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Microstructure evolution and mechanical behavior of a new microalloyed high Mn austenitic steel during compressive deformation
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Microstructure evolution and mechanical behavior of a new microalloyed High Mn austenitic steel during compressive deformation M. Eskandari, A. Zarei-Hanzaki, J.A. Szpunar, M. A. Mohtadi-Bonab, A.R. Kamali, M. NazarianSamani www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(14)00951-4 http://dx.doi.org/10.1016/j.msea.2014.07.084 MSA31403
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Materials Science & Engineering A
Received date: 31 May 2014 Revised date: 19 July 2014 Accepted date: 22 July 2014 Cite this article as: M. Eskandari, A. Zarei-Hanzaki, J.A. Szpunar, M.A. MohtadiBonab, A.R. Kamali, M. Nazarian-Samani, Microstructure evolution and mechanical behavior of a new microalloyed High Mn austenitic steel during compressive deformation, Materials Science & Engineering A, http://dx.doi.org/ 10.1016/j.msea.2014.07.084 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Nomenclature SFE
stacking fault energy
TRIP
transformation induced plasticity steel
TWIP
twinning induced plasticity steel
DRX
dynamic recrystallization
TEM
transmission electron microscopy
EBSD
electron backscatter diffraction
XRD
X-ray diffraction
TB
twin boundaries
SPD
severe plastic deformation
FCC
face centered body
Microstructure evolution and mechanical behavior of a new microalloyed high Mn austenitic steel during compressive deformation M. Eskandari a,b , A. Zarei-Hanzaki b, J. A. Szpunar a, M.A. Mohtadi-Bonab a, A. R. Kamali c, M. Nazarian-Samani b a
) Advanced Materials for Clean Energy, Department of Mechanical Engineering, University of Saskatchewan, Canada
b)
Hot Deformation & Thermo-mechanical Processing of High Performance Engineering Materials, School of Metallurgy and Materials Engineering, College of Engineering, University of Tehran, Tehran, Iran
c)
Department of Materials Science and Metallurgy, University of Cambridge, Cambridge, United Kingdom E-mail address: [email protected], [email protected]
+1 306 262 7302
Abstract To study the microstructural and mechanical behavior of a new microalloyed TRansformation-Twinning induced plasticity steel, the compression tests were conducted from 25 to 1000°C. Five different categories of true stress-true strain curve, exhibiting distinctive strain hardening/strain softening and stress fluctuation, have been characterized in the range of deformation temperature. The experimental steel exhibits the strength of about 1.28 GPa after straining to 0.6 along with an outstanding strain hardening rate (~14000 MPa) during compression at
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25°C.
This excellent combination of mechanical properties is attributed to the occurrence of strain-induced
martensite transformation. The martensitic transformation continues occurring up to temperatures as high as 150°C, while mechanical twinning is considered as the chief deformation mechanism in the temperature range of 150 to 600°C. Furthermore, a banded-like structure is observed in the temperature range of 600 to 1000°C as a result of precipitate formation. This banded structure is mainly composed of planar slip as well as deformation twins with misorientation of about 59°.
Keywords: EBSD; deformation twinning; martensite; steels; recrystallization
1. Introduction In the past two decades, new groups of low stacking fault energy (SFE) alloys such as transformation induced plasticity (TRIP) steels, twinning induced plasticity (TWIP) steels and bainitic-martensitic functionally graded steels have been paid more attention due to their good combination of high strength and ductility [1-6]. Modeling of flow stress of such steels during hot compression has been investigated by Berto and Lazzarin [4]. Depending on the composition and deformation temperature, the characterized plasticity in TRIP/TWIP steels have been related to the martensitic transformation, mechanical twinning and dislocation glide [3].
Such
mechanisms are strongly dictated by the alloy SFE [4-6]. Allain et al. [7] believe that the strain-induced martensite activates where the SFE is lower than 18 mJ/m2, while the mechanical twinning occurs where the SFE value ranges from 12 to 35 mJ/m2. Consequently due to the observed SFE overlapping, the co-activating of the mechanical twinning and martensite formation is highly logical in a certain SFE range [8-10]. As is well reported the deformed austenite microstructure in TRIP-TWIP steel is highly affected by the alloy SFE and the deformation conditions thereby exhibiting a wide range of mechanical properties. In addition, the plastic deformation reveals the macroscopic plastic instabilities in a wide range of strain-rate and/or temperature. This is characterized by the various shape of stress fluctuation in the true stress–true strain curves [11]. This follows the temperature-SFE dependence of the plastic flow micro-mechanisms in this material [12]. The twin formation is believed to be accompanied with pronounced stress fluctuation, although it is usually relevant to low-temperature deformation [13].
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As is well understood where the deformation temperature is decreased to room temperature, the twinning and martensitic transformations are successively observed in accordance with the diminishing SFE in common TRIP and TWIP steels. In addition, it is believed that there is a transition in twin formation with increasing temperature [14]. This transition temperature is increased where SFE decreases [13,14]. The recent studies [15,16] have demonstrated that this temperature is around to 350°C in TWIP steels. Nevertheless Meyers et al. [13] have studied the twinning stress as a function of temperature for a number of polycrystalline alloys, and found that the twinning stress is insensitive to temperature. The latter is yet a matter of significant debates [17,18]. It is worth mentioning that the positive or negative temperature dependence of SFE is dictated by various factors [9,18]. The main ones are the changes in elastic constants, changes in lattice friction forces, solute or impurity pinning of dislocations, Suzuki segregation to the stacking fault and changes in short range order, and local or overall chemical composition due to the solubility changes of the alloying elements with temperature. The local composition of the alloy may be changed during hot deformation, for instance due to the strain-induced precipitation or decomposition, thereby resulting in SFE changes.
All things considered, in TWIP steels with positive
temperature dependence of SFE, the increase in the deformation temperature retards the mechanical twinning due to the difficult dissociation of perfect dislocations.
Conversely, in the TWIP steels with negative temperature
dependence of SFE, it is logical to deduct that depending on the composition and deformation temperature, the occurrence of mechanical twinning at high temperature is theoretically possible. Therefore, it is interesting to investigate, the effect of microalloying elements such as Ti, Nb, and V, particularly their carbides on the formation of mechanical twinning during deformation at high temperatures. To date the mechanical twinning has been observed at relatively low temperature in TWIP steels however there is no reporting on the high temperature mechanical twinning. It is to emphasize that there is few reports on the formation of high temperature mechanical twins in Ti–6Al–4V at 800°C, and 316L austenitic stainless steel in the range of 800-900°C during severe plastic deformation (SPD) [20,21]. It is believed that the mechanical twins would be formed at any temperature during SPD processes. Such deformation twins show a nano thickness of about 2040 nm. To the best of the present authors’ knowledge, there has been no study reporting significant mechanical twinning or deformation band activities during deformation at such high temperature (600 to 1000°C) under low to moderate strain rates in TRIP-TWIP steels. Furthermore there are little studies [10,15] on the relationship of plasticity to microstructural evolution in steels exhibiting both TRIP and TWIP effects.
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The structure of this paper is arranged as is follows. Firstly, the mechanical behavior of a new microalloying TRIP-TWIP steel in the temperature range of 25 to 1000°C is briefly described. Then the variation of flow behavior and the corresponding microstructures are explained. Finally the remarkable activity of mechanical twinning at elevated temperatures will be described as a first time observation in TRIP/TWIP high manganese steels.
2. Experimental procedure 2.1. Alloy preparation and thermomechanical treatments The experimental material was received in as-cast condition with chemical composition of Fe-0.10C-21Mn2.50Si-1.60Al-0.02Nb-0.02Ti-0.01V (in wt.%). The schematic diagram illustrating the various processing steps in the present work is provided in Fig. 1. The as-cast material was first homogenized at 1180°C for 8 hours under air atmosphere. Then it was heated to 1200°C, held for 30 minutes, and rolled to total strain of 1.5. The material was immediately quenched in water upon being hot-rolled. Finally, after annealing the two specimens with different grain size of about 40±5 (A-type) and 500±50 (B-type) µm were produced. The true stress–true strain behavior of the steel has been studied through compression testing scheme from 25 to 1000°C under constant strain rate of 0.01 s-1 using a Gotech AI-7000 universal testing machine. The compression tests have been conducted according to ASTM E209 standard [22].
2.2. Microstructure characterization The microstructural analysis was performed via an optical microscopy (MEIJI) and a scanning electron microscopy (FEG-SEM, Carl Zeiss Ultra Plus, Carl Zeiss NTS GmbH, Oberkochen, Germany) operating at 20 kV, equipped with an electron backscatter diffraction (EBSD). To prepare the specimens, the mounted samples were mechanically polished with diamond pastes of 6 and 1 μm after grinding. The specimens were further electropolished with electrolyte A2 from Struers (Struers GmbH, Willich, Germany). The electro-polishing was performed between 3 and 6 K by applying a voltage potential of 32 V for 20 seconds. For optical microscopic observations, the specimens were mechanically polished using diamond paste and then electro-etched at room temperature in 65% nitric acid solution. The transmission electron microscopy (TEM) was conducted on a 200 kV JEOL 2000FX instrument equipped with an electron diffractometer. The TEM specimens were prepared by manual polishing of the samples to a nominal thickness of 100 μm, and then further
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thinned to electron transparency by an ion beam milling technique. Electron transparency was obtained by twin-jet electro-polishing using 20 V DC at a temperature of -30 °C. A solution of 5 vol.% perchloric acid in methanol served as the electrolyte. The Clemex software was used to calculate the grain size. In order to investigate the occurrence of the TRIP effect, a Ferriteoscope MP30 (to estimate the magnetic phase volume fraction) has been utilized. Moreover, the actual martensite content was determined using the following equation [23]: Vol. % martensite = 1.75 × Ferriteoscope reading
(1)
The phase identification analysis by X-ray diffraction (XRD) was carried out using a Philips X'Pert PRO apparatus (Cu Kα1 radiation, λ=1.54056 Å, operated at 40 kV).
3. Results 3.1. Mechanical properties Figure 2a shows the true stress-true strain variation of the A-type material with temperature (in the range of 25 to 1000°C). As is seen, the experimental TRIP-TWIP steel in A-type condition exhibits a high compressive strength of about 1280±10 MPa for corresponding strain of 0.6 (true strain) and the yield strength of 385±5 MPa at 25°C. The variation of flow stress versus temperature after straining to 0.5 (true strain) is plotted in Fig. 2b. As is expected the strength level varies with deformation temperature and the variation profile may be classified into three temperature ranges (i.e., 25 to 200°C, 200 to 700 °C and 700 to 1000°C). The rigorous variation of strength (in the temperature range of 25 to 200°C) is attributed to the change of deformation mechanism from TRIP to merely TWIP effect during straining. This is clearly understood from Fig. 3 where the relationship between volume fraction of martensite and deformation temperature is demonstrated. As is seen, the volume fraction of deformation-induced martensite reaches 54%, 18% and 3% after straining to 0.6 (true strain) at 25°C, 100°C and 150°C, respectively.
In addition there is no martensite in all deformed specimens at 200°C to 1000°C.
Consequently, the Md temperature of the present steel appears to be in the range of 150 to 200°C. The true stress-true strain behavior of material at three temperature ranges observed in Fig. 2b will be discussed in details in the next paragraphs. The variation of strain hardening rate (dσ/dε) with true strain is superimposed on the true stress-true strain curve of the A-type specimens at 25°C in Fig. 4a. The high values of
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strain hardening rate is attributed to the strain-induced martensite. Likewise, the variation of strain hardening rate is classified into different stages. The curve drops rapidly in the initial period (elastic stage). However, in the plastic deformation stage, it first slightly decreases (stage A), then becomes constant in the middle stage (stage B) and finally starts monotonically reducing (stage C). To better clarification of these three stages, the variation of volume fraction of martensite with true strain is shown in Fig. 4b. The stage A (the true strain of 0.05 to 0.10) is assigned to the planar glide type of dislocation slip as the dominant deformation mechanism [24]. This idea is supported with the fact that the martensite has not been detected in this interval of strain (see Fig. 4b). In stage B (the true strain of 0.10 to 0.31) the rate of strain hardening is leveled off owning to the strain-induced martensite. The amount of martensite reaches 39 percent in this stage. The martensite platelets may divide the original austenite grains into submicron-scale regions thereby playing an important role in restricting the dislocation movement. This would result in a higher rate of strain hardening. Finally, in stage C (the true strain of 0.31 to 0.60), the rate of strain hardening starts reducing due to the lower rate of martensite formation. The martensite/twinning activity reduction in this stage is clearly understood from Fig. 4b. A different type of stress-strain behavior is observed during straining at 600°C (Fig. 5) in the A-type specimens. A continuous serration type behavior in true stress-true strain curve is realized. This pronounced serration is not seen at lower temperatures. The variation of dσ/dε with true strain is superimposed on the true stress-true strain curve of the experimental alloy at 800°C in Fig. 6. The different strain hardening behavior is observed in the B and A-types specimens (Fig. 6a and 6b, respectively). As is seen in Fig. 6a, in stage A, the strain hardening rate becomes plateau, same as stage B in Fig. 4a, eventually the strain hardening rate starts reducing. Indeed, the stress-strain curve of this specimen shows a moderate steep at early stage of deformation followed by a flat-plateau behavior. This is the third type of stress-strain curve. The fourth type of stress-strain curve is shown in Fig. 6b, which shows a continuous reduction in strain hardening rate. The fifth type of stress-strain curve at 1000 °C has been depicted in Fig. 7. This curve displays a strain hardening behavior after a period of strain softening. It sounds that during primary stage of deformation at 1000°C, the restoration processes are dominant (Fig 7a, B-type specimen and Fig 7b, A-type specimen). However, this is followed by a pronounced strain hardening behavior.
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3.2. Microstructural evolution The microstructures of the A and B-types experimental material consist of fully austenitic structure with a grain size of 40±5 and 500±50 μm, respectively, (Fig. 8). The typical micrographs of A-type specimens after compression to different strain at 25°C are shown in Fig. 9. The morphologies of martensite after straining to 0.15 are demonstrated in Fig. 9a. As is well established the martensite/austenite interface is not strong enough to prevent the growth of martensite [28], hence, the martensite continues to grow up to form the martensite intersections. Such intersections are seen in Fig. 9c after straining to 0.6 (true strain). Figure 9d demonstrates the overall shape change of a deformed specimen. As is seen, some ripples are formed which probably pertains to strain-induced martensite. Figure 10 shows EBSD image of compressed specimen at 25°C after straining to 0.30. As is seen both άmartensite (blue color, BCC) and
-martensite (yellow color, HCP) have been formed. Figure 3 properly indicates
that the TRIP effect as one of the main deformation mechanisms has been finished during compression from 150°C to 200°C.
After that deformation twinning acts as main deformation mechanism. Figure 11 displays the
microstructure of martensite and mechanical twinning during compression at 100, 300, 400 and 600 °C after straining to 0.60 (true strain). Interestingly, as is shown in Fig. 11d, a banded-like structure is formed during straining at 600°C. It appears that the secondary twin system has been activated at 600°C thereby leading to stress fluctuation in the true stresstrue strain curve (see Fig. 5). Theoretically such stress fluctuation is attributed to either dynamic strain aging or twin formation. The dynamic strain aging is more prominent in alloys with a carbon concentration greater than 0.5 mass%. [29,30]. Figure 11e indicates EBSD image of compressed specimen at 300°C. The deformation twinning can be seen in later EBSD image. Figure 12 displays the high temperature banded-like structure in A-type specimens after hot compression at 700, 800 and 900°C to true strain of 0.60. Figure 12d demonstrates the overall shape change of a deformed specimen after hot compression at 800°C. The ripples may well pertain to twinning activity [31]. It should be mentioned that such high temperature banded-like structure have been also observed (Fig.13) in B-type specimens during hot compression at 700, 800 and 900°C after straining to 0.60 (true strain). As is seen, there some few lines not bridging the austenite grains. On the other hand, it is well believed that in the ample of time the deformation bands bridge the entire diameter of the grains [32]. The majority of observed lines do not connect to the primary austenite grain boundary thereby forming a banded-like structure inside the grains. The connection/disconnection of
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twin boundary to austenite grain boundaries depends on the state of deformation (i.e. deformation temperature, deformation mode) [33]. Nevertheless, these lines are configured in a way that a severe grain refinement in austenite grains is occurred. As is seen in Figs.13d and 13e, the formation of banded-like structure has led to overall shape change like a ripple structure which has been already seen during compression at 25°C, too (see Fig. 9d). The SEM-EBSD micrographs of A and B-types specimens after hot compression at 800°C (0.60 true strain) are given in Figs. 14 and 15, respectively. As is seen in Fig.14, a few fine dynamic recrystallization grains (DRX) have been formed at the vicinity of primary austenite grain boundaries. The majority of such grains contain annealing and/or deformation twinning (indicated by arrows), which is called DRX twin. The other feature, which is depicted in Fig. 14 and 15, is a banded-like structure inside the austenite grains. A delicate look at the microstructure reveals that the most of grains compose of intersected lines. Such intersections have also been observed in Figs. 12 and 13. As is seen in Fig. 15, deformation twinning is formed. TEM observation was employed in order to study the planar structure inside the grains observed in Figs. 14 and 15. Figure 16 shows a TEM micrograph of B-type specimen after straining to 0.10 at 900 °C. As observed, the deformation lines occur in arrays of planar dislocation at high temperature. The planar arrays of dislocation are expected at room temperature deformation. However, existence of short-range order precipitation may trigger slip planarity at high temperature. The inset in Fig. 16 is the selected area diffraction pattern of the corresponding region. As is seen, the pattern consists of diffraction rings and spots. The rings, indicated by A in the micrograph, were identified to be the reflections of gamma-austenite. The spots, indicated by B, can be assigned to the diffractions from (200), (210) and (310) planes of single-crystalline cubic AlFe3C phase (k-carbide as a short-range order precipitation). In addition, some other diffraction spots, indicated by C, were detectable which may pertain to complex carbides.
3.3. XRD studies Figure 17 displays the XRD pattern of the specimens before and after hot compression. The XRD pattern shown in Fig. 17a confirms the A-type specimen prior to hot compression contains single phase with nearly complete austenitic structure. Figure 17b shows carbides formation after hot compression at 800°C. Characterizing the VC (200) reflection is difficult due to the probable overlapping with the high-intensity γ-austenite (111) peak located at 2θ=43.75°. There is a same problem to differentiate the TiC (111) reflection with NbC (200) peak located
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at 2θ=41.55° and TiC (220) reflection with NbC peak located at 2θ=74.81°. It is to emphasize that no martensite was detected by Ferriteoscope.
4. Discussion 4.1. Mechanical behavior The strain-induced martensite and twinning are two major competing mechanisms assisting the plastic deformation of TRIP/TWIP steels. The dominating deformation mechanism of an alloy is determined by both the deformation conditions (such as strain, strain rate, temperature) and the nature of materials such as SFE. The latter plays the most important role on the mechanical behavior of the present steel. Taking this fact into consideration, the plastic deformation in experimental steel is more controlled by TRIP effect in the range of 25-150°C. Martensitic phase transformation in the present steel is suppressed by increasing deformation temperature to higher than 150°C. Therefore, the experimental steel shows higher temperature TRIP effect in comparison to the previous TRIP steels reported in the literature [1,14-15]. The strain hardening behavior of experimental steel (Fig. 4) is significantly different from that of medium to high SFE FCC metals, where the plastic deformation is carried by dislocation glide, without the interference of twinning/martensite. The strain hardening behavior in such metals commonly exhibits a continuous decrease in dσ/dε with strain (stage C), followed by a plateau-like stage hardening at large strains, eventually leading to asymptotic stress saturation [25]. Idrissi et al. [26] have identified specific dislocation reactions (i.e., so-called deviation mechanism) taking place in a low SFE FCC metal at an early stage of deformation, which bring about glide and storage of partial dislocations on a plane other than the primary glide plane. Such reactions provide an extra source of dislocation multiplication, which can boost the strain hardening rate. Drobnjak and Parr [27] also have suggested that the interaction of dislocations with stacking faults might be responsible for the increased strain hardening rate in low SFE FCC metals. Surprisingly, the observed banded-like structure shows an appropriate strain hardening behavior at high temperature in B-type specimens (see Fig. 6a). Comparing Fig. 12 and 13 shows a higher deformation band or twin activity in B-type specimen (with coarser initial grain size, 500±50 μm) than that of A-type one.
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The strengthening effect of twin boundaries (TB) have been recently addressed by experimental measurements as well as molecular dynamic simulations in FCC metals [21–23,42,43]. It is believed that TBs are effective in blocking dislocation motion, like conventional grain boundaries, and a Hall–Petch-type relationship exists for TB strengthening. Thus, it seems the banded-like structure acts as barriers to slip for dislocation at high temperature. In contrast, corresponding to Fig. 6b, the dσ/dε shows continuous reduction. According to Fig. 14, the occurrence of DRX in A-type specimen (with initial finer grain size) has resulted in strain softening behavior. At this condition, the DRX phenomenon has dominant effect on strain hardening rate and compensates the banded structure effect. Therefore, it appears the banded structure and DRX process are competitive mechanisms at elevated temperature mechanical behavior. The coexistent features related to these mechanisms have been shown in Fig. 13c. As was shown in Fig. 7, a strain hardening appears after a strain softening. In the other word, the strain hardening occurs after specific strain, which will be discussed in section 4.3. Figure 7a, indicates a slight strain softening behavior due to the insufficient DRX (B-type specimen with coarser initial grain size and higher DRX critical strain, see Fig. 13c) followed by a pronounced strain hardening. In this condition the effect of banded structure overcome DRX effect. In contrast, Fig. 7b (A-type specimen with finer initial grain size) indicates a pronounced strain softening due to the sufficient DRX (see Fig. 14) followed by a slight strain hardening effect. Thus, DRX is predominant mechanism in this condition.
4.2. Microstructural evolution 4.2.1. Nature of the banded-like structure There are a couple of rationales on the nature of the banded structure formation during high temperature deformation in the present TRIP/TWIP steel. The first idea covers that these lines are deformation/planar slip, while the second one is based on deformation twinning. The studies on deformation bands are not numerous [32,34]. In common, the deformation band can be seen in pancaked austenite structure deformed in non-recrystallization region, many of which bridge the grains by roughly parallel lines [32]. In addition, the previous works on hot deformation of TWIP steels did not report the formation of deformation bands [3,16,35-37]. The presence of the short-range order as well as carbide precipitates significantly affect plastic deformation. Frommeyer and Brux [44, 45] observed the occurrence of planar slip in Fe-28Mn-10Al-0.5C steel with high SFE
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during room temperature straining. They suggested that the formation of k-carbide precipitates at high temperature is in charge of formation of planar bands during ambient deformation. The plastic deformation within the austenite matrix of the present steel at high temperature occurs by the formation and propagation of planar slip after precipitates formation. Considering the presence of ordered k-carbide precipitates (Fig. 16 and 17), it may be suggested that the formation of k-carbide precipitates results in slip planarity at high temperature. In the other hand, formation of other carbide precipitations (i.e. VC, TiC) led to carbon depletion in austenite matrix [46,47]. This is led to decreasing SFE locally result in formation of deformation twin (shown in Fig. 15). Therefore, it appears that based on the nature of precipitations the banded-like structure is formed of planar slip along with deformation twins.
4.3. Mechanisms of banded-like structure formation Although it is the first time that the occurrence of mechanical twinning at such high temperature in TRIPTWIP steel is reported, G. Yapici et al. [20,21] have already demonstrated the occurrence of mechanical twinning at 800°C in Ti alloy and 316L austenitic stainless steel during SPD process. The observed twin thickness was smaller than 30 nm. They believe that mechanical twinning can be happening during SPD processing with no temperature limitation. This was attributed to the severe state of plastic deformation (i.e., high stress, strain and strain rate) imposed during SPD process, along with the favorable texture evolved during deformation. The mechanical twinning was also observed in Ti single crystals [48] above 400°C under compression loading. This was attributed to the fact that, while the critical shear stress for nucleation of a definite mode of twin formation would increase with increasing temperature, for the other mode of twin formation might decrease by increasing temperature. It is believed that [49-50] as the required stress level increases, the separation distance between twinning partial dislocations creating a so-called effective SFE enhances. In addition it is believed [50] that at higher strain rates and/or some other loading conditions where high stress levels can be achieved, the SFE might be altered due to the effect of applied stress on the partial dislocation separation. In FCC materials, the applied stress plays a significant role in partial dislocation separation resulting in an effective SFE. Marcinkowski et al. [51] and Copley et al. [52] have indicated that in FCC materials with low-SFE, the applied stress changes the equilibrium
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separation distance of the Shockley partials. Since the partial dislocation separation and the SFE are inversely related, the applied stress creates an effective SFE as experimentally shown in low-SFE austenitic stainless steels [53]. Depending on texture and the stress level, the effective SFE can be quite low and lead to the mechanical twin nucleation [54-55]. It is to mention that the temperature dependence of SFE depends on various factors [19]. The main factors are solute or impurity pinning of dislocations, Suzuki segregation to the stacking fault and changes in short range order and local or global chemical composition due to changes in the solubility of the alloying element with temperature. During hot compression, changes in the global or local composition of the alloy (strain-induced precipitation or resolution) may occur and contribute to change in SFE (Fig. 16 and 17). Therefore, another proposed mechanism is carbide precipitation at high temperature that depletes carbon from austenite matrix, and leads to the decrease in SFE. Decreasing the SFE facilitate the mechanical twinning even at such high temperatures. Due to the presence of Ti, V and Nb along with carbon content (0.11 wt.℅) in the present steel, this mechanism is more feasible. It should be noted that the latter mechanism was seen in stable 316 and 304 austenitic stainless steels which boost the Ms temperature up to higher temperature and the formation of thermal martensite during air cooling [23]. Ishida et al. [56] confirmed the precipitation of M3C ((Fe,Mn)3C) and ((Fe,Mn)3AlC) carbides in the Fe-MnAl-C TWIP steel at high temperature region of 900–1200°C.
In addition, Kim [57] observed many partial
dislocations in contact with hcp AlN precipitates in Fe-24.8Mn-0.28C-1.3Al-0.044N TWIP steel and considered that the precipitations would provide nucleation sites for mechanical twins under external stress. It is interesting to note that there are no DRX grains at the lines intersections in the banded structure. The majority of fine DRX grains have located at the vicinity of primary austenite grain boundaries (Fig. 14). Tamura et al. [32] stated that to accommodate the strain compatibility across the twin boundary, the region at the vicinity of twin boundary is highly strained and thereby becomes a preferred site for nucleation. The relative importance of nucleation sites is considered to be dependent on the amount of deformation. Where the amount of deformation is large, the TBs along with deformation bands become more important nucleation sites.
According to the
aforementioned mechanisms for high temperature mechanical twinning, the formation of strain-induced precipitation and/or Suzuki effect mechanism needs time to occur during hot compression. Therefore, an incubation time is needed to form high temperature mechanical twins. This rationalize that the TBs experience only few fraction of total straining of 0.60, not whole strain. Thus, those TBs, which form the banded structure does not
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suffer large amount of deformation to be nominee as nucleation sites. To support this, the strain hardening due to banded structure occurs after a specific strain (see Fig. 7). It can be concluded that the results of the present study have provided useful insights into the evolution of twinning and the possible mechanisms contributing to strain hardening of TRIP/TWIP steels. This may assist the design of new steel grades and the modeling of their deformation behavior.
5. Conclusion
The present work has been conducted on a new microalloying TRIP-TWIP steel with the principal conclusions as are follows:
•
The experimental TRIP-TWIP steel shows a compressive strength of about 1280±10 MPa after straining to 0.60 with yield strength of 385±5 MPa as well as an outstanding strain hardening rate during compression at the 25 °C.
•
The TRIP effect is diagnosed from 25°C to 150°C. The mechanism of deformation is varied from TRIP effect to solely TWIP one in the temperature range of 150°C to 600°C.
•
A banded structure is observed in the range of 600 to 1000°C. This banded structure consists of planar slip as well as deformation twin with a misorientation angle of about 59°.
•
The observed ripple structure in outer surface of strained specimens refers to the strain-induced martensite at room temperature and twinning activity at 800°C and 900°C.
•
The DRX grain has been formed at the vicinity of primary austenite grain boundaries at 800°C. The banded structure is free of DRX grains.
•
The banded structure and restoration processes are competitive mechanisms at 800°C and 900°C thereby changing the rate of strain hardening.
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References [1] A. Zarei-Hanzaki, P.D. Hodgson and S. Yue, Metall. Mater. Trans. A 28 (1997) 2405-14. [2] A.S. Hamada and L.P. Karjalainen, Mater. Sci. Eng. A 528 (2011) 1819-27. [3] V. Torabinejad, A. Zarei-Hanzaki, S. Moemeni and A. Imandoust, Mater. Des. 32 (2011) 5015-21. [4] F. Berto, P. Lazzarin, Mater. Sci. Eng. R 75 (2014) 1-48. [5] M. Abolghasemzadeh, H. S. Salavati Pour, F. Berto, Y. Alizadeh, Mater. Sci. Eng. A 534 (2012) 329-338. [6] F. Berto, P. Lazzarin, P. Gallo, J. Strain Anal. Eng. 49 (2014) 244-256. [7] T.S. Byun, Acta Mater. 51 (2003) 3063-71. [8] P.J. Ferreira, P. Mullner, Acta Mater. 46 (1998) 4479-84. [9] I. Karaman, H. Sehitoglu, K. Gall, Y.I. Chumlyakov, H.J. Maier, Acta Mater. 48 (2000) 1345-59. [10] S. Allain, J.-P. Chateau, O. Bouaziz, S. Migot, N. Guelton, Mater. Sci. Eng. A 387 (2004) 158-62. [11] C. Scott, S. Allain, M. Faral, N. Guelton, Rev. Metall. 6 (2006) 293-98. [12] A. Saeed-Akbari, J. Imlau, U. Prahl, W. Bleck, Metall. Mater. Trans. A 40 (2009) 3076-90. [13] M. Eskandari, A. Zarei-Hanzaki, A. Marandi, Mater. Des. 39 (2012) 279-84. [14] T.A. Lebedkina, M.A. Lebyodkin, J.-Ph. Chateau, A. Jacquesm S. Allain, Mater. Sci. Eng. A 519 (2009) 14754. [15] S. Allain, J.P. Chateau, O. Bouaziz, Steel Res. 73 (2002) 299-304. [16] M.A. Meyers, O. Vohringer, V.A. Lubarda, Acta Mater. 49 (2001) 4025-39. [17] Y.T. Zhu, X.Z. Liao, X.L. Wu, Prog. Mater. Sci. 57 (2012) 1-62. [18] A. Kovalev, A. Jahn, A. Wei, P.R. Scheller, Steel Res. Int. 82 (2011) 45-49. [19] M. Sabet, A. Zarei-Hanzaki, S.H. Khoddam, J. Eng. Mater. Technol. 131 (2009) 502-08. [20] G.F. Bolling, R.H. Richman, Acta Metall. 13 (1965) 709-22. [21] J.W. Christian, S. Mahajan, Prog. Mater. Sci. 39 (1995) 1-157. [22] X. Tian, Y. Zhang, Mater. Sci. Eng. A 516 (2009) 73-7. [23] G.G. Yapici, I. Karaman, Z. Luo, Acta Mater. 54 (2006) 3755-71. [24] G.G. Yapici, I. Karaman, J. Mater. Res. 19 (2004) 8-13.
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[25] ASTM E209. Standard practice for compression tests of metallic materials at elevated temperatures with conventional or rapid heating rates and strain rates. Annual book of ASTM standards. 03 (2010) 01-05. [26] M. Eskandari, A. Kermanpur, A. Najafizadeh, Metall. Mater. Trans. A 40 (2009) 2241-49. [27] H. Ding, H. Ding, D. Song, Z. Tang, Mater. Sci. Eng. A 528 (2011) 868-73. [28] U.F. Kocks, H. Mecking, Prog. Mater. Sci. 48 (2003) 171-273. [29] H. Idrissi, K. Renard, L. Ryelandt, D. Schryvers, P.J. Jacques, Acta Mater. 58 (2010) 2464-76. [30] D.J. Drobnjak, J. Gordon Parr, Metall. Trans. 1 (1970) 759-65. [31] F. Lu, P. Yang, L. Meng, F. Cui, H. Ding, Mater. Sci. Technol. 27 (2011) 257-60. [32] L. Chen, H.-S. Kim, S.-K. Kim, B.C. De Cooman, ISIJ Int. 47 (2007) 1804-12. [33] Y.N. Dastur, W.C. Leslie, Metall. Trans. A 12 (1981) 749-59. [34] G. Dini, A. Najafizadeh, R. Ueji, S.M. Monir-Vaghefi, Mater. Lett. 64 (2010) 15-8. [35] I. Tamura, H. Sekine, T. Tanaka, C. Ouchi: Thermomechanical Processing of High-Strength Low-Alloy Steels, Butterworth & Co., 1988. [36] K. Renard, H. Idrissi, D. Schryvers, P.J. Jacques, Scr. Mater. 66 (2012) 966-71. [37] H. Sekine, T. Maruyama, Seitetsu Kenkyu 289 (1976) 11920-25. [38] V. Torabinejad, A. Zarei-Hanzaki, M. Sabet, H.R. Abedi, Mater. Des. 32 (2011) 2345-49. [39] A.S. Hamada, L.P. Karjalainen, M.C. Somani, R.M. Ramadan, Mater. Sci. Forum 550 (2007) 217-22. [40] J. Hajkazemi, A. Zarei-Hanzaki, M. Sabet, S. Khoddam, Mater. Sci. Eng. A 530 (2011) 233-8. [41] C.S. Barrett: Structure of Metals, 2nd ed., McGraw-Hill, New York, 1952. [42] A. Kreisler, R.D. Doherty, Metal. Sc. 12 (1978) 551-57. [43] H. Inagaki, Trans Iron and Steel Inst. Japan 23 (1983) 1059-65. [44] M. Umemoto, I. Tamura, H. Otsuka, Tetsu-to-Hagane 68 (1982) 1381-87. [45] O. Bouaziz, N. Guelton, Mater. Sci. Eng. A 319 (2001) 246-9. [46] H. Beladi, I.B. Timokhina, Y. Estrin, J. Kim, B.C. De Cooman, S.K. Kim, Acta Mater.59 ( 2011) 7787-99. [47] G. Frommeyer, U. Brux, Steel Res. Int. 77 (2006) 627-633. [48] J.D. Yoo, K.-T. Park, Mater. Sci. Eng. A 496 (2008) 417-24. [49] H. Beladi, P. Cizek, P.D. Hodgson, Metall. Mater. Trans. A 40 (2009) 1175-89. [50] X. Zhang, T. Sawaguchi, K. Ogawa, F. Yin, X. Zhao, Philos. Mag. 91 (2011) 4410-26.
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[51] S. Poulat, B. Decamps, L. Priester, Philos. Mag. A 77 (1998) 1381-97. [52] S. Poulat, B. Decamps, L. Priester, Philos. Mag. A 79 (1999) 2655-80. [53] N.E.Paton, W.A. Backofen, Metall. Trans. A 1 (1970) 2841-45. [54] M.H. Yoo, Metall. Trans. A 12 (1981) 409-18. [55] M.J. Marcinkowski, D.S. Miller, Philos. Mag. 6 (1961) 871-93. [56] S.M. Copley, B.H. Kear, Acta Metall. 16 (1968) 227-31. [57] D. Goodchild, W.T. Roberts, D.V. Wilson, Acta Metall. 18 (1970) 1137-45. [58] I. Karaman, H. Sehitoglu, H.J. Maier, Y.I. Chumlyakov, Acta Mater. 49 (2001) 3919-33. [59] K. Ishida, H. Ohtani, N. Satoh, R. Kainuma, T. Nishizawa, ISIJ Int. 30 (1990) 680-6. [60] T.W.Kim, Ph.D. thesis. 1993, KAIST, Korea.
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Figure Captions Fig. 1. schematic diagram of the entire process to produce primary specimens. Fig. 2. (a) The true stress–true strain response of the A-type material at different temperature under strain rate of 0.01 s-1, (b) the variation of corresponding flow stress of 0.5 true strain with temperature. Fig. 3.The effect of temperature on the strain-induced martensite formation in the experimental steel. Fig. 4. (a) The variation of strain hardening rate with true strain at 25 °C is superimposed on its corresponding true stress-true strain curve, (b) the volume fraction of strain-induced martensite against true strain in type-A specimen. Fig. 5. The true stress–true strain response of the A-type specimen at 600°C. Fig. 6. The variation of strain hardening rate with true strain at 800°C is superimposed on its corresponding true stress-true strain curve, (a) B-type specimen with grain size of 500±50μm, (b) A-type specimen with grain size of 40±5 μm. Fig. 7. The true stress–true strain response of the specimen at 1000 °C, (a) B-type specimen under strain rate of 0.01 s-1, (b) A-type specimen under strain rate of 0.01 s-1. Fig. 8. The microstructure of the specimen in, (a) A-type condition, and (b) B-type condition. Fig. 9. The SEM microstructures of the A-type specimen after compression at 25°C: (a) after straining to 0.15, composing of martensite platelets, (b) after straining to 0.30, (c) after straining to 0.60, martensite intersections are seen, (d) outer surface, indicating uneven surface like ripple structure. Fig. 10. EBSD micrograph of compressed sample at 25°C after straining to 0.3: (a) band contrast, (b) map color. blue and yellow displays ά and
martensite, respectively.
Fig. 11. The microstructure of A-type specimens after straining to 0.60 at: (a)100°C, indicates the austenite partitioning by deformation-induced martensite, (b) 300°C, and (c) 400°C merely indicates the mechanical twinning, (d) 600°C, containing a banded-like structure links to activated secondary deformation twin, (e) EBSD image of compressed sample at 300°C showing deformation twinning. Fig. 12. The microstructure of A-type specimens after straining to 0.60 at (a) 700 °C,(b) 800°C, the high temperature banded structure links to mechanical twins/deformation band are seen, (c) 900°C, (d) the specimen outer surface after hot compression at 800°C. Fig. 13. The microstructure of B-type specimens after straining to 0.60 at: (a) 700°C, indicating high temperature mechanical twinning/deformation band, (b) 800°C, (c) 900°C, (d) and (e) outer surface after hot compression at 800°C and 900°C, respectively. Fig. 14. SEM-EBSD micrograph of A-type specimens after hot compression at 800°C to 0.60 true strain under strain rate of 0.01 s-1 : (a) and (c) band contrast image, (b) and (d) color map. Black and red lines represent high angle grain boundaries (θ> 15°) and twin boundaries (55°), respectively. Fig. 15. EBSD micrograph of the B-type specimen after hot compression at 900°C to 0.60 true strain under strain rate of 0.01 s-1 . Fig. 16. TEM micrograph of the B-type specimen after hot compression to 0.60 true strain at 900°C; taken from an area of the banded-like structure. The inset is the selected area diffraction pattern of the corresponding region. The
17
rings, indicated by A, are the reflections of gamma-austenite. The spots, indicated by B, can be attributed to the diffractions from single-crystalline cubic AlFe3C phase. The diffraction spots, indicated by C, could not be indexed. Fig. 17. The XRD patterns: (a) the A-type specimen before hot compression, (b) the A-type specimen after hot compression at 800°C under strain rate of 0.01 s-1.
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Figure(s)
Fig. 1. schematic chematic diagram of the entire process to produce primary specimens. specimens
1200 Compressive Strenght (MPa)
(a)
(b)
1000 800 600 400 200 0 0
200
400
600
800
1000
Temperature (°C)
Fig. 2. (a) The true stress–true true strain response of the A-type material at different temperature under strain rate of 0.01 s-1, (b) the variation of corresponding flow stress of 0.5 true strain with temperature.
Fig. 3.The effect of temperature on the strain-induced martensite formation in the experimental steel.
(a)
(b)
Fig. 4. (a) The variation of strain hardening rate with true strain at 25 °C is superimposed on its corresponding true stress-true strain curve, (b) the volume fraction of strain-induced martensite against true strain in type-A specimen.
Fig. 5. The true stress–true strain response of the A-type specimen at 600°C.
(b)
(a)
Fig. 6. The variation of strain hardening rate with true strain at 800°C is superimposed on its corresponding true stress-true strain curve, (a) B-type specimen with grain size of 500±50µm, (b) A-type specimen with grain size of 40±5 µm.
(a)
(b)
Fig. 7. The true stress–true strain response of the specimen at 1000 °C, (a) B-type specimen under strain rate of 0.01 s-1, (b) A-type specimen under strain rate of 0.01 s-1.
(a)
(b)
200 µm
Fig. 8. The microstructure of the specimen in, (a) A-type condition, and (b) B-type condition.
(a)
(b)
(c)
(d)
1 cm
A specimen after compression at 25°C: (a) after straining to 0.1 0.15, Fig. 9. The SEM microstructures of the A-type composing of martensite platelets, (b)) after straining to 0.30, 0.3 (c) after straining to 0.60, martensite intersections are seen, (d) outer surface, indicating uneven surface like ripple structure. structure
(a)
(b)
Fig. 10. EBSD micrograph of compressed sample at 25°C after straining to 0.3: (a) band contrast, (b) map color. blue and yellow displays ά and Ԑ martensite, respectively.
(a)
(b)
(c)
(d)
(e)
Fig. 11. The microstructure of A-type type specimens after straining to 0.60 at: (a)100°C, indicates the austenite partitioning by deformation-induced induced martensite, martensite (b) 300°C, and (c) 400°C merely indicates the mechanical twinning twinning, (d) 600°C, containing a banded-like structure links to activated activat secondary deformation twin, (e) EBSD image of compressed sample at 300°C showing deformation twinning.
(a)
(b)
(c)
(d)
Fig. 12. The microstructure of A-type specimens after straining to 0.60 at (a) 700 °C,(b) 800°C, the high temperature banded structure links to mechanical twins/deformation band are seen, (c) 900°C, (d) the specimen outer surface after hot compression at 800°C.
(b)
(a)
(c)
(d)
(e)
1 cm
Fig. 13. The microstructure of B-type type specimens after straining to 0.60 at:: (a) 700°C, indicating high temperature mechanical twinning/deformation band,, (b) 800°C, 800 (c) 900°C, (d) and (e) outer surface after hot compression at 800°C and 900°C, respectively.
(a)
(b)
Grain boundary serration
Partial dynamic recrystallization
20 µm
Fig. 14. SEM-EBSD micrograph of A-type specimens after hot compression at 800°C to 0.60 true strain under strain rate of 0.01 s-1 : (a) band contrast, (b) color map. Black and red lines represent high angle grain boundaries (θ> 15°) and twin boundaries (55°), respectively.
Twins
Twin-like structure
300 µm
Fig. 15. EBSD micrograph of the B-type specimen after hot compression at 900°C to 0.60 true strain under strain rate of 0.01 s-1.
Fig. 16. TEM micrograph of the B-type specimen after hot compression to 0.60 true strain at 900°C; taken from an area of the banded-like structure. The inset is the selected area diffraction pattern of the corresponding region. The rings, indicated by A, are the reflections of gamma-austenite. The spots, indicated by B, can be attributed to the diffractions from single-crystalline cubic AlFe3C phase. The diffraction spots, indicated by C, could not be indexed.
(a)
(b)
Fig. 17. The XRD patterns: (a) the A-type specimen before hot compression, (b) the A-type specimen after hot compression at 800°C under strain rate of 0.01 s-1.
Microstructure evolution and mechanical behavior of a new microalloyed High Mn austenitic steel during compressive deformation M. Eskandari, A. Zarei-Hanzaki, J.A. Szpunar, M. A. Mohtadi-Bonab, A.R. Kamali, M. NazarianSamani www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(14)00951-4 http://dx.doi.org/10.1016/j.msea.2014.07.084 MSA31403
To appear in:
Materials Science & Engineering A
Received date: 31 May 2014 Revised date: 19 July 2014 Accepted date: 22 July 2014 Cite this article as: M. Eskandari, A. Zarei-Hanzaki, J.A. Szpunar, M.A. MohtadiBonab, A.R. Kamali, M. Nazarian-Samani, Microstructure evolution and mechanical behavior of a new microalloyed High Mn austenitic steel during compressive deformation, Materials Science & Engineering A, http://dx.doi.org/ 10.1016/j.msea.2014.07.084 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Nomenclature SFE
stacking fault energy
TRIP
transformation induced plasticity steel
TWIP
twinning induced plasticity steel
DRX
dynamic recrystallization
TEM
transmission electron microscopy
EBSD
electron backscatter diffraction
XRD
X-ray diffraction
TB
twin boundaries
SPD
severe plastic deformation
FCC
face centered body
Microstructure evolution and mechanical behavior of a new microalloyed high Mn austenitic steel during compressive deformation M. Eskandari a,b , A. Zarei-Hanzaki b, J. A. Szpunar a, M.A. Mohtadi-Bonab a, A. R. Kamali c, M. Nazarian-Samani b a
) Advanced Materials for Clean Energy, Department of Mechanical Engineering, University of Saskatchewan, Canada
b)
Hot Deformation & Thermo-mechanical Processing of High Performance Engineering Materials, School of Metallurgy and Materials Engineering, College of Engineering, University of Tehran, Tehran, Iran
c)
Department of Materials Science and Metallurgy, University of Cambridge, Cambridge, United Kingdom E-mail address: [email protected], [email protected]
+1 306 262 7302
Abstract To study the microstructural and mechanical behavior of a new microalloyed TRansformation-Twinning induced plasticity steel, the compression tests were conducted from 25 to 1000°C. Five different categories of true stress-true strain curve, exhibiting distinctive strain hardening/strain softening and stress fluctuation, have been characterized in the range of deformation temperature. The experimental steel exhibits the strength of about 1.28 GPa after straining to 0.6 along with an outstanding strain hardening rate (~14000 MPa) during compression at
1
25°C.
This excellent combination of mechanical properties is attributed to the occurrence of strain-induced
martensite transformation. The martensitic transformation continues occurring up to temperatures as high as 150°C, while mechanical twinning is considered as the chief deformation mechanism in the temperature range of 150 to 600°C. Furthermore, a banded-like structure is observed in the temperature range of 600 to 1000°C as a result of precipitate formation. This banded structure is mainly composed of planar slip as well as deformation twins with misorientation of about 59°.
Keywords: EBSD; deformation twinning; martensite; steels; recrystallization
1. Introduction In the past two decades, new groups of low stacking fault energy (SFE) alloys such as transformation induced plasticity (TRIP) steels, twinning induced plasticity (TWIP) steels and bainitic-martensitic functionally graded steels have been paid more attention due to their good combination of high strength and ductility [1-6]. Modeling of flow stress of such steels during hot compression has been investigated by Berto and Lazzarin [4]. Depending on the composition and deformation temperature, the characterized plasticity in TRIP/TWIP steels have been related to the martensitic transformation, mechanical twinning and dislocation glide [3].
Such
mechanisms are strongly dictated by the alloy SFE [4-6]. Allain et al. [7] believe that the strain-induced martensite activates where the SFE is lower than 18 mJ/m2, while the mechanical twinning occurs where the SFE value ranges from 12 to 35 mJ/m2. Consequently due to the observed SFE overlapping, the co-activating of the mechanical twinning and martensite formation is highly logical in a certain SFE range [8-10]. As is well reported the deformed austenite microstructure in TRIP-TWIP steel is highly affected by the alloy SFE and the deformation conditions thereby exhibiting a wide range of mechanical properties. In addition, the plastic deformation reveals the macroscopic plastic instabilities in a wide range of strain-rate and/or temperature. This is characterized by the various shape of stress fluctuation in the true stress–true strain curves [11]. This follows the temperature-SFE dependence of the plastic flow micro-mechanisms in this material [12]. The twin formation is believed to be accompanied with pronounced stress fluctuation, although it is usually relevant to low-temperature deformation [13].
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As is well understood where the deformation temperature is decreased to room temperature, the twinning and martensitic transformations are successively observed in accordance with the diminishing SFE in common TRIP and TWIP steels. In addition, it is believed that there is a transition in twin formation with increasing temperature [14]. This transition temperature is increased where SFE decreases [13,14]. The recent studies [15,16] have demonstrated that this temperature is around to 350°C in TWIP steels. Nevertheless Meyers et al. [13] have studied the twinning stress as a function of temperature for a number of polycrystalline alloys, and found that the twinning stress is insensitive to temperature. The latter is yet a matter of significant debates [17,18]. It is worth mentioning that the positive or negative temperature dependence of SFE is dictated by various factors [9,18]. The main ones are the changes in elastic constants, changes in lattice friction forces, solute or impurity pinning of dislocations, Suzuki segregation to the stacking fault and changes in short range order, and local or overall chemical composition due to the solubility changes of the alloying elements with temperature. The local composition of the alloy may be changed during hot deformation, for instance due to the strain-induced precipitation or decomposition, thereby resulting in SFE changes.
All things considered, in TWIP steels with positive
temperature dependence of SFE, the increase in the deformation temperature retards the mechanical twinning due to the difficult dissociation of perfect dislocations.
Conversely, in the TWIP steels with negative temperature
dependence of SFE, it is logical to deduct that depending on the composition and deformation temperature, the occurrence of mechanical twinning at high temperature is theoretically possible. Therefore, it is interesting to investigate, the effect of microalloying elements such as Ti, Nb, and V, particularly their carbides on the formation of mechanical twinning during deformation at high temperatures. To date the mechanical twinning has been observed at relatively low temperature in TWIP steels however there is no reporting on the high temperature mechanical twinning. It is to emphasize that there is few reports on the formation of high temperature mechanical twins in Ti–6Al–4V at 800°C, and 316L austenitic stainless steel in the range of 800-900°C during severe plastic deformation (SPD) [20,21]. It is believed that the mechanical twins would be formed at any temperature during SPD processes. Such deformation twins show a nano thickness of about 2040 nm. To the best of the present authors’ knowledge, there has been no study reporting significant mechanical twinning or deformation band activities during deformation at such high temperature (600 to 1000°C) under low to moderate strain rates in TRIP-TWIP steels. Furthermore there are little studies [10,15] on the relationship of plasticity to microstructural evolution in steels exhibiting both TRIP and TWIP effects.
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The structure of this paper is arranged as is follows. Firstly, the mechanical behavior of a new microalloying TRIP-TWIP steel in the temperature range of 25 to 1000°C is briefly described. Then the variation of flow behavior and the corresponding microstructures are explained. Finally the remarkable activity of mechanical twinning at elevated temperatures will be described as a first time observation in TRIP/TWIP high manganese steels.
2. Experimental procedure 2.1. Alloy preparation and thermomechanical treatments The experimental material was received in as-cast condition with chemical composition of Fe-0.10C-21Mn2.50Si-1.60Al-0.02Nb-0.02Ti-0.01V (in wt.%). The schematic diagram illustrating the various processing steps in the present work is provided in Fig. 1. The as-cast material was first homogenized at 1180°C for 8 hours under air atmosphere. Then it was heated to 1200°C, held for 30 minutes, and rolled to total strain of 1.5. The material was immediately quenched in water upon being hot-rolled. Finally, after annealing the two specimens with different grain size of about 40±5 (A-type) and 500±50 (B-type) µm were produced. The true stress–true strain behavior of the steel has been studied through compression testing scheme from 25 to 1000°C under constant strain rate of 0.01 s-1 using a Gotech AI-7000 universal testing machine. The compression tests have been conducted according to ASTM E209 standard [22].
2.2. Microstructure characterization The microstructural analysis was performed via an optical microscopy (MEIJI) and a scanning electron microscopy (FEG-SEM, Carl Zeiss Ultra Plus, Carl Zeiss NTS GmbH, Oberkochen, Germany) operating at 20 kV, equipped with an electron backscatter diffraction (EBSD). To prepare the specimens, the mounted samples were mechanically polished with diamond pastes of 6 and 1 μm after grinding. The specimens were further electropolished with electrolyte A2 from Struers (Struers GmbH, Willich, Germany). The electro-polishing was performed between 3 and 6 K by applying a voltage potential of 32 V for 20 seconds. For optical microscopic observations, the specimens were mechanically polished using diamond paste and then electro-etched at room temperature in 65% nitric acid solution. The transmission electron microscopy (TEM) was conducted on a 200 kV JEOL 2000FX instrument equipped with an electron diffractometer. The TEM specimens were prepared by manual polishing of the samples to a nominal thickness of 100 μm, and then further
4
thinned to electron transparency by an ion beam milling technique. Electron transparency was obtained by twin-jet electro-polishing using 20 V DC at a temperature of -30 °C. A solution of 5 vol.% perchloric acid in methanol served as the electrolyte. The Clemex software was used to calculate the grain size. In order to investigate the occurrence of the TRIP effect, a Ferriteoscope MP30 (to estimate the magnetic phase volume fraction) has been utilized. Moreover, the actual martensite content was determined using the following equation [23]: Vol. % martensite = 1.75 × Ferriteoscope reading
(1)
The phase identification analysis by X-ray diffraction (XRD) was carried out using a Philips X'Pert PRO apparatus (Cu Kα1 radiation, λ=1.54056 Å, operated at 40 kV).
3. Results 3.1. Mechanical properties Figure 2a shows the true stress-true strain variation of the A-type material with temperature (in the range of 25 to 1000°C). As is seen, the experimental TRIP-TWIP steel in A-type condition exhibits a high compressive strength of about 1280±10 MPa for corresponding strain of 0.6 (true strain) and the yield strength of 385±5 MPa at 25°C. The variation of flow stress versus temperature after straining to 0.5 (true strain) is plotted in Fig. 2b. As is expected the strength level varies with deformation temperature and the variation profile may be classified into three temperature ranges (i.e., 25 to 200°C, 200 to 700 °C and 700 to 1000°C). The rigorous variation of strength (in the temperature range of 25 to 200°C) is attributed to the change of deformation mechanism from TRIP to merely TWIP effect during straining. This is clearly understood from Fig. 3 where the relationship between volume fraction of martensite and deformation temperature is demonstrated. As is seen, the volume fraction of deformation-induced martensite reaches 54%, 18% and 3% after straining to 0.6 (true strain) at 25°C, 100°C and 150°C, respectively.
In addition there is no martensite in all deformed specimens at 200°C to 1000°C.
Consequently, the Md temperature of the present steel appears to be in the range of 150 to 200°C. The true stress-true strain behavior of material at three temperature ranges observed in Fig. 2b will be discussed in details in the next paragraphs. The variation of strain hardening rate (dσ/dε) with true strain is superimposed on the true stress-true strain curve of the A-type specimens at 25°C in Fig. 4a. The high values of
5
strain hardening rate is attributed to the strain-induced martensite. Likewise, the variation of strain hardening rate is classified into different stages. The curve drops rapidly in the initial period (elastic stage). However, in the plastic deformation stage, it first slightly decreases (stage A), then becomes constant in the middle stage (stage B) and finally starts monotonically reducing (stage C). To better clarification of these three stages, the variation of volume fraction of martensite with true strain is shown in Fig. 4b. The stage A (the true strain of 0.05 to 0.10) is assigned to the planar glide type of dislocation slip as the dominant deformation mechanism [24]. This idea is supported with the fact that the martensite has not been detected in this interval of strain (see Fig. 4b). In stage B (the true strain of 0.10 to 0.31) the rate of strain hardening is leveled off owning to the strain-induced martensite. The amount of martensite reaches 39 percent in this stage. The martensite platelets may divide the original austenite grains into submicron-scale regions thereby playing an important role in restricting the dislocation movement. This would result in a higher rate of strain hardening. Finally, in stage C (the true strain of 0.31 to 0.60), the rate of strain hardening starts reducing due to the lower rate of martensite formation. The martensite/twinning activity reduction in this stage is clearly understood from Fig. 4b. A different type of stress-strain behavior is observed during straining at 600°C (Fig. 5) in the A-type specimens. A continuous serration type behavior in true stress-true strain curve is realized. This pronounced serration is not seen at lower temperatures. The variation of dσ/dε with true strain is superimposed on the true stress-true strain curve of the experimental alloy at 800°C in Fig. 6. The different strain hardening behavior is observed in the B and A-types specimens (Fig. 6a and 6b, respectively). As is seen in Fig. 6a, in stage A, the strain hardening rate becomes plateau, same as stage B in Fig. 4a, eventually the strain hardening rate starts reducing. Indeed, the stress-strain curve of this specimen shows a moderate steep at early stage of deformation followed by a flat-plateau behavior. This is the third type of stress-strain curve. The fourth type of stress-strain curve is shown in Fig. 6b, which shows a continuous reduction in strain hardening rate. The fifth type of stress-strain curve at 1000 °C has been depicted in Fig. 7. This curve displays a strain hardening behavior after a period of strain softening. It sounds that during primary stage of deformation at 1000°C, the restoration processes are dominant (Fig 7a, B-type specimen and Fig 7b, A-type specimen). However, this is followed by a pronounced strain hardening behavior.
6
3.2. Microstructural evolution The microstructures of the A and B-types experimental material consist of fully austenitic structure with a grain size of 40±5 and 500±50 μm, respectively, (Fig. 8). The typical micrographs of A-type specimens after compression to different strain at 25°C are shown in Fig. 9. The morphologies of martensite after straining to 0.15 are demonstrated in Fig. 9a. As is well established the martensite/austenite interface is not strong enough to prevent the growth of martensite [28], hence, the martensite continues to grow up to form the martensite intersections. Such intersections are seen in Fig. 9c after straining to 0.6 (true strain). Figure 9d demonstrates the overall shape change of a deformed specimen. As is seen, some ripples are formed which probably pertains to strain-induced martensite. Figure 10 shows EBSD image of compressed specimen at 25°C after straining to 0.30. As is seen both άmartensite (blue color, BCC) and
-martensite (yellow color, HCP) have been formed. Figure 3 properly indicates
that the TRIP effect as one of the main deformation mechanisms has been finished during compression from 150°C to 200°C.
After that deformation twinning acts as main deformation mechanism. Figure 11 displays the
microstructure of martensite and mechanical twinning during compression at 100, 300, 400 and 600 °C after straining to 0.60 (true strain). Interestingly, as is shown in Fig. 11d, a banded-like structure is formed during straining at 600°C. It appears that the secondary twin system has been activated at 600°C thereby leading to stress fluctuation in the true stresstrue strain curve (see Fig. 5). Theoretically such stress fluctuation is attributed to either dynamic strain aging or twin formation. The dynamic strain aging is more prominent in alloys with a carbon concentration greater than 0.5 mass%. [29,30]. Figure 11e indicates EBSD image of compressed specimen at 300°C. The deformation twinning can be seen in later EBSD image. Figure 12 displays the high temperature banded-like structure in A-type specimens after hot compression at 700, 800 and 900°C to true strain of 0.60. Figure 12d demonstrates the overall shape change of a deformed specimen after hot compression at 800°C. The ripples may well pertain to twinning activity [31]. It should be mentioned that such high temperature banded-like structure have been also observed (Fig.13) in B-type specimens during hot compression at 700, 800 and 900°C after straining to 0.60 (true strain). As is seen, there some few lines not bridging the austenite grains. On the other hand, it is well believed that in the ample of time the deformation bands bridge the entire diameter of the grains [32]. The majority of observed lines do not connect to the primary austenite grain boundary thereby forming a banded-like structure inside the grains. The connection/disconnection of
7
twin boundary to austenite grain boundaries depends on the state of deformation (i.e. deformation temperature, deformation mode) [33]. Nevertheless, these lines are configured in a way that a severe grain refinement in austenite grains is occurred. As is seen in Figs.13d and 13e, the formation of banded-like structure has led to overall shape change like a ripple structure which has been already seen during compression at 25°C, too (see Fig. 9d). The SEM-EBSD micrographs of A and B-types specimens after hot compression at 800°C (0.60 true strain) are given in Figs. 14 and 15, respectively. As is seen in Fig.14, a few fine dynamic recrystallization grains (DRX) have been formed at the vicinity of primary austenite grain boundaries. The majority of such grains contain annealing and/or deformation twinning (indicated by arrows), which is called DRX twin. The other feature, which is depicted in Fig. 14 and 15, is a banded-like structure inside the austenite grains. A delicate look at the microstructure reveals that the most of grains compose of intersected lines. Such intersections have also been observed in Figs. 12 and 13. As is seen in Fig. 15, deformation twinning is formed. TEM observation was employed in order to study the planar structure inside the grains observed in Figs. 14 and 15. Figure 16 shows a TEM micrograph of B-type specimen after straining to 0.10 at 900 °C. As observed, the deformation lines occur in arrays of planar dislocation at high temperature. The planar arrays of dislocation are expected at room temperature deformation. However, existence of short-range order precipitation may trigger slip planarity at high temperature. The inset in Fig. 16 is the selected area diffraction pattern of the corresponding region. As is seen, the pattern consists of diffraction rings and spots. The rings, indicated by A in the micrograph, were identified to be the reflections of gamma-austenite. The spots, indicated by B, can be assigned to the diffractions from (200), (210) and (310) planes of single-crystalline cubic AlFe3C phase (k-carbide as a short-range order precipitation). In addition, some other diffraction spots, indicated by C, were detectable which may pertain to complex carbides.
3.3. XRD studies Figure 17 displays the XRD pattern of the specimens before and after hot compression. The XRD pattern shown in Fig. 17a confirms the A-type specimen prior to hot compression contains single phase with nearly complete austenitic structure. Figure 17b shows carbides formation after hot compression at 800°C. Characterizing the VC (200) reflection is difficult due to the probable overlapping with the high-intensity γ-austenite (111) peak located at 2θ=43.75°. There is a same problem to differentiate the TiC (111) reflection with NbC (200) peak located
8
at 2θ=41.55° and TiC (220) reflection with NbC peak located at 2θ=74.81°. It is to emphasize that no martensite was detected by Ferriteoscope.
4. Discussion 4.1. Mechanical behavior The strain-induced martensite and twinning are two major competing mechanisms assisting the plastic deformation of TRIP/TWIP steels. The dominating deformation mechanism of an alloy is determined by both the deformation conditions (such as strain, strain rate, temperature) and the nature of materials such as SFE. The latter plays the most important role on the mechanical behavior of the present steel. Taking this fact into consideration, the plastic deformation in experimental steel is more controlled by TRIP effect in the range of 25-150°C. Martensitic phase transformation in the present steel is suppressed by increasing deformation temperature to higher than 150°C. Therefore, the experimental steel shows higher temperature TRIP effect in comparison to the previous TRIP steels reported in the literature [1,14-15]. The strain hardening behavior of experimental steel (Fig. 4) is significantly different from that of medium to high SFE FCC metals, where the plastic deformation is carried by dislocation glide, without the interference of twinning/martensite. The strain hardening behavior in such metals commonly exhibits a continuous decrease in dσ/dε with strain (stage C), followed by a plateau-like stage hardening at large strains, eventually leading to asymptotic stress saturation [25]. Idrissi et al. [26] have identified specific dislocation reactions (i.e., so-called deviation mechanism) taking place in a low SFE FCC metal at an early stage of deformation, which bring about glide and storage of partial dislocations on a plane other than the primary glide plane. Such reactions provide an extra source of dislocation multiplication, which can boost the strain hardening rate. Drobnjak and Parr [27] also have suggested that the interaction of dislocations with stacking faults might be responsible for the increased strain hardening rate in low SFE FCC metals. Surprisingly, the observed banded-like structure shows an appropriate strain hardening behavior at high temperature in B-type specimens (see Fig. 6a). Comparing Fig. 12 and 13 shows a higher deformation band or twin activity in B-type specimen (with coarser initial grain size, 500±50 μm) than that of A-type one.
9
The strengthening effect of twin boundaries (TB) have been recently addressed by experimental measurements as well as molecular dynamic simulations in FCC metals [21–23,42,43]. It is believed that TBs are effective in blocking dislocation motion, like conventional grain boundaries, and a Hall–Petch-type relationship exists for TB strengthening. Thus, it seems the banded-like structure acts as barriers to slip for dislocation at high temperature. In contrast, corresponding to Fig. 6b, the dσ/dε shows continuous reduction. According to Fig. 14, the occurrence of DRX in A-type specimen (with initial finer grain size) has resulted in strain softening behavior. At this condition, the DRX phenomenon has dominant effect on strain hardening rate and compensates the banded structure effect. Therefore, it appears the banded structure and DRX process are competitive mechanisms at elevated temperature mechanical behavior. The coexistent features related to these mechanisms have been shown in Fig. 13c. As was shown in Fig. 7, a strain hardening appears after a strain softening. In the other word, the strain hardening occurs after specific strain, which will be discussed in section 4.3. Figure 7a, indicates a slight strain softening behavior due to the insufficient DRX (B-type specimen with coarser initial grain size and higher DRX critical strain, see Fig. 13c) followed by a pronounced strain hardening. In this condition the effect of banded structure overcome DRX effect. In contrast, Fig. 7b (A-type specimen with finer initial grain size) indicates a pronounced strain softening due to the sufficient DRX (see Fig. 14) followed by a slight strain hardening effect. Thus, DRX is predominant mechanism in this condition.
4.2. Microstructural evolution 4.2.1. Nature of the banded-like structure There are a couple of rationales on the nature of the banded structure formation during high temperature deformation in the present TRIP/TWIP steel. The first idea covers that these lines are deformation/planar slip, while the second one is based on deformation twinning. The studies on deformation bands are not numerous [32,34]. In common, the deformation band can be seen in pancaked austenite structure deformed in non-recrystallization region, many of which bridge the grains by roughly parallel lines [32]. In addition, the previous works on hot deformation of TWIP steels did not report the formation of deformation bands [3,16,35-37]. The presence of the short-range order as well as carbide precipitates significantly affect plastic deformation. Frommeyer and Brux [44, 45] observed the occurrence of planar slip in Fe-28Mn-10Al-0.5C steel with high SFE
10
during room temperature straining. They suggested that the formation of k-carbide precipitates at high temperature is in charge of formation of planar bands during ambient deformation. The plastic deformation within the austenite matrix of the present steel at high temperature occurs by the formation and propagation of planar slip after precipitates formation. Considering the presence of ordered k-carbide precipitates (Fig. 16 and 17), it may be suggested that the formation of k-carbide precipitates results in slip planarity at high temperature. In the other hand, formation of other carbide precipitations (i.e. VC, TiC) led to carbon depletion in austenite matrix [46,47]. This is led to decreasing SFE locally result in formation of deformation twin (shown in Fig. 15). Therefore, it appears that based on the nature of precipitations the banded-like structure is formed of planar slip along with deformation twins.
4.3. Mechanisms of banded-like structure formation Although it is the first time that the occurrence of mechanical twinning at such high temperature in TRIPTWIP steel is reported, G. Yapici et al. [20,21] have already demonstrated the occurrence of mechanical twinning at 800°C in Ti alloy and 316L austenitic stainless steel during SPD process. The observed twin thickness was smaller than 30 nm. They believe that mechanical twinning can be happening during SPD processing with no temperature limitation. This was attributed to the severe state of plastic deformation (i.e., high stress, strain and strain rate) imposed during SPD process, along with the favorable texture evolved during deformation. The mechanical twinning was also observed in Ti single crystals [48] above 400°C under compression loading. This was attributed to the fact that, while the critical shear stress for nucleation of a definite mode of twin formation would increase with increasing temperature, for the other mode of twin formation might decrease by increasing temperature. It is believed that [49-50] as the required stress level increases, the separation distance between twinning partial dislocations creating a so-called effective SFE enhances. In addition it is believed [50] that at higher strain rates and/or some other loading conditions where high stress levels can be achieved, the SFE might be altered due to the effect of applied stress on the partial dislocation separation. In FCC materials, the applied stress plays a significant role in partial dislocation separation resulting in an effective SFE. Marcinkowski et al. [51] and Copley et al. [52] have indicated that in FCC materials with low-SFE, the applied stress changes the equilibrium
11
separation distance of the Shockley partials. Since the partial dislocation separation and the SFE are inversely related, the applied stress creates an effective SFE as experimentally shown in low-SFE austenitic stainless steels [53]. Depending on texture and the stress level, the effective SFE can be quite low and lead to the mechanical twin nucleation [54-55]. It is to mention that the temperature dependence of SFE depends on various factors [19]. The main factors are solute or impurity pinning of dislocations, Suzuki segregation to the stacking fault and changes in short range order and local or global chemical composition due to changes in the solubility of the alloying element with temperature. During hot compression, changes in the global or local composition of the alloy (strain-induced precipitation or resolution) may occur and contribute to change in SFE (Fig. 16 and 17). Therefore, another proposed mechanism is carbide precipitation at high temperature that depletes carbon from austenite matrix, and leads to the decrease in SFE. Decreasing the SFE facilitate the mechanical twinning even at such high temperatures. Due to the presence of Ti, V and Nb along with carbon content (0.11 wt.℅) in the present steel, this mechanism is more feasible. It should be noted that the latter mechanism was seen in stable 316 and 304 austenitic stainless steels which boost the Ms temperature up to higher temperature and the formation of thermal martensite during air cooling [23]. Ishida et al. [56] confirmed the precipitation of M3C ((Fe,Mn)3C) and ((Fe,Mn)3AlC) carbides in the Fe-MnAl-C TWIP steel at high temperature region of 900–1200°C.
In addition, Kim [57] observed many partial
dislocations in contact with hcp AlN precipitates in Fe-24.8Mn-0.28C-1.3Al-0.044N TWIP steel and considered that the precipitations would provide nucleation sites for mechanical twins under external stress. It is interesting to note that there are no DRX grains at the lines intersections in the banded structure. The majority of fine DRX grains have located at the vicinity of primary austenite grain boundaries (Fig. 14). Tamura et al. [32] stated that to accommodate the strain compatibility across the twin boundary, the region at the vicinity of twin boundary is highly strained and thereby becomes a preferred site for nucleation. The relative importance of nucleation sites is considered to be dependent on the amount of deformation. Where the amount of deformation is large, the TBs along with deformation bands become more important nucleation sites.
According to the
aforementioned mechanisms for high temperature mechanical twinning, the formation of strain-induced precipitation and/or Suzuki effect mechanism needs time to occur during hot compression. Therefore, an incubation time is needed to form high temperature mechanical twins. This rationalize that the TBs experience only few fraction of total straining of 0.60, not whole strain. Thus, those TBs, which form the banded structure does not
12
suffer large amount of deformation to be nominee as nucleation sites. To support this, the strain hardening due to banded structure occurs after a specific strain (see Fig. 7). It can be concluded that the results of the present study have provided useful insights into the evolution of twinning and the possible mechanisms contributing to strain hardening of TRIP/TWIP steels. This may assist the design of new steel grades and the modeling of their deformation behavior.
5. Conclusion
The present work has been conducted on a new microalloying TRIP-TWIP steel with the principal conclusions as are follows:
•
The experimental TRIP-TWIP steel shows a compressive strength of about 1280±10 MPa after straining to 0.60 with yield strength of 385±5 MPa as well as an outstanding strain hardening rate during compression at the 25 °C.
•
The TRIP effect is diagnosed from 25°C to 150°C. The mechanism of deformation is varied from TRIP effect to solely TWIP one in the temperature range of 150°C to 600°C.
•
A banded structure is observed in the range of 600 to 1000°C. This banded structure consists of planar slip as well as deformation twin with a misorientation angle of about 59°.
•
The observed ripple structure in outer surface of strained specimens refers to the strain-induced martensite at room temperature and twinning activity at 800°C and 900°C.
•
The DRX grain has been formed at the vicinity of primary austenite grain boundaries at 800°C. The banded structure is free of DRX grains.
•
The banded structure and restoration processes are competitive mechanisms at 800°C and 900°C thereby changing the rate of strain hardening.
13
References [1] A. Zarei-Hanzaki, P.D. Hodgson and S. Yue, Metall. Mater. Trans. A 28 (1997) 2405-14. [2] A.S. Hamada and L.P. Karjalainen, Mater. Sci. Eng. A 528 (2011) 1819-27. [3] V. Torabinejad, A. Zarei-Hanzaki, S. Moemeni and A. Imandoust, Mater. Des. 32 (2011) 5015-21. [4] F. Berto, P. Lazzarin, Mater. Sci. Eng. R 75 (2014) 1-48. [5] M. Abolghasemzadeh, H. S. Salavati Pour, F. Berto, Y. Alizadeh, Mater. Sci. Eng. A 534 (2012) 329-338. [6] F. Berto, P. Lazzarin, P. Gallo, J. Strain Anal. Eng. 49 (2014) 244-256. [7] T.S. Byun, Acta Mater. 51 (2003) 3063-71. [8] P.J. Ferreira, P. Mullner, Acta Mater. 46 (1998) 4479-84. [9] I. Karaman, H. Sehitoglu, K. Gall, Y.I. Chumlyakov, H.J. Maier, Acta Mater. 48 (2000) 1345-59. [10] S. Allain, J.-P. Chateau, O. Bouaziz, S. Migot, N. Guelton, Mater. Sci. Eng. A 387 (2004) 158-62. [11] C. Scott, S. Allain, M. Faral, N. Guelton, Rev. Metall. 6 (2006) 293-98. [12] A. Saeed-Akbari, J. Imlau, U. Prahl, W. Bleck, Metall. Mater. Trans. A 40 (2009) 3076-90. [13] M. Eskandari, A. Zarei-Hanzaki, A. Marandi, Mater. Des. 39 (2012) 279-84. [14] T.A. Lebedkina, M.A. Lebyodkin, J.-Ph. Chateau, A. Jacquesm S. Allain, Mater. Sci. Eng. A 519 (2009) 14754. [15] S. Allain, J.P. Chateau, O. Bouaziz, Steel Res. 73 (2002) 299-304. [16] M.A. Meyers, O. Vohringer, V.A. Lubarda, Acta Mater. 49 (2001) 4025-39. [17] Y.T. Zhu, X.Z. Liao, X.L. Wu, Prog. Mater. Sci. 57 (2012) 1-62. [18] A. Kovalev, A. Jahn, A. Wei, P.R. Scheller, Steel Res. Int. 82 (2011) 45-49. [19] M. Sabet, A. Zarei-Hanzaki, S.H. Khoddam, J. Eng. Mater. Technol. 131 (2009) 502-08. [20] G.F. Bolling, R.H. Richman, Acta Metall. 13 (1965) 709-22. [21] J.W. Christian, S. Mahajan, Prog. Mater. Sci. 39 (1995) 1-157. [22] X. Tian, Y. Zhang, Mater. Sci. Eng. A 516 (2009) 73-7. [23] G.G. Yapici, I. Karaman, Z. Luo, Acta Mater. 54 (2006) 3755-71. [24] G.G. Yapici, I. Karaman, J. Mater. Res. 19 (2004) 8-13.
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[25] ASTM E209. Standard practice for compression tests of metallic materials at elevated temperatures with conventional or rapid heating rates and strain rates. Annual book of ASTM standards. 03 (2010) 01-05. [26] M. Eskandari, A. Kermanpur, A. Najafizadeh, Metall. Mater. Trans. A 40 (2009) 2241-49. [27] H. Ding, H. Ding, D. Song, Z. Tang, Mater. Sci. Eng. A 528 (2011) 868-73. [28] U.F. Kocks, H. Mecking, Prog. Mater. Sci. 48 (2003) 171-273. [29] H. Idrissi, K. Renard, L. Ryelandt, D. Schryvers, P.J. Jacques, Acta Mater. 58 (2010) 2464-76. [30] D.J. Drobnjak, J. Gordon Parr, Metall. Trans. 1 (1970) 759-65. [31] F. Lu, P. Yang, L. Meng, F. Cui, H. Ding, Mater. Sci. Technol. 27 (2011) 257-60. [32] L. Chen, H.-S. Kim, S.-K. Kim, B.C. De Cooman, ISIJ Int. 47 (2007) 1804-12. [33] Y.N. Dastur, W.C. Leslie, Metall. Trans. A 12 (1981) 749-59. [34] G. Dini, A. Najafizadeh, R. Ueji, S.M. Monir-Vaghefi, Mater. Lett. 64 (2010) 15-8. [35] I. Tamura, H. Sekine, T. Tanaka, C. Ouchi: Thermomechanical Processing of High-Strength Low-Alloy Steels, Butterworth & Co., 1988. [36] K. Renard, H. Idrissi, D. Schryvers, P.J. Jacques, Scr. Mater. 66 (2012) 966-71. [37] H. Sekine, T. Maruyama, Seitetsu Kenkyu 289 (1976) 11920-25. [38] V. Torabinejad, A. Zarei-Hanzaki, M. Sabet, H.R. Abedi, Mater. Des. 32 (2011) 2345-49. [39] A.S. Hamada, L.P. Karjalainen, M.C. Somani, R.M. Ramadan, Mater. Sci. Forum 550 (2007) 217-22. [40] J. Hajkazemi, A. Zarei-Hanzaki, M. Sabet, S. Khoddam, Mater. Sci. Eng. A 530 (2011) 233-8. [41] C.S. Barrett: Structure of Metals, 2nd ed., McGraw-Hill, New York, 1952. [42] A. Kreisler, R.D. Doherty, Metal. Sc. 12 (1978) 551-57. [43] H. Inagaki, Trans Iron and Steel Inst. Japan 23 (1983) 1059-65. [44] M. Umemoto, I. Tamura, H. Otsuka, Tetsu-to-Hagane 68 (1982) 1381-87. [45] O. Bouaziz, N. Guelton, Mater. Sci. Eng. A 319 (2001) 246-9. [46] H. Beladi, I.B. Timokhina, Y. Estrin, J. Kim, B.C. De Cooman, S.K. Kim, Acta Mater.59 ( 2011) 7787-99. [47] G. Frommeyer, U. Brux, Steel Res. Int. 77 (2006) 627-633. [48] J.D. Yoo, K.-T. Park, Mater. Sci. Eng. A 496 (2008) 417-24. [49] H. Beladi, P. Cizek, P.D. Hodgson, Metall. Mater. Trans. A 40 (2009) 1175-89. [50] X. Zhang, T. Sawaguchi, K. Ogawa, F. Yin, X. Zhao, Philos. Mag. 91 (2011) 4410-26.
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[51] S. Poulat, B. Decamps, L. Priester, Philos. Mag. A 77 (1998) 1381-97. [52] S. Poulat, B. Decamps, L. Priester, Philos. Mag. A 79 (1999) 2655-80. [53] N.E.Paton, W.A. Backofen, Metall. Trans. A 1 (1970) 2841-45. [54] M.H. Yoo, Metall. Trans. A 12 (1981) 409-18. [55] M.J. Marcinkowski, D.S. Miller, Philos. Mag. 6 (1961) 871-93. [56] S.M. Copley, B.H. Kear, Acta Metall. 16 (1968) 227-31. [57] D. Goodchild, W.T. Roberts, D.V. Wilson, Acta Metall. 18 (1970) 1137-45. [58] I. Karaman, H. Sehitoglu, H.J. Maier, Y.I. Chumlyakov, Acta Mater. 49 (2001) 3919-33. [59] K. Ishida, H. Ohtani, N. Satoh, R. Kainuma, T. Nishizawa, ISIJ Int. 30 (1990) 680-6. [60] T.W.Kim, Ph.D. thesis. 1993, KAIST, Korea.
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Figure Captions Fig. 1. schematic diagram of the entire process to produce primary specimens. Fig. 2. (a) The true stress–true strain response of the A-type material at different temperature under strain rate of 0.01 s-1, (b) the variation of corresponding flow stress of 0.5 true strain with temperature. Fig. 3.The effect of temperature on the strain-induced martensite formation in the experimental steel. Fig. 4. (a) The variation of strain hardening rate with true strain at 25 °C is superimposed on its corresponding true stress-true strain curve, (b) the volume fraction of strain-induced martensite against true strain in type-A specimen. Fig. 5. The true stress–true strain response of the A-type specimen at 600°C. Fig. 6. The variation of strain hardening rate with true strain at 800°C is superimposed on its corresponding true stress-true strain curve, (a) B-type specimen with grain size of 500±50μm, (b) A-type specimen with grain size of 40±5 μm. Fig. 7. The true stress–true strain response of the specimen at 1000 °C, (a) B-type specimen under strain rate of 0.01 s-1, (b) A-type specimen under strain rate of 0.01 s-1. Fig. 8. The microstructure of the specimen in, (a) A-type condition, and (b) B-type condition. Fig. 9. The SEM microstructures of the A-type specimen after compression at 25°C: (a) after straining to 0.15, composing of martensite platelets, (b) after straining to 0.30, (c) after straining to 0.60, martensite intersections are seen, (d) outer surface, indicating uneven surface like ripple structure. Fig. 10. EBSD micrograph of compressed sample at 25°C after straining to 0.3: (a) band contrast, (b) map color. blue and yellow displays ά and
martensite, respectively.
Fig. 11. The microstructure of A-type specimens after straining to 0.60 at: (a)100°C, indicates the austenite partitioning by deformation-induced martensite, (b) 300°C, and (c) 400°C merely indicates the mechanical twinning, (d) 600°C, containing a banded-like structure links to activated secondary deformation twin, (e) EBSD image of compressed sample at 300°C showing deformation twinning. Fig. 12. The microstructure of A-type specimens after straining to 0.60 at (a) 700 °C,(b) 800°C, the high temperature banded structure links to mechanical twins/deformation band are seen, (c) 900°C, (d) the specimen outer surface after hot compression at 800°C. Fig. 13. The microstructure of B-type specimens after straining to 0.60 at: (a) 700°C, indicating high temperature mechanical twinning/deformation band, (b) 800°C, (c) 900°C, (d) and (e) outer surface after hot compression at 800°C and 900°C, respectively. Fig. 14. SEM-EBSD micrograph of A-type specimens after hot compression at 800°C to 0.60 true strain under strain rate of 0.01 s-1 : (a) and (c) band contrast image, (b) and (d) color map. Black and red lines represent high angle grain boundaries (θ> 15°) and twin boundaries (55°), respectively. Fig. 15. EBSD micrograph of the B-type specimen after hot compression at 900°C to 0.60 true strain under strain rate of 0.01 s-1 . Fig. 16. TEM micrograph of the B-type specimen after hot compression to 0.60 true strain at 900°C; taken from an area of the banded-like structure. The inset is the selected area diffraction pattern of the corresponding region. The
17
rings, indicated by A, are the reflections of gamma-austenite. The spots, indicated by B, can be attributed to the diffractions from single-crystalline cubic AlFe3C phase. The diffraction spots, indicated by C, could not be indexed. Fig. 17. The XRD patterns: (a) the A-type specimen before hot compression, (b) the A-type specimen after hot compression at 800°C under strain rate of 0.01 s-1.
18
Figure(s)
Fig. 1. schematic chematic diagram of the entire process to produce primary specimens. specimens
1200 Compressive Strenght (MPa)
(a)
(b)
1000 800 600 400 200 0 0
200
400
600
800
1000
Temperature (°C)
Fig. 2. (a) The true stress–true true strain response of the A-type material at different temperature under strain rate of 0.01 s-1, (b) the variation of corresponding flow stress of 0.5 true strain with temperature.
Fig. 3.The effect of temperature on the strain-induced martensite formation in the experimental steel.
(a)
(b)
Fig. 4. (a) The variation of strain hardening rate with true strain at 25 °C is superimposed on its corresponding true stress-true strain curve, (b) the volume fraction of strain-induced martensite against true strain in type-A specimen.
Fig. 5. The true stress–true strain response of the A-type specimen at 600°C.
(b)
(a)
Fig. 6. The variation of strain hardening rate with true strain at 800°C is superimposed on its corresponding true stress-true strain curve, (a) B-type specimen with grain size of 500±50µm, (b) A-type specimen with grain size of 40±5 µm.
(a)
(b)
Fig. 7. The true stress–true strain response of the specimen at 1000 °C, (a) B-type specimen under strain rate of 0.01 s-1, (b) A-type specimen under strain rate of 0.01 s-1.
(a)
(b)
200 µm
Fig. 8. The microstructure of the specimen in, (a) A-type condition, and (b) B-type condition.
(a)
(b)
(c)
(d)
1 cm
A specimen after compression at 25°C: (a) after straining to 0.1 0.15, Fig. 9. The SEM microstructures of the A-type composing of martensite platelets, (b)) after straining to 0.30, 0.3 (c) after straining to 0.60, martensite intersections are seen, (d) outer surface, indicating uneven surface like ripple structure. structure
(a)
(b)
Fig. 10. EBSD micrograph of compressed sample at 25°C after straining to 0.3: (a) band contrast, (b) map color. blue and yellow displays ά and Ԑ martensite, respectively.
(a)
(b)
(c)
(d)
(e)
Fig. 11. The microstructure of A-type type specimens after straining to 0.60 at: (a)100°C, indicates the austenite partitioning by deformation-induced induced martensite, martensite (b) 300°C, and (c) 400°C merely indicates the mechanical twinning twinning, (d) 600°C, containing a banded-like structure links to activated activat secondary deformation twin, (e) EBSD image of compressed sample at 300°C showing deformation twinning.
(a)
(b)
(c)
(d)
Fig. 12. The microstructure of A-type specimens after straining to 0.60 at (a) 700 °C,(b) 800°C, the high temperature banded structure links to mechanical twins/deformation band are seen, (c) 900°C, (d) the specimen outer surface after hot compression at 800°C.
(b)
(a)
(c)
(d)
(e)
1 cm
Fig. 13. The microstructure of B-type type specimens after straining to 0.60 at:: (a) 700°C, indicating high temperature mechanical twinning/deformation band,, (b) 800°C, 800 (c) 900°C, (d) and (e) outer surface after hot compression at 800°C and 900°C, respectively.
(a)
(b)
Grain boundary serration
Partial dynamic recrystallization
20 µm
Fig. 14. SEM-EBSD micrograph of A-type specimens after hot compression at 800°C to 0.60 true strain under strain rate of 0.01 s-1 : (a) band contrast, (b) color map. Black and red lines represent high angle grain boundaries (θ> 15°) and twin boundaries (55°), respectively.
Twins
Twin-like structure
300 µm
Fig. 15. EBSD micrograph of the B-type specimen after hot compression at 900°C to 0.60 true strain under strain rate of 0.01 s-1.
Fig. 16. TEM micrograph of the B-type specimen after hot compression to 0.60 true strain at 900°C; taken from an area of the banded-like structure. The inset is the selected area diffraction pattern of the corresponding region. The rings, indicated by A, are the reflections of gamma-austenite. The spots, indicated by B, can be attributed to the diffractions from single-crystalline cubic AlFe3C phase. The diffraction spots, indicated by C, could not be indexed.
(a)
(b)
Fig. 17. The XRD patterns: (a) the A-type specimen before hot compression, (b) the A-type specimen after hot compression at 800°C under strain rate of 0.01 s-1.