Metallurgical and Materials Transactions A

, Volume 38, Issue 3, pp 450–463

The Effect of Multiple Deformations on the Formation of Ultrafine Grained Steels

Authors

    • Centre for Material and Fibre InnovationDeakin University
  • Georgina L. Kelly
    • Centre for Material and Fibre InnovationDeakin University
  • Peter D. Hodgson
    • Centre for Material and Fibre InnovationDeakin University
Article

DOI: 10.1007/s11661-006-9080-7

Cite this article as:
Beladi, H., Kelly, G.L. & Hodgson, P.D. Metall and Mat Trans A (2007) 38: 450. doi:10.1007/s11661-006-9080-7

A C-Mn-Nb-Ti steel was deformed by hot torsion to study ultrafine ferrite formation through dynamic strain-induced transformation (DSIT) in conjunction with air cooling. A systematic study was carried out first to evaluate the effect of deformation temperature and prior austenite grain size on the critical strain for ultrafine ferrite formation (εC,UFF) through single-pass deformation. Then, multiple deformations in the nonrecrystallization region were used to study the effect of thermomechanical parameters (i.e., strain, deformation temperature, etc.) on εC,UFF. The multiple deformations in the nonrecrystallization region significantly reduced εC,UFF, although the total equivalent strain for a given thermomechanical condition was higher than that required in single-pass deformation. The current study on a Ni-30Fe austenitic model alloy revealed that laminar microband structures were the key intragranular defects in the austenite for nucleation of ferrite during the hot torsion test. The microbands were refined and overall misorientation angle distribution increased with a decrease in the deformation temperature for a given thermomechanical processing condition. For nonisothermal multipass deformation, there was some contribution to the formation of high-angle microband boundaries from strains at higher temperature, although the strains were not completely additive.

Introduction

Dynamic strain-induced transformation (DSIT) is one of the routes that has recently been used to produce ultrafine ferrite (UFF) grains in steels. The DSIT mechanism is clearly different from conventional grain refinement practices, which can be viewed as strain-assisted transformation.

In the latter case, the austenite is deformed above the transformation temperature and deformation is stored in the microstructure until the start of transformation (i.e., there is no recrystallization between the deformation and transformation events). Here, the ferrite grain size transformed from the deformed austenite is significantly refined (5 μm) compared with the ferrite grains transformed from the undeformed austenite, with the role of the deformation being the introduction of new nucleation sites.[1] The important feature here is that this is a static process and the deformation sets up the necessary conditions, but does not directly interact with the transformation.

In the case of DSIT, most of the work to date suggests that the ferrite refinement can be further enhanced through the combined effect of deformation and transformation (1 to 3 μm), i.e., a dynamic process.[26] Here, the steel is deformed at temperatures significantly below the Ae3 (i.e., equilibrium temperature of austenite to ferrite transformation) and above the Ar3 (i.e., the empirical temperature of the austenite to ferrite transformation during continuous cooling), and in most cases, it is clear that the transformation occurs during deformation.

There are some serious limitations that prevent mass production of UFF steels, although some reports are available that indicate UFF strip steels have been produced with pilot scale rolling.[6] One of the most important issues is the strain required for UFF formation through DSIT, because this is quite high compared with other thermomechanical processes (i.e., conventional controlled rolling). According to recent studies,[2,4,7] even with the addition of shear, this requires at least 40 pct reduction for a single rolling pass, although this could be affected by steel composition.[8] Furthermore, the critical strain for UFF formation (i.e., εC,UFF) varies between 1.5 and 3 in hot torsion testing and depends on the thermomechanical parameters.[9]

One proposal to overcome this barrier is to accumulate strain through multiple deformations in the nonrecrystallization region and then to use a lower reduction in the intercritical region to activate the dynamic process. This requires an appropriate thermomechanical schedule designed to avoid any recrystallization between passes (i.e., the Ae3-Tnr region, where Tnr is the nonrecrystallization temperature) and to accumulate the strain before the final deformation at, or just above, the Ar3. Essentially, the concept is to precondition the austenite in the nonrecrystallization region to establish potential nucleation sites. The final deformation is then applied to activate the dynamic transformation process.

There are some factors, though, that are unclear and that could potentially limit this approach. In particular, it is possible that deformation in the nonrecrystallization temperature range may not set up the intragranular high-angle defects to act as nucleation sites.[10] The effect of thermomechanical parameters on the nature of the deformed structure, therefore, was also evaluated using a Ni-30 wt pct Fe austenitic model alloy. This allowed a wider range of temperature and deformation conditions to be studied, because this alloy is austenitic to room temperature.

Experimental procedure

A hot torsion simulator with an induction furnace and controlled cooling facilities was employed to conduct the experiments. Specimens with a gage length of 20 mm and gauge diameter of 6.7 mm were used. The sample was enclosed in a quartz tube with a positive pressure of argon gas to prevent decarburization. The maximum temperature difference between the two ends of the sample was 10 °C. Equivalent true stress-true strain values were calculated from the torque-twist data using the method based on the analysis by Fields and Backofen.[11] The microstructural study was made on tangential sections at a depth of ∼100 μm below the surface of the gage length. The grain size was measured using the intercept method.

Samples for electron backscattered diffraction (EBSD) were prepared by mechanical polishing and finished with a colloidal silica slurry polish. The EBSD was employed to measure the size and volume fraction of recrystallized grains as well as the misorientation angles between adjacent grains and subgrains using HKL Technology Channel 5 (Denmark). The EBSD images were constructed either using a step size of 0.05 μm for characterization of intragranular defects or 0.7 μm to study the overall microstructure (i.e., ∼20 grains). Backscattered imaging was also used to reveal the intragranular defect structure.

Two sets of experiments were carried out to study the effect of thermomechanical parameters not only on ferrite refinement but also on the nature of deformation features.

Effect of Thermomechanical Parameters on Ferrite Refinement

A 0.12 pct C-1.14 pct Mn-0.24 pct Si-0.027 pct Al-0.022 pct Nb-0.005 pct V-0.012 pct Ti (in wt pct) steel was used for the multiple deformation schedule. To design an appropriate thermomechanical schedule, the austenite grain structure was first studied over a wide range of austenitization temperatures. Then, the Tnr and Ar3 were determined for two prior austenite grain sizes.

Austenitization

Specimens with dimensions of 10 × 10 × 5 mm were reheated at a range of temperature between 900 °C and 1300 °C for 15 minutes in a muffle furnace, followed by immediate water quenching. The austenite grain boundaries were revealed using a picric acid-based etchant. As will be discussed in detail later, the temperatures of 1000 °C and 1200 °C were chosen to study the effect of thermomechanical parameters on ferrite refinement. The austenite grain size was 14 and 71 μm at 1000 °C and 1200 °C, respectively.

Tnr determination

A multipass torsion test was carried out during cooling (1 °C/s) to determine the Tnr temperature for both sizes of prior austenite grains (i.e., 14 and 71 μm). This follows the method developed by Jonas and co-workers.[12] Pass strains of 0.15 and a strain rate of 1 s−1 were applied. Sixteen and ten passes were performed for coarse and fine prior austenite grains, respectively. The different number of passes is due to the different reheating temperatures. The temperature decreased from pass to pass by ∼20 °C.

Thermomechanical processing

The schedule involved austenitization at either 1200 °C or 1000 °C for 5 minutes, which is comparable to reheating the sample for 15 minutes using a muffle furnace, resulting in austenite grain sizes of 71 and 14 μm, respectively. Specimens were then controlled cooled at a rate of 1 °C/s to a given deformation temperature (i.e., in the range of 775 °C to 900 °C) followed by single-pass deformation at a true strain rate of 1 s−1. The strain varied between 0 and 4 followed by air cooling to determine εC,UFF.

To assess the impact of multiple deformations in the nonrecrystallization region on εC,UFF, two deformation temperatures were selected between the Tnr and Ae3 temperatures (i.e., 900 °C and 860 °C). The Ae3 temperature was calculated to be 854 °C using the Chemsage program. The pass strain was either 0.5 or 1 for each deformation temperature in the nonrecrystallization region. The final strain varied between 0.5 and 2.5 at the deformation temperature of 800 °C (i.e., between the Ae3 and Ar3 temperatures). The Ar3 temperature was determined using the continuous cooling deformation method for both initial austenite grain sizes.[13] The Ar3 was 793 °C and 769 °C for steel with a prior austenite grain size of 14 and 71 μm, respectively.

Softening fraction determination

Two sets of multipass deformation were carried out in the current work: isothermal multipass deformation and nonisothermal multipass deformation. Hence, the estimation of softening fraction was different for each set of experiments.

Isothermal multipass deformation
The fractional softening (Xs) before the last strain in the isothermal multipass test was calculated using Eq. [1].[1] The yield stress, \( \sigma _{1} \), of the steel was taken from the single pass curve. The stress,\( \sigma _{2} \), was also taken from the single pass curve at a strain equal to the total of the preceding strains in the multipass test. The \( \sigma _{3} \) value was taken to be the yield stress in the last deformation of the multipass tests. This is shown schematically in Figure 1(a).
$$ {\rm X}_{s} = \frac{{\sigma _{3} - \sigma _{1} }} {{\sigma _{3} - \sigma _{2} }} \times 100 $$
(1)
where
https://static-content.springer.com/image/art%3A10.1007%2Fs11661-006-9080-7/MediaObjects/11661_2006_9080_Fig1_HTML.gif
Fig. 1

Schematic representation of the determination of the softening fraction in different multipass tests using stress-strain curves from multipass tests and single-pass tests at equivalent finishing temperatures; T and σ correspond to the deformation temperature (T1 > T2 > T3) and stress, respectively: (a) isothermal deformation and (b) nonisothermal deformation (σ1 and σ2 are the yield stress of the first pass and final pass, respectively, and σ3 is the stress for total strain to the last pass)

\( \sigma _{1} \)= the yield stress of the final pass of the multipass deformation,

\( \sigma _{2} \)= the yield stress of the single pass at the finishing temperature, and

\( \sigma _{3} \)= the stress at the finishing temperature for the total strain to the last pass.

Nonisothermal multipass deformation
The deformation temperatures were different in each pass for nonisothermal deformation. Therefore, the fractional softening (X′s) during the multiple deformations (in the Tnr-Ae3 region) and before the final strain at 800 °C (in the Ae3-Ar3 region) could not be calculated by comparing the flow curves obtained from multiple deformations with the flow curve of final strain at 800 °C using Eq. [1]. Hence, a total equivalent strain (i.e., sum of strains at 900 °C and 860 °C) was applied at 800 °C. Then, the total equivalent stress-strain curve at a given thermomechanical processing condition was considered to measure \( \sigma ^{'}_{2} \)and \( \sigma ^{'}_{3} \)values (Figure 1(b) and Eq. [2]). It is worth mentioning that there is a possibility of dynamic strain-induced ferrite transformation at 800 °C when the total equivalent strain is applied. The dynamic strain-induced transformation of ferrite could affect the estimation of the softening fraction. This was, however, ignored in the current study.
$$ {\rm X}^{'}_{s} = \frac{{\sigma ^{'}_{3} - \sigma ^{'}_{2} }} {{\sigma ^{'}_{3} - \sigma ^{'}_{1} }} \times 100 $$
(2)

where

\( \sigma ^{'}_{1} \)= the yield stress of the final pulse at 800 °C,

\( \sigma ^{'}_{2} \)= the yield stress of the total equivalent strain at 800 °C, and

\( \sigma ^{'}_{3} \)= the stress at 800 °C for the total equivalent strain.

Effect of Thermomechanical Parameters on the Nature of Deformation Features

A 70Ni-30 wt pct Fe austenitic model alloy was used to study the nature of deformation features at a given thermomechanical condition. Because this alloy retains a stable austenitic structure after cooling to room temperature, it allows the high-temperature deformation microstructures to be readily characterized. In addition, this alloy has a stacking fault energy similar to that of austenitic iron at elevated temperatures.[14] Therefore, it is expected to deform in a similar fashion to steel during hot deformation. Hurley et al.[4,15,16] and Adachi et al.[10] have previously used this alloy to monitor the evolution of ferrite nucleation sites during hot deformation. They have shown a strong agreement between the deformed structure in the model alloy and ferrite nucleation sites at an early stage of transformation for hot compression, torsion, and rolling deformation techniques.

This set of experiments was undertaken to determine the effect of strain on the deformed structure characteristics at deformation temperatures between 700 °C and 900 °C. Specimens were initially reheated at 1200 °C for 2 minutes followed by two strains of 0.4 with an interpass time of 40 seconds (i.e., roughing stage) resulting in an equiaxed recrystallized grain size of 110 μm. The specimens were then cooled at 1 °C/s to a given deformation temperature, held for 2 minutes to allow the temperature to stabilize, and deformed to strains ranging from 0 to 1 at a strain rate of 1 s−1. Once the deformation ceased, the samples were immediately water quenched. Backscattered imaging was performed to reveal the structure of the intragranular defects supported by EBSD analysis to study the deformation characteristics in more detail.

Results

Austenitization

The austenite grain size increased with an increase in the reheating temperature (Figure 2). The rate of austenite grain growth changed at different temperature ranges. Two sharp deviations were seen at temperatures of 1000 °C and 1200 °C. Abnormal grain growth was observed at soaking temperatures of 1050 °C and 1250 °C. This suggested that the grain growth behavior was controlled by particles with different thermal stabilities. The austenite grains were homogenous at 1000 °C and 1200 °C. Therefore, these temperatures were selected to study the effect of prior austenite grain sizes on ferrite refinement.
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Fig. 2

Austenite grain size as a function of reheating temperature

The steel composition shows that this steel was killed by Al during steel making. The first deviation can be attributed to AlN particles, which start to be dissolved above 1000 °C. Abnormal grain growth was also observed at a soaking temperature of 1050 °C. With increasing temperature, the AlN precipitates completely dissolved and the grain size distribution became homogenous. The grain growth at temperatures of 1000 °C to 1200 °C, therefore, was controlled by other particles, which have higher thermal stability than that of AlN. According to the composition of the steel, the second deviation is likely to be due to Nb(C,N) or TiN precipitates.

Killmore et al.[17] evaluated the austenite grain growth in a steel composition similar to the current steel. However, the current result is different from Killmore et al.’s result, particularly at temperatures higher than 1200 °C (Figure 2). This could be caused by differences in steel processing, which affect the size and distribution of precipitates in the initial steel microstructure.[18]

Tnr Determination

The multipass torsion test showed that the flow stress increased from pass to pass with decreasing temperature for both prior austenite grains. There were also two changes in the rate of flow stress increase for both prior austenite grains (indicated by SDE and SIP in Figure 3). For example, in Figure 3(a), there was strain accumulation between the deformations at 919 °C and 900 °C, while there was substantial softening between 900 °C and 880 °C. Below this temperature, there was little interdeformation softening. These fluctuations could be explained by the solute drag effect of Nb (SDE)[12,19] and the pinning effect of strain-induced Nb(C,N) precipitates (SIP),[12,19] respectively. The mean flow stress (MFS) for each pass was calculated from the flow curves using Eq. [3].
$$ \sigma _{{{\text{MFS}}}} = \frac{1} {{\varepsilon _{{i + 1}} - \varepsilon _{i} }}{\int_{\varepsilon _{i} }^{\varepsilon _{{i + 1}} } {\overline{\sigma } d\varepsilon } } $$
(3)
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Fig. 3

Stress-strain curves corresponding to multipass torsion tests for different prior austenite grains: (a) l4 μm and (b) 7l μm. The temperatures of the passes are shown in the figure. (SDE and SIP correspond to the solute drag effect and strain-induced precipitation effect, respectively.)

There was a slope change in MFS as a function of inverse temperature, which allowed the graph to be divided into two regions for both prior austenite grains (Figure 4). In region I (which corresponds to high-temperature deformation), full recrystallization took place because there was no strain accumulation. The increase in MFS was solely due to the decrease in temperature. In region II (which corresponds to deformation below the Tnr), there was only partial recrystallization or no recrystallization at all. Here, the strain accumulated from pass to pass, so that the MFS increased more rapidly with decreasing temperature. The retardation of recrystallization first occurred by solute drag at higher temperatures and was then completely halted by strain-induced precipitation, which was clearly shown in region II (Figure 4).
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Fig. 4

Dependence of the MFS on the inverse absolute temperature during a multipass torsion test of the steel for different prior austenite grains: (a) l4 μm and (b) 7l μm

The Tnr temperature was determined to be 910 °C and 950 °C for 14- and 71-μm prior austenite grains, respectively. A dramatic decrease in the Tnr with a decrease in prior austenite grain size was expected due to the lower Nb in solution with a decrease in the austenitization temperature, as described by Irvine et al.[20] The finer austenite grain size also accelerates the rate of recrystallization.[12] Therefore, deformation temperatures of 900 °C and 860 °C were chosen to study the effect of multiple deformations in the nonrecrystallization region on the critical strain for UFF formation. These were also the temperatures where there was clear strain accumulation in Figure 3.

Effect of Thermomechanical Parameters on Ferrite Refinement

Single-pass deformation

The ferrite refinement was strongly affected by the deformation temperature for a given thermomechanical condition in the nonrecrystallization region. The level of refinement that could be achieved at a strain of 3 increased with decreasing deformation temperature. In addition, the ferrite refinement in the Ae3-Ar3 region was much greater compared with refinement resulting from deformation in the Tnr-Ae3 region (Figure 5). The ferrite grain size was reduced by a factor of ∼2 in the Ae3-Ar3 region at the strain of 3.
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Fig. 5

Ferrite grain size as a function of strain at different deformation temperatures in the nonrecrystallization region; W is Widmanstätten ferrite, and Td is the deformation temperature

The ferrite grain size decreased with an increase in the strain in the Ae3-Ar3 region for both initial austenite grain sizes (Figure 5). At strains lower than 1, there was nonpolygonal ferrite (Widmanstätten ferrite) in the air-cooled microstructures. This formed during cooling in the coarse-grained austenite (indicated by “W” in Figure 5) after deformation at 800 °C and 825 °C (Figures 6(a) and (b)). This disappeared beyond a strain of 1, and the microstructure consisted of polygonal ferrite, fine carbides, and fine pearlite colonies (Figures 6(c) through (f)). The ferrite grain size was inhomogeneous and it became more uniform with increasing strain. There were regions of very fine ferrite grains (less than 2 μm) where the carbides were mainly observed (Figure 6(d)). The pearlite colonies were mostly located among the coarser ferrite grains (greater than 2 μm). The volume fraction of fine grains increased with the strain, and their distribution became more homogenous. It also appeared that the volume fraction of fine ferrite grains increased with a decrease in deformation temperature in the Ae3-Ar3 region (Figure 6).
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Fig. 6

Air-cooled microstructure steel after deformation at different temperatures and strains. The strain rate and austenitization temperature were 1 s–1 and 1200 °C, respectively: (a), (c), and (e) at Td = 825 °C and (b), (d), and (f) at Td = 800 °C. Parallel arrows represent the deformation direction; Td is the deformation temperature

A mean ferrite grain size of less than 3 μm was considered as an ultrafine ferrite microstructure in this study. The minimum strain required to produce UFF (εC,UFF) was 2.8 for a deformation temperature of 800 °C for the steel with coarse prior austenite grains (Figure 5 and Table I). At strains higher than εC,UFF, the microstructure did not change markedly and there was still inhomogenity in the distribution of the ferrite grain size, carbides, and pearlite colonies (Figure 6(e)). No critical strain for UFF formation was detected at 825 °C. The mean ferrite grain size was greater than 3 μm at 825 °C, even after a strain of 4, which was the limiting strain as some cracks appeared on the gauge length of the sample (Figure 5). At 775 °C, there were work-hardened ferrite grains in the microstructure. This indicated that deformation occurred in the two-phase region. Therefore, the processing window for UFF formation through DSIT in conjunction with air cooling for this steel requires a deformation temperature between 825 °C and 775 °C.
Table I

Effect of Thermomechanical Parameters on εC,UFF*

Ta (°C)

Td1 (°C)

ε1

Td2 (°C)

ε2

Td3 (°C)

εC,UFF

εeq

XS (Pct)

1200

900

no UFF

1200

860

no UFF

1200

825

no UFF

1200

800

2.8

2.8

1200

775

no UFF

1000

800

2.5

2.5

1000

900

0.5

860

0.5

800

2

3

48

1000

900

1

860

1

800

1.5

3.5

77

1200

900

1

860

1

800

1

3

18

*Ta and Td correspond to austenitization temperature and deformation temperature, respectively

The εC,UFF was slightly reduced from 2.8 to 2.5 with a decrease in the prior austenite grain size from 71 to 14 μm at a deformation temperature of 800 °C (Figure 5 and Table I). For the fine prior austenite grain, the UFF microstructure was more uniform than for the coarse initial grain size (Figure 7).
https://static-content.springer.com/image/art%3A10.1007%2Fs11661-006-9080-7/MediaObjects/11661_2006_9080_Fig7_HTML.gif
Fig. 7

Air-cooled microstructure after deformation at 800 °C for different austenitization temperatures and strains: (a) Ta = 1200 °C, ε = 2.8; and (b) Ta = 1000 °C, ε = 2.5. Parallel arrows represent the deformation direction (Tais the austenitization temperature)

Multiple deformation

Multiple deformation in the nonrecrystallization region significantly reduced the strain required for UFF formation at 800 °C (εC,UFF), although the total equivalent strain (εeq) for the total thermomechanical processing below the Tnr (i.e., sum of strains at 900 °C, 860 °C, and 800 °C) was higher than that of single-pass deformation (Figure 8 and Table I). The fractional softening between passes was calculated using Eq. [2]. The result showed that there was softening between multiple deformations and the final deformation pass (i.e., during controlled cooling). In addition, the softening fraction (Xs) was significantly affected by the prior austenite grain and strain value of the pass (Table I). The results also suggested that the coarse prior austenite grain size has a greater effect on εC,UFF through multiple deformation in the nonrecrystallization region compared with the fine prior austenite grain size at a given pass strain in the nonrecrystallization region.
https://static-content.springer.com/image/art%3A10.1007%2Fs11661-006-9080-7/MediaObjects/11661_2006_9080_Fig8_HTML.gif
Fig. 8

Ferrite grain size as a function of the final strain at 800 °C for different prior austenite grain sizes and different deformed conditions (pass strain). Two pass strains were applied at 900 °C and 860 °C

Effect of Thermomechanical Parameters on the Nature of Intragranular Defects

The initial microstructure of the model alloy prior to deformation consisted of a coarse-grained structure containing annealing twins. The initial grain size was 110 μm. Stress-strain curves showed no clear peak for deformation temperatures between 700 °C and 900 °C at 1 s−1 (Figure 9). The curves had a period of clear power-law-type work hardening followed by an apparent linear behavior. The maximum stress increased significantly with a decrease in the deformation temperature from 900 °C to 700 °C (Table II).
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Fig. 9

Stress-strain curves of the model alloy at different deformation temperatures; εc is the critical strain for dynamic recrystallization

Table II

Peak Strain (εp) and Critical Strain for Dynamic Recrystallization (εC,DRX) at Different Deformation Temperatures (Td)

Td

εp

εC,DRX

εC,DRX/εp

900 °C

0.58

0.28

0.48

800 °C

0.64

0.44

0.69

700 °C

1

0.75

0.75

The critical strain for dynamic recrystallization (εC) was determined using an approach recently presented by Jonas et al.[21] (Figure 9 and Table II). The strain of the peak stress was termed the peak strain (εP) in this study. The critical strain for dynamic recrystallization and peak strain sharply increased with a decrease in the deformation temperature from 1000 °C to 700 °C. In fact, a strain of 1 was greater than the critical strain for dynamic recrystallization for all deformation temperatures.

The as-quenched microstructures at different deformation temperatures also revealed that new recrystallized grains formed on pre-existing grain boundaries during deformation (indicated by the arrows in Figure 10). The volume fraction of recrystallized grains and their size increased with an increase in deformation temperature. In addition, twin boundaries were more serrated at higher deformation temperatures for a given thermomechanical condition.
https://static-content.springer.com/image/art%3A10.1007%2Fs11661-006-9080-7/MediaObjects/11661_2006_9080_Fig10_HTML.jpg
Fig. 10

As-quenched microstructures of Ni-30Fe alloy at a strain of 1 for different deformation temperatures: (a) 700 °C, (b) 800 °C, and (c) 900 °C. Examples of new, recrystallized grains are indicated by the arrows. Parallel arrows represent the macroscopic shear. (TW represents annealing twin.)

The main aim of this set of experiments was to study the development of the intragranular defects in the austenite during thermomechanical processing. Backscattered electron imaging allowed the intragranular deformed structure for the different thermomechanical processing conditions to be more clearly seen. Parallel, bandlike structures were observed extensively within individual austenite grains throughout the microstructure for all deformation temperatures (Figure 11). This structure, known as microbands[22] or geometrically necessary boundaries,[23] formed at an angle of up to 45 deg to the direction of macroscopic shear, depending on parent grain crystallography orientation. The microbands were refined with a decrease in the deformation temperature for a given set of thermomechanical processing parameters (Figure 12).
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Fig. 11

Microstructure of a grain interior of the model alloy at different deformation temperatures: (a), (c), and (e) backscattered electron images at deformation temperatures of 900 °C, 800 °C, and 700 °C, respectively. (b), (d), and (f) EBSD images at deformation temperatures of 900 °C, 800 °C, and 700 °C, respectively. Parallel arrows represent the macroscopic shear; A and B are the higher magnification of two areas in (f) shown in detail with higher angle boundaries highlighted. The black and red lines represent low-angle (<15 deg) and high-angle (>15 deg) boundaries, respectively

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Fig. 12

Mean microband thickness as a function of deformation temperature at a strain of 1 and a strain rate of 1 s–1 for the Ni-3OFe alloy

The EBSD was also carried out on the same area on which backscattered electron imaging was performed. At a deformation temperature of 900 °C, the misorientation angle of the intragranular structure was less than 3 deg over the entire microstructure (Figures 11 and 13). However, this significantly increased as the deformation temperature decreased. Some microband boundaries were observed with misorientation angles as high as 16 and 18 deg at 800 °C and 700 °C, respectively (Figure 11). The misorientation angles across microband boundaries also increased with an increase in the strain at a given deformation temperature (Figure 14). This suggested that the deformation temperature and strain had a strong effect on the nature of the intragranular deformed structure.
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Fig. 13

Misorientation angle distribution as a function of deformation temperature at a strain of 1 for the Ni-30Fe alloy

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Fig. 14

Misorientation angle distribution as a function of strain at different deformation temperatures for the Ni-30Fe alloy: (a) 700 °C, (b) 800 °C, and (c) 900 °C

Discussion

Effect of Thermomechanical Parameters on Ferrite Refinement

Controlled rolling, where the steel is deformed in the nonrecrystallization region (i.e., lower than Tnr temperature), is known as the most effective industrial method for ferrite refinement from deformation. The deformation probably has a minor effect on the free energy for the ferrite transformation and the diffusion of elements. However, the main effect of deformation in the nonrecrystallization region is the introduction of serrations in the austenite grain boundary (Figure 15(a)-ii), which then act as nucleation sites for the transformation. The retained strain (i.e., the strain accumulated below the Tnr) enhances the number of nuclei per unit length of grain boundary (Sv). The deformation can also activate intragranular nucleation sites (i.e., deformation bands,[23,24,25] twin bands,[26] and dislocation arrays[27]), and furthermore, there may be an indirect geometry effect due to the reduction in the distance between adjacent austenite grain boundaries (Figure 15). The large flattening of the austenite grains through deformation without recrystallization means that separation between these boundaries is reduced, and it is possible that soft impingement of the ferrite growing in from each of the boundaries may reduce the growth of ferrite (Figure 15). Controlled rolling, therefore, can be seen as a strain-assisted transformation, but where the effect of strain is more from the geometry that is created through the flattening and roughening of the boundaries.
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Fig. 15

Schematic presentation of austenite grain behavior during straining in the nonrecrystallization region followed by transformation: (a) strain-assisted transformation (Ae3 < Td < Tnr) and (b) dynamic strain-induced transformation (Ar3 < Td < Ae3). (i) A strain-free austenite grain, (ii) roughening the austenite grain boundaries, (iii) intragranular defect formation and flattening of the austenite grain, and (iv) a temperature lower than Ar3 (γ and α are austenite and ferrite grains, respectively, and Td is the deformation temperature)

The current results reveal that the level of ferrite refinement in the Tnr-Ae3 region increases with a decrease in the deformation temperature. For example, for the steel with coarse prior austenite grain, the ferrite grain size reduces from 6.5 to 4.6 μm at a strain of 3 with a decrease in deformation temperature from 900 °C to 860 °C (Figure 5). Isothermal multipass deformations at 900 °C and 860 °C show that there is softening between passes (Figure 16). The deformation temperature has a strong effect on the softening fraction. Recrystallization can occur at a deformation temperature of 900 °C beyond a strain of 2 and during cooling before transformation takes place (Figure 16), which significantly decreases the retained strain. The recrystallized grains remove the grain boundary steps, intragranular defects, and flattened austenite grain structure. However, it would be expected that a small amount of Ti in the steel (i.e., 0.012 wt pct) could effectively reduce the growth of recrystallized grains. Hence, the surface area of austenite per unit volume also increases at 900 °C through recrystallization. However, the increase in the surface area of austenite per unit volume through recrystallization is much smaller than for controlled rolling, and this leads to a lower level of ferrite grain refinement when only recrystallization occurs.
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Fig. 16

Effect of deformation temperature on the softening fraction through isothermal multipass deformation of the steel using a pass strain of 0.3 and an interpass time of 1 s

There is still a ferrite grain size limit (∼5 μm) through deformation without recrystallization in the nonrecrystallization region even at a temperature just above the Ae3 (i.e., 860 °C). It is more likely that recovery takes place here during cooling before transformation occurs (Figure 16). Prior to the 1990s, it was believed that the recovery during controlled cooling before transformation could have a major effect on the potential of intragranular defects for ferrite nucleation.[28,29,30] On the other hand, it has been shown that, in the absence of recrystallization, there will be no change in Sv even after softening due to recovery.[1] Those authors reported a ferrite grain size limit of 5 μm,[1] similar to the current result. It was shown later that there is a coarsening of ferrite grains during postdeformation cooling, which significantly limits ferrite grain refinement through controlled rolling.[31]

The level of ferrite refinement significantly increases when the deformation is applied in the Ae3-Ar3 region where dynamic-strain-induced transformation occurs. However, there is a significant inhomogenity in the ferrite grain size distribution, and it becomes more uniform with strain (Figure 6). The difference in ferrite grain size and distribution of carbides and pearlite colonies suggests that there is a different mechanism for formation of these phases during deformation and postdeformation cooling.

The fine grains (less than 2 μm) represent ferrite grains induced dynamically during deformation in the Ae3-Ar3 region. They mainly have three-dimensional impingement. The full impingement of DSIT grains at the early stage of transformation controls the coarsening of ferrite during postdeformation cooling.[32] This led to the formation of raftlike regions along the shear direction for the air-cooled conditions (Figure 6).

The austenite did not fully transform to ferrite during deformation and the volume fraction of DSIT ferrite depended upon strain. Therefore, the ferrite grains also statically formed during postdeformation cooling. There is a lack of full impingement in these statically formed ferrite grains that leads to coarsening in this region. This resulted in coarser ferrite grains and inhomogenity in the air-cooled microstructure (Figure 6). With an increase in the strain, the volume fraction of DSIT ferrite grains increased. Therefore, DSIT ferrite grains were uniformly distributed at the early stage of transformation resulting in a more homogenous air-cooled microstructure. The distribution of DSIT ferrite grains can also be promoted by a decrease in prior austenite grain size, because the fine austenite grain promotes a more homogenous strain distribution (Figure 6).

The current result reveals that deformation in the Ae3-Ar3 region is necessary to dynamically induce ferrite formation during deformation, but this is not a sufficient factor for UFF formation. The volume fraction of DSIT ferrite decreases with an increase in the deformation temperature in the Ae3-Ar3 region at a given strain resulting in a greater coarsening rate. The ferrite coarsening kinetics also depend upon the temperature and rapidly increase with an increase in temperature. Hence, the deformation temperature range for UFF formation could be significantly limited based on thermomechanical parameters. The current results suggest that the deformation temperature should be close to 800 °C to produce an UFF microstructure. For instance, no critical strain for UFF formation was detected at 825 °C (i.e., ferrite grain size >3 μm), even at a strain of 4 at which some cracks appeared on the gauge length (Table I). However, the cooling below 800 °C is also limited, because, at 775 °C, the material had already partially transformed to ferrite prior to deformation.

Effect of Multiple Deformations in the Nonrecrystallization Region on εC,UFF

The current results confirm that εC,UFF is very high for single-pass deformation (strain of >2.5 for both initial austenite grain sizes). However, it is possible to significantly reduce the strain required in a single pass for UFF formation (εC,UFF) through multiple deformations in the nonrecrystallization region. However, the total equivalent strain (εeq) in multiple deformations is still higher than that of εC,UFF in a single-pass deformation, although this could be reduced by optimizing the thermomechanical parameters (Figure 8 and Table I).

The enhancement of nucleation sites in deformed austenite is the main reason for the remarkable refinement in ferrite grain size for most thermomechanical processes. Therefore, it could be proposed that the nature of the intragranular defects has been altered through partial deformation in the nonrecrystallization region or the potential of intragranular defects has reduced during controlled cooling (before final straining) compared with single-pass deformation.

The current results show that there is a significant softening (Xs) between multiple deformations and the final pass deformation (during controlled cooling), although this is strongly affected by pass strain value and prior austenite grain size (or Nb in solution) (Figure 8 and Table I). For a fine prior austenite grain, the softening fraction is much higher than 20 pct. This indicates the strong possibility that static recrystallization has occurred before the final deformation at 800 °C (Table I). This suggests that new recrystallized austenite grains removed the deformed structures (i.e., serrated boundaries and intragranular defects) formed during multiple deformations. Although these recrystallized grains increase the surface area of austenite per unit volume, it is much smaller compared with the fully deformed structure at a given condition. This, therefore, leads to a lower level of ferrite refinement. Consequently, no significant effect is seen on εC,UFF through multipass deformations for the finer prior austenite grain size (Figure 8 and Table I).

In the samples with a coarse prior austenite grain size, there is, however, a remarkable reduction in εC,UFF, although the total equivalent strain (εeq) is slightly higher than εC,UFF for single-pass deformation (Figure 8 and Table I). The interdeformation softening fraction (i.e., Xs ∼18 pct) suggests that recovery is more likely to have occurred during cooling. This seems to slightly affect the nucleation site potential before the final deformation. In fact, multiple deformations in the nonrecrystallization region mainly increase the surface area of austenite per unit volume (i.e., ferrite nucleation sites) for the coarse austenite grains, and the final strain will mostly be employed to induce the dynamic transformation rather than the introduction of ferrite nucleation sites. This suggests that the coarse austenite grains are more effective in reducing εC,UFF through multiple deformations in the nonrecrystallization region compared with the fine prior austenite grains. However, as noted elsewhere, this effect also may be due to the lower Nb, and it may be possible to achieve a similar benefit if we combine a fine austenite grain with high Nb.

Effect of Thermomechanical Parameters on Substructure Formation in the Ni-30Fe Austenite

In the model alloy, the strain introduces serrations in the grain boundaries at an early stage of transformation. The twin grain boundaries also lose their coherency with increasing strain for a given thermomechanical condition. These geometry changes in grain and twin boundaries are appropriate sites for initiation of recrystallization beyond the critical strain for dynamic recrystallization (Figure 10). In the absence of recrystallization, they would also be favourite sites for ferrite nucleation during austenite to ferrite transformation in a steel with lower alloying.

The strain also introduces microband structures in the grain interior. There is a misorientation angle across the microband boundaries at a given strain. The misorientation angle represents the dislocation density (i.e., elastic strain energy) in this region. In other words, a higher misorientation angle indicates greater elastic strain energy. As strain is increased, the misorientation angles across microband boundaries increase to accommodate the strain (Figure 14). This is similar to other observations by Hughes and Nix[22] and Hurley.[16] This would result in a gradual increase in the elastic strain energy within the microband boundaries. These regions could reach a sufficient energy level to become active for a transformation event (i.e., recrystallization[33] or ferrite transformation[16]). For ferrite transformation, the undercooling is a further necessary parameter to activate these sites for nucleation. The high misorientation angle areas (e.g., >10 deg) will require a lower level of undercooling for ferrite nucleation compared with low misorientation angle areas (e.g., <3 deg). This is consistent with the effect of strain on ferrite grain size in both controlled rolling and DSIT (Figure 5).

For this set of experiments, the laminar microband structures are observed for all thermomechanical processing conditions. This suggests that the intragranular defects in austenite that act as nucleation sites for ferrite during hot torsion are more likely to be microbands (i.e., >10 deg) than other deformation features such as microshear bands. This is similar to findings of another research group using hot torsion[16] and hot compression testing.[10]

The results also reveal that the deformation temperature has a strong effect on the microband characteristics. At a given strain, the microbands refine (Figure 12) and the overall misorientation angle distribution increases (Figure 13) with a decrease in the deformation temperature. This suggests that there is a critical temperature region where the dislocation structure changes significantly for a given strain. The current observation indicates that this critical temperature region (i.e., <800 °C) is consistent with that used for the DSIT route. This could be another explanation for increasing ferrite refinement at lower deformation temperatures (Figure 5).

A multiple deformation experiment was also carried out to assess the impact of deformation temperature on the nature of the intragranular defects. A strain of 0.5 was applied at 900 °C followed by rapid cooling to 700 °C to avoid any recrystallization. Then, another strain of 0.5 was applied at 700 °C and the specimen immediately water quenched. The EBSD result reveals that there is some contribution from strains at higher temperature to the formation of these higher angle microband boundaries in the multipass case, although it would appear that the strains are not completely additive (Figure 17). Very few high-angle boundaries were evident after a strain of 1 at 900 °C (Figure 13). However, a strain of 0.5 applied at 900 °C then followed by another strain of 0.5 at 700 °C resulted in more high-angle boundaries than a single strain of 0.5 applied at 700 °C (Figure 17). The strain applied at the lower temperature is much more effective at generating the higher angle boundaries. This suggests that the increase in the total equivalent strain (εeq) for UFF formation during multiple deformations for initially coarse austenite is partly because of the deformation temperature effect on intragranular defect characteristics apart from recovery.
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Fig. 17

Misorientation angle distribution as a function of strain for different thermomechanical processing conditions for the Ni-30Fe alloy

Summary

  1. 1.

    The ferrite grain refinement for deformation in the Ae3-Ar3 region was significantly greater compared with deformation in the Tnr-Ae3 region. The ferrite grain size was reduced by factor of ∼2 in the Ae3-Ar3 region at a strain of 3.

     
  2. 2.

    The εC,UFF value was large for UFF formation through single-pass deformation in the Ae3-Ar3 region for both initial austenite grain sizes (>2.5), although it could be affected by deformation temperature as well as prior austenite grain size.

     
  3. 3.

    It was possible to significantly reduce εC,UFF (to ∼1) through multiple deformation in the nonrecrystallization region, although the total equivalent strain (εeq) was higher than that for single-pass deformation.

     
  4. 4.

    The current results revealed that the laminar microband structures were the main intragranular feature at all thermomechanical processing conditions using hot torsion testing. This suggested that the key intragranular defects in austenite for nucleation of ferrite during hot torsion test are more likely to be microbands rather than other deformation features such as microshear bands.

     
  5. 5.

    The microbands refined and overall misorientation angle distribution increased with a decrease in deformation temperature for a given thermomechanical processing condition. This suggested that there is a critical temperature region where the generation of the dislocation structure changes remarkably for a given strain.

     
  6. 6.

    A nonisothermal multipass deformation of the model alloy revealed that there is some contribution from strains at higher temperature to the formation of high-angle microband boundaries in the multipass case. However, the strains are not completely additive.

     

Acknowledgments

This research has been supported by grants through the Australian Research Council. The authors are grateful to NEDO, Japan, the Ferrous Nano Metal Project, and Professor Setsuo Takaki, Kushu University, Japan, for supply of the Ni-30Fe model alloy. Hodgson acknowledges the provision of a Federation Fellowship by the Australian Research Council. Beladi also acknowledges Deakin University for providing a research scholarship for this work.

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© THE MINERALS, METALS & MATERIALS SOCIETY and ASM INTERNATIONAL 2007