
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microstructure may have a significant impact the formation of ultra-fine grained and sub-micron
grain structures, particularly in steel alloys where the microstructures may be comprised of
ferrite, bainite or martensite. For this reason, the objective of the current study was to quantify
and elucidate the effects of initial microstructure in 4140 steel on its recrystallization kinetics
during warm rolling at sub-AC1 temperatures.
2. Experimental:
Billet bar stock of 4140 steel was cut into nine 10.16 cm long sections with a cross
section of 2.286 cm in width by 0.94 cm in height. The initial microstructure was confirmed via
scanning electron microscopy (SEM) to be predominantly pearlitic, Figure 1a and 2a. Nominal
composition values for the 4140 steel used in this investigation is listed in Table I.
Table I: Nominal Elemental Composition for 4140 Steel, by % wt. Cr Mn C Si Mo S P Fe 0.80-1.10%
0.75-1.00%
0.38-0.43%
0.15-0.30%
0.15-0.25%
0.040% 0.035% Bal.
In order to form the distinct starting microstructures, the 4140 material was austenitized at 800
which corresponded to 15 K (15 °C) above the austenitizing temperature (AC3) and cooled at
different rates. Thermocouples were spot welded to the ends of the bars to monitor their
temperature during the heat treatment process. The bars reached the target temperature in
approximately 25 minutes, and were held at this temperature for an additional 20 minutes to
ensure a uniformly austenitized microstructure. Martensitic structures were formed by water
quench to room temperature at a cooling rates of ~30 K/s (30 °C/s).
Bainitic microstructures were formed by removing the bars from the furnace and placing them on
323 K (50 °C) at a rate of
approximately 0.83 K/s (0.83 °C/s).
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Each of the three different initial microstructures was warm rolled to 33%, 50%, 67%,
and 80% height reductions at the temperatures of 866 K, 922 K, and 977 K (593 °C, 649 °C, and
704 °C). Bars were individually isothermally heat treated to the target rolling temperature for
approximately 20 minutes to ensure a uniform internal temperature. The bars were then rolled to
the designated height reduction while being placed back into the furnace for 10 minutes after
every 2 rolling passes to retain the target rolling temperature. [roller speed and strain rate] The
rolling height reductions ranged from 0.076 cm down to 0.0076 cm per pass as deformation
progressed and the height of the samples decreased. After each bar was rolled to the appropriate
height reduction, it was immediately water quenched. Specimens were excised from each of the
samples for detailed microstructural characterization.
Bars in their heat treated state, and those samples that were rolled to various height
reductions were prepared for characterization using standard metallographic techniques and a
finishing polish using a 0.06 micron colloidal silica suspension. Samples were characterized and
quantitatively analyzed using electron backscatter diffraction (EBSD) in a JEOL-5900LV SEM
equipped with an Oxford Instruments Nordlys-HKL EBSD detector. For all samples, the rolling
direction was analyzed at all height reductions less than 80%. Samples reduced by 80% in height
were analyzed along the normal direction. Measurements were made to ensure that the center of
the cross section on the rolled-direction face was being characterized on each sample for
comparative purposes. 80% height reduced samples were likewise analyzed in the center of the
bar along the normal direction. EBSD scans were conducted by searching for both BCC and FCC
iron with the settings of 7 bands detection, 67 reflectors for the BCC iron, and 56 reflectors for
the FCC iron. Post processing was conducted within the Oxford HKL Channel 5 – Tango
software program. Conservative noise reduction and extrapolation was performed prior to the
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quantitative analysis of the inverse pole figure (IPF) maps. Grain boundaries were defined as
boundaries exhibiting greater than 15 degree critical misorientation with respect to their
neighbors. Average grain sizes were determined by the mean circle equivalent diameter. For a
grain to be considered recrystallized, size and shape requirements would need to be met,
depending on the initial microstructure.
Flat, dog-bone tensile samples were machined from the final 80% height reduction of
each rolled bar. Two tensile samples were obtained from each bar. Dimensions of the tensile
samples were 0.635 cm in width by approximately 0.188 cm in thickness. Gage sections were
2.54 cm long, and samples were tested in tension at a rate of 0.127 cm/min using an Instron
tension test machine with a 200 kN (45000 lbf) load cell. Tensile force and elongation were
recorded for each bar until failure. These data were used to construct stress-strain curves for each
bar, from which yield strength, ultimate tensile strength, elastic modulus, and percent ductility
were calculated. Because two samples were tested for each bar, these values were found by
averaging the two data sets.
3. Results:
3.1 Microstructural Characterization:
To obtain three different initial microstructures, sample bars with a mixed pearlitic and
ferritic microstructure were heat treated to obtain fully martensitic and fully bainitic
microstructures, Figure 1. Each of the starting microstructures was carefully characterized using
both optical and electron microscopy and evaluated using EBSD. For the pearlitic and ferritic
microstructure, the average grain size was 10.7 microns, Figures 1a. It was determined through
both optical microscopy and SEM that the phase fraction of pearlite was approximately 75%,
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Figure 2a. K/s, the microstructure of 4140 was
found to completely transform into martensite, Figure 1b. The average lath size measured via
EBSD was 2.4 microns and the average grain size was observed to range between 5-10 microns,
Figure 2b. A predominately bainitic microstructure was formed after cooling the 4140 samples at
a cooling rate of 0.83K/s, Figure 1c. A mixture of both upper and lower bainite was observed
within the microstructure, Figure 2c, and the area fraction of bainite grains within the
microstructure was nominally at least 85%. The remaining regions of the microstructure were
comprised of martensite and ferrite. In this instance, the average grain size was determined to be
3.4 microns.
Figure 1: Starting microstructures prior to warm rolling: a) the as received pearlitic and ferritic microstructure, b) the martensitic microstructure obtained through annealing and quenching, and c) the mixed bainite microstructure obtained through annealing and air quenching on a brass plate.
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Figure 2: SEM images verifying the phase components for the a) pearlitic microstructure, b) martensitic microstructure, and c) bainitic microstructure.
After warm rolling to reduce the height of all samples by ~33%, intragranular lattice
rotations and the formation of subgrain structures was observed in all samples. For example, in
the initially pearlitic and ferritic sample rolled at 866 K (593 °C), Figure 3, the local variations in
color within the grains indicate that subgrain formation was occurring. Recrystallized grains
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were also observed, which possess a reduced grain size relative to the rest of the microstructure.
With a warm rolling strain rate of X and a process that utilizes multiple rolling passes of the
deformed structures, both dynamic, meta-dynamic recrystallization mechanisms are likely
responsible for the refinement of the grain structure [ref]. When characterizing the as-deformed
microstructures, distinguishing between these recrystallization mechanisms is challenging and
beyond to scope of the present investigation. As a result, for the purposes of this investigation,
all of grain refinement was considered to occur dynamically.
Figure 3: Initially pearlitic and ferritic sample reduced by 33% in height via rolling at 866 K (593 °C), exhibiting a partially recrystallized structure. To best compare the effects of temperature and initial microstructure on the
recrystallization response of the 4140 steel, the samples warm rolled to 50% height reductions
for all temperatures and initial microstructures were selected, Figure 4. These observations
clearly show that the size of the recrystallized grains typically increased with increasing
temperature. For instance, following refinement of the pearlitic microstructure, the mean
recrystallized grain size was 0.8 microns at 866 K (593 °C), Figure 4a, which increased to 0.9
microns at 922 K (649 °C), Figure 4b, and finally to 1.0 microns at 977 K (704 °C), Figure 4c. In
the samples with the martensitic starting microstructure, the average recrystallized grain size
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increased from just under 0.6 microns, Figure 4d, to just over 0.6 microns, Figure 4e, and finally
to 0.7 microns, Figure 4f, at 866 K, 922 K, and 977 K (593 °C, 649 °C, and 704 °C) respectively.
Finally, the initially bainitic samples showed an increase in dynamically recrystallized grain size
from 0.7 microns, Figure 4g, to 0.9 microns, Figure 4h, and remaining at 0.9 microns, Figure 4i,
as the temperature increased from 866 K, to 922 K, and to 977 K (593 °C, to 649 °C, and to 704
°C) respectively. Effects of initial microstructure were also evident, as the initial microstructure
which appeared to generate the smallest dynamically recrystallized grains at a given temperature
was martensite, and the largest dynamically recrystallized grains were generated by the initially
pearlitic samples. In addition, differences in fraction of dynamic recrystallization were observed
between different starting microstructures. For instance, the pearlitic samples exhibited the
lowest fraction of dynamic recrystallization following warm rolling to a 50% height reduction,
Figure 4a-c, with values varying only slightly with temperature from 14%, to 10%, and to 5% as
the temperature increased from 866 K, to 922 K, and finally to 977 K (593 °C, 649 °C, and 704
°C). However, the martensite samples, Figure 4d-f, exhibited the largest fraction of dynamic
recrystallization at this height reduction. These values ranged from 78% at 866 K (593 °C), to
76% at 922 K (649 °C), and finally to 65% at 977 K (704 °C). Finally, the bainitic samples warm
rolled to 50% height reduction, Figure 4g-i, were observed to exhibit dynamically recrystallized
area fractions of 63%, 44%, and 30% at 866 K, 922 K, and 977 K (593 °C, 649 °C, and 704 °C)
respectively. Although the general trend at this height reduction indicated that samples rolled at
lower temperatures exhibited a larger area fraction of dynamically recrystallized grains, this
trend was least pronounced in the samples with the pearlitic starting microstructures. However,
this trend held for all instances in the initially martensitic samples. It also should be noted that
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the bainitic samples behaved differently at higher strain levels, where the difference between
rolling temperatures was not significant.
Figure 4: Comparison of and dynamic recrystallization and dynamically recrystallized grain size trends as a function of rolling temperature and initial microstructure. Pearlitic samples rolled at a) 866 K (593 °C), b) 922 K (649 °C), and c) 977 K (704 °C). Martensitic samples rolled at d) 866 K (593 °C), e) 922 K (649 °C), and f) 977 K (704 °C). Bainitic samples rolled at g) 866 K (593 °C), h) 922 K (649 °C), and i) 977 K (704 °C). Following warm rolling, some of the samples exhibited significant grain refinement as
the microstructures were found to completely or nearly completely dynamically recrystallize
(>75% area fraction recrystallization). These samples included the initially bainitic samples
rolled at 977 K (704 °C) to height reductions of 67% and 80%, and the martensitic samples
rolled at 866 K (593 °C) and 922 K (649 °C) to height reductions of 67% and 80%. Figure 5
shows the microstructure of an initially bainitic sample rolled at 977K (704 °C) to a height
reduction of 67%. This microstructure exhibited full dynamic recrystallization with a mean grain
size of 0.6 microns. These grains were observed to be nominally equiaxed with little difference
in size.
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Figure 5: The initially bainitic microstructure warm rolled at 977 K (704 °C) such that its height was reduced by 67%. The processed microstructure shown exhibited full dynamic recrystallization. 3.2 Tensile Results:
Two tensile specimens were machined from samples warm rolled to a height reduction of
80%, and loaded in tension until failure at room temperature. Values for key data points such as
the yield strength and elongation to failure were averaged between the two samples and reported
as such. As expected, the yield stress for all samples decreased as a function of rolling
temperature. For example, the initially pearlitic microstructure exhibited a yield stress of 1090
MPa after being rolled to 80% height reduction at 866 K (593 °C), which was reduced to 869
MPa as the rolling temperature was increased to 922 K (649 °C), and finally to 627 MPa at the
rolling temperature of 977 K (704 °C), Figure 6a. In Figure 6b, the initially martensitic
microstructure had yield strengths of 1160 MPa, 894 MPa, and 651 MPa for the corresponding
rolling temperatures of 866 K, 922 K, and 977 K (593 °C, 649 °C, and 704 °C) respectively.
Finally, Figure 6c shows the stress-strain curves for the initially bainitic microstructure rolled to
80% height reduction with respective yield strengths of 1087 MPa, 828 MPa, and 613 MPa at the
rolling temperatures of 866 K, 922 K, and 977 K (593 °C, 649 °C, and 704 °C). In addition,
ductility was also observed to increase in both the initially pearlitic and bainitic tensile samples.
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The initially bainitic sample warm rolled at 977 K (704 °C) had the best ductility of all samples,
reaching 11.7% strain. However, it was observed that the initially martensitic microstructures
experienced a decrease in ductility as the rolling temperature increased. These samples typically
had the least ductility. For example, the initially martensitic microstructure rolled at 977 K (704
°C) only withstood elongation to 1.0% strain, which was the lowest value recorded.
Figure 6: Tensile stress strain curves and results for samples rolled to 80% height reduction with: a) the initially pearlitic microstructure, b) the initially martensitic microstructure, and c) the initially bainitic microstructure.
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3.3 Development of a Microstructural Process Model:
Based on the observed microstructural changes collected during this investigation,
constitutive relationships were developed to describe the microstructural evolution of the
materials as a function of the initial microstructure and the processing parameters. This set of
expressions considers the effects of temperature, effective strain, and the starting microstructure
before rolling.
In order to best relate the rolling process to the microstructural evolution of the material,
the various height reduction steps were quantified in terms of their effective strain on the
material, Equation 1.
𝜀𝜀 = 𝜀𝜀 + 𝜀𝜀 + 𝜀𝜀/
(1)
In Equation 1, 𝜀𝜀 is the effective strain, and 𝜀𝜀 , 𝜀𝜀 , and 𝜀𝜀 are the respective strains in each
direction of the rolled bar, calculated from the change in dimensions of the bar as it was
processed. The strain based on the change in dimension in each respective direction was
calculated by Equation 2, where 𝜀𝜀 is the strain, DL is the change in length for the respective
direction, and L0 is the original length of that dimension prior to processing.
𝜀𝜀 = ∆ (2)
Typically, the fraction of dynamic recrystallization, 𝑋𝑋 , occurring at a specific effective
strain and temperature is described by an Avrami relation, Equation 3:
𝑋𝑋 = 1− 𝑒𝑒𝑒𝑒𝑒𝑒 −ln (𝛼𝛼 ) .
. .
.
(3)
Where the effective strain required to initialize dynamic recrystallization is assumed to be the
effective strain required for 5% dynamic recrystallization, 𝜀𝜀 . , and is defined as:
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𝜀𝜀 . = 𝐴𝐴 𝑇𝑇 (4)
Where T is the rolling temperature in degrees Celsius, and A1, k1 are constants. Additionally, the
effective strain required for 50% dynamic recrystallization is defined by:
𝜀𝜀 . = 𝐴𝐴 𝑇𝑇 (5)
Where T is the rolling temperature in degrees Celsius, and A2, k2 are constants. The remaining
terms in Equation 3, a1, and n, are constants.
The size of the dynamically recrystallized grains, 𝐷𝐷 , in microns, may be described by
a decaying exponential function, Equation 6:
𝐷𝐷 = 𝐵𝐵exp (6)
Where b1 is a constant, T is the rolling temperature and the lead term B can be described by:
𝐵𝐵 = 𝐴𝐴 𝑇𝑇 (7)
Where A3, k3 are constants, and T is the rolling temperature.
Once the constants from Equations 4 and 5 are determined, the parameters from Equation
3 can be back calculated from the data. Likewise, the constants from Equation 6 can be
determined from the data after Equation 7 is developed. Equations 4, 5, and 7 are used to
calibrate the initialization point, inflection point, and y-intercept within the data, respectively,
and to ensure that the proper variation with respect to temperature is achieved. The fully solved
equations for each respective starting microstructure are shown as follows:
Initially Pearlitic and Ferritic Microstructure:
𝑋𝑋 = 1− 𝑒𝑒𝑒𝑒𝑒𝑒 −𝑙𝑙𝑙𝑙 (1.834) .
. .
. . (8)
𝜀𝜀 . = 2.358 ∗ 10 𝑇𝑇 . (9)
𝜀𝜀 . = 0.6464 𝑇𝑇 . (10)
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𝐷𝐷 = 𝐵𝐵exp . (11)
𝐵𝐵 = 9.196 ∗ 10 𝑇𝑇 . (12)
Initially Bainitic Microstructure:
𝑋𝑋 = 1− 𝑒𝑒𝑒𝑒𝑒𝑒 −ln (2.088) .
. .
. . (13)
𝜀𝜀 . = 8.353 ∗ 10 𝑇𝑇 . (14)
𝜀𝜀 . = 9.995 ∗ 10 𝑇𝑇 . (15)
𝐷𝐷 = 𝐵𝐵exp . (16)
𝐵𝐵 = 5.448 ∗ 10 𝑇𝑇 . (17)
Initially Martensitic Microstructure:
𝑋𝑋 = 1− 𝑒𝑒𝑒𝑒𝑒𝑒 −ln (2.019) .
. .
. . (18)
𝜀𝜀 . = 8.079 ∗ 10 𝑇𝑇 . (19)
𝜀𝜀 . = 9.031 ∗ 10 𝑇𝑇 . (20)
𝐷𝐷 = 𝐵𝐵exp . (21)
𝐵𝐵 = 2.138 ∗ 10 𝑇𝑇 . (22)
These expressions can be visualized in Figures 7 and 8, which describe the fraction of dynamic
recrystallization as a function of the processing parameters (Figure 7), and the dynamically
recrystallized grain size as a function of the processing parameters (Figure 8). In each figure, a),
b), and c) are shown for the initially pearlitic and ferritic, bainitic, and martensitic
microstructure, respectively. These figures also compare the model predictions and the
experimental data. Error bars are included in these plots as Standard Error(s) versus each
individual trend produced from the global model.
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Figure 7: The fraction of dynamic recrystallization for the initially a) pearlitic and ferritic, b) bainitic, and c) martensitic microstructure as a function of the processing parameters. A comparison of the model with the experimental data is shown.
Figure 8: The dynamically recrystallized grain size for the initially a) pearlitic and ferritic, b) bainitic, and c) martensitic microstructure as a function of the processing parameters. A comparison of the model with the experimental data is shown.
3.4 Correlation between Microstructure and Mechanical Properties:
Following the development of the microstructural model, the resulting microstructures
were incorporated into constitutive expressions for predicting the flow behavior of the material.
Typically, the Hall-Petch equation is used to relate grain size to the yield strength of a given
material. The form of this equation is shown as Equation 23:
𝜎𝜎 = 𝜎𝜎 + 𝑘𝑘 𝐷𝐷 / (23)
Where 𝜎𝜎 is the yield stress, D is the grain size of the material, and 𝜎𝜎 , 𝑘𝑘 are constants.
However, for very fine grain sizes (<1 µm) a departure from the typical Hall-Petch trend occurs.
Therefore, a modified equation valid for fine grain sizes is necessary, Equation 24 [35]:
𝜎𝜎 = + 𝛽𝛽 𝐷𝐷 (24)
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Where 𝛼𝛼 ,𝛽𝛽 are constants. For purposes of this study, a global equation of the following form is
used:
𝜎𝜎 = + 𝛽𝛽 𝐷𝐷 𝑋𝑋 (25)
Where 𝐷𝐷 ,𝑋𝑋 are defined previously in section 3.2. These terms are combined into the D
term to weight the effects of the fraction of dynamic recrystallization on the grain size and
ending yield stress. When this equation is solved for the experimental data, using model values
for 𝐷𝐷 , and 𝑋𝑋 , it takes the form:
𝜎𝜎 = . + 0.0001 𝐷𝐷 𝑋𝑋 (26)
Where 𝐷𝐷 is in microns. This solved equation is visualized in Figure 9, which shows the yield
stress as a function of 𝑋𝑋 ∗ 𝐷𝐷 (which are both dependent on the material’s processing
parameters). Error bars are included on the plot as Standard Error(s).
Figure 9: The yield stress for material as a function of the processing parameters. A comparison of the model with the experimental data is shown. Note that this modeling equation is applicable to all three initial microstructures.
4. Discussion:
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In this investigation, samples of 4140 steel billet material were processed to contain three
distinct microstructures prior to warm rolling. By varying the warm rolling process parameters
to induce varying levels of effective strain at temperatures between 866K and 973K, the effect of
the starting microstructure on the dynamic grain refinement characteristics was quantified. It was
shown that the initially pearlitic and ferritic microstructure could be transformed to a fully
martensitic microstructure when the cooling rate was on the order of 30 K/s (30 °C/s), while a
predominately bainitic microstructure was obtained when the cooling rate was on the order of
0.83 K/s (0.83 °C/s). Consistent with literature [36], 4140 steel transforms completely to
martensite when cooled from above 1023 K (750 °C) at cooling rates exceeding 16.7 K/s (16.7
°C/s). For a 100% bainitic transformation, a linear cooling rate between 0.083 K/s (0.083 °C/s)
and 0.0083/s (0.0083 °C/s) would be required. When producing the bainitic microstructures,
however, the actual measured cooling rate of approximately 0.83 K/s (0.83 °C/s) resulted in the
formation of 80-85% bainite, with the remainder transforming to martensite. The SEM images
shown in Figure 2 are consistent with the measured cooling rates as the fraction of bainite is
estimated to be approximately 85%, with the remainder being martensitic and retained austenite.
During the warm rolling process, the three initial microstructures behaved differently
with respect to their dynamic recrystallization kinetics. The initially pearlitic and ferritic
microstructure exhibited the most sluggish recrystallization kinetics. Although cementite
precipitates present in pearlite grains effectively pin dislocations, the comparatively large
spacings between precipitates form regions of ferrite where subgrain formation and rotation into
continuously dynamically recrystallized grains may occur. Moreover, within the ferrite grains,
no features exist to pin dislocations, and significant recovery is able to occur. These
microstructural features contribute to increasing the level of effective strain required to produce
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CDRX in those regions as the kinetics of deformation and recovery are in competition. As such,
the final levels of CDRX in these pearlitic/ferritic microstructures tends to be limited for the
range of effective strains evaluated in this study. On the other hand, the bainitic microstructure
contains a fine and homogeneous distribution of carbides throughout the microstructure. The
small interparticle spacings and homogeneous distribution of fine carbides confers a high degree
of dislocation pinning throughout the microstructure. This also serves to limit the degree of
degree of dislocation recovery occurring during deformation since the magnitude of dislocation
travel is limited by the interparticle spacings. This limits the size of the subgrain structures that
can form and promotes relatively high levels of dynamic recrystallization as a function of
effective strain. Finally, the 4140 samples with the martensitic starting microstructures exhibited
the highest tendency to dynamically recrystallize during warm rolling. Heating the martensitic
samples up to the warm rolling temperatures resulted in tempering of the martensite and
precipitates of carbides with extremely fine interparticle spacings. Not only did this
microstructure possess a homogeneous dispersion of carbides to pin dislocations, but the
increased grain boundary surface formed by the interlocking martensitic sheaves also promoted
CDRX during deformation. In addition to forming sub-grain structures that are limited by the
size of the carbide spacings, these elongated sheaves also appear to fragment and form discrete
grains during deformation. The combination of these factors decreased the amount of effective
strain required to dynamically recrystallize the majority of the volume fraction. It should be
noted that despite the difference in starting microstructure, the dynamically recrystallized grains
appeared to be ferritic and the resulting microstructure consisted of ferrite and carbides in all
three cases.
23
As expected, the tensile tests demonstrated that the yield strength of the warm rolled
4140 samples decreased as the warm rolling temperature increased [37]. In addition, the initially
pearlitic and bainitic samples experienced an increase in ductility corresponding with the
decrease in yield strength. Interestingly, the initially bainitic sample rolled at 977 K (704 °C)
exhibited the most tensile ductility out of all the samples tested with 11.7% strain. Following
yield, the plastic flow behavior in this sample was characterized by a lack of both hardening and
softening behavior. The increased ductility associated with this sample could likely be attributed
to the microstructure being comprised of a nearly homogeneous distribution of sub-micron
grains, Figure 5. The observed tensile flow behavior is characteristic of superplastic flow
behavior where macroscopic strains are largely being accommodated via grain boundary sliding
and rotation with minimal dislocation activity. Under these conditions, the ductility of the
microstructures was effectively extended. Although the samples with initially martensitic
microstructures were found to exhibit similar grain sizes and magnitudes of the flow stresses
when compared to the initially bainitic samples, the warm rolled martensitic samples
experienced a decrease in the overall ductility as the rolling temperature increased. It is possible
that the decreases in ductility as a function of temperature could be attributed to temper
embrittlement of 4140. Since molybdenum and chromium are known to react to form carbides
(most effectively at temperatures above 813 K (540 °C) [38], softening of the alloy is mitigated
at all tempering temperatures. In conventional Cr-Mo steels, carbides are often utilized to assist
in preventing grain growth and alloying the steel with chromium and molybdenum additions
retards the coalescence of carbides. However, molybdenum additions are also utilized to mitigate
the effects of temper embrittlement. When Cr-Mo carbides form in alloy steels containing low
levels of molybdenum, such as 4140, the corresponding depletion of the Mo from the ferrite
24
renders the alloy susceptible to temper embrittlement. As a result, 4140 is susceptible to temper
embrittlement when the martensitic structure is tempered at temperatures between 728 K and 868
K (455 °C and 595 °C). Since this coincides with temperature range used for warm rolling and
embrittlement could not be avoided in 4140 at the lower rolling temperatures. In order to remove
the susceptibility for temper embrittlement in this alloy, a slight increase in the molybdenum
content would be required to prevent harmful inclusions from segregating to the grain
boundaries. For example, 4340 steel is not susceptible to temper embrittlement as its
molybdenum content is slightly higher than that of 4140, and its chromium content slightly
lower. As such, the molybdenum is able to sufficiently counteract the effects of the chromium,
while the addition of nickel provides some weak solid solution strengthening [38].
Constitutive models developed to describe the microstructural transformations occurring
via dynamic recrystallization during warm rolling of 4140 steel were observed to be in good
agreement with the experimental data as the effects of both deformation temperature and
effective strain were effectively captured. These models also correctly predict the difference in
behavior associated with varying the initial microstructures. Although some discrepancies in the
models can be observed at higher temperatures, this can likely be attributed to the occurrence of
grain growth. Model predictions for the size of the dynamically recrystallized grains are
consistent with the observations that lower rolling temperatures produce finer dynamically
recrystallized grains. It should be noted that the error bars shown in Figures 7 and 8 were
generated versus each individual trend computed by the global model; the error bars shown may
be larger than expected as a result.
The yield stress model can be observed to be in good agreement with the experimental
data. Assuming complete recrystallization, this model can be used to accurately predict the yield
25
stress of warm rolled 4140 steel containing grain sizes between ~0.2 and 1.0 microns. For grain
sizes above 1 micron, the typical Hall-Petch equation likely be valid. In the yield stress model,
Xdrx is factored in as a correction for the level of recrystallization present in the material.
Although the degree of dynamic recrystallization was limited in the samples possessing the
initially pearlitic and ferritic microstructure due to recrystallization predominately occurring
within the pearlite phase fraction (unless significantly higher effective strains are provided), the
corresponding yield stresses were still comparable to the other two initial microstructures.
Successfully coupling the microstructural process models with the yield stress model is
significant, in that the strength of the material can be directly estimated based on the initial
microstructure and the processing parameters.
5. Conclusions:
This investigation showed that the starting microstructure of 4140 steel plays a major role
in the development of thermal – mechanical processes capable of producing sub-micron grain
sizes. Starting with a tempered martensitic microstructure, warm rolling resulted in recrystallized
grain sizes in the range of 0.5-0.8 microns, and required the least effective strain to achieve high
levels of recrystallization. Warm rolling of bainitic 4140 samples also induced high levels of
recrystallization and had dynamically recrystallized grain sizes slightly larger than that of the
initially martensitic samples, in the range of 0.6-0.9 microns. When the starting microstructure
was predominately comprised of pearlite and ferrite, the resulting warm rolled microstructure
appeared to exhibit a limited level of dynamic recrystallization (~50%) unless significant
effective strain would be supplied at a high rate. Dynamically recrystallized grain sizes for these
samples ranged from 0.7-1.0 microns. Constitutive models were developed that describe the
26
microstructural changes occurring in 4140 steel as a function of deformation temperature,
effective strain and starting microstructure.
The dynamically recrystallized microstructures containing submicron grains were found
to possess yield strengths ranging from approximately 600 MPa up to 1160 MPa following warm
rolling to an overall height reduction of 80%. As expected, the magnitudes of the yield strengths
were observed to increase as the rolling temperature decreased. A modified Hall-Petch model
was developed to estimate the yield stress as a function of grain size. Ductility in these samples
increased with temperature, except in the samples with the martensitic starting structures, where
temper embrittlement likely occurred. It should be noted that the reported yield strength values
were measured from as-rolled material and these strength levels may potentially be enhanced
even further with appropriate heat treatments.
Acknowledgements:
The authors would like to acknowledge the Forging Foundation (FIERF)/Forging Industry
Association and the Armour College of Engineering, a college of the Illinois Institute of
Technology, for providing funding to support this research.
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