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Strain-induced phase transformation of an Mn-Al alloy during hot compression
Materials Science & Engineering A 751 (2019) 271–282
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Strain-induced phase transformation of an Mn-Al alloy during hot compression
T
H. Dehghana, S.A. Seyyed Ebrahimia, , M. Noskob ⁎
a b
Advanced Magnetic Materials Research Center, School of Metallurgy and Materials, College of Engineering, University of Tehran, Tehran 1439957131, Iran Institute of Materials and Machine Mechanics, Slovak Academy of Sciences, Dúbravská cesta 9/6319, 845 13 Bratislava, Slovak Republic
ARTICLE INFO
ABSTRACT
Keywords: Mn-Al permanent magnet alloy Rare earth free Hot deformation Strain-Induced phase transformation Processing map Bulk Ferro magnetic property
The high-temperature deformation of an Mn-Al alloy with a chemical composition of Mn51Al47C2 was assessed using hot compression testing at the temperature range of 600–800 °C and the strain rates ranging 0.001 s−1 to 1 s−1. Optical microscopy (OM), scanning electron microscopy (SEM), electron backscattering diffraction (EBSD) and X-ray diffraction (XRD) analyses were implemented to characterize the samples. It was found that during deformation, the τ phase dynamically transformed into γ2+ β + ε phases and the transformation temperature reduced from the equilibrium value of about 790 °C down to < 750 °C. In such temperature range, considering the lower strength of the produced phases, the dynamic strain-induced phase transformations caused a flow softening in the stress-strain curves. However, deformation at temperatures lower than 700 °C, in the state of τ phase stability, led to dynamic recrystallization which was responsible for flow softening. The standard approach was applied to calculate the processing maps which represented an optimal working region in the temperature range of 650 °C up to 700 °C with the strain rates lower than 0.1 s−1 in order to avoid flow localization and dynamic phase transformations.
1. Introduction The Mn-Al permanent magnet alloys have been the subject of researches based on their potentially better magnetic properties than hard ferrites and also more affordable prices than rare earth magnets [1–3]. The ferromagnetic L10 phase τ-Mn-Al is the only magnetic phase in the Mn-Al system [4,5]. A high theoretical magneto-crystalline anisotropy field about 38 kOe, saturation magnetisation about 144 emu/g, and (BH) max of 12.64 MGOe have been reported for this alloy [6]. This meta-stable phase can be stabilized by adding small amounts of carbon that leads to increasing its coercivity, as well [3,7]. It has been shown that the magnetic properties strongly depend on the microstructural characteristics, including phase transformations and stability of different phases, grain structure, and crystallographic texture. Considering this matter, several methods have been introduced to produce and/or improve the magnetic properties of Mn-Al permanent magnet alloy, such as mechanical alloying [8–10], mechanical milling [11–15], melt spinning [16,17], gas atomization [17], and consolidation via sintering [18,19] or ECAE [20], etc. However, some disadvantages can restrict the application of these processing technologies. For instance, one of the essential steps of the fabrication process in powder metallurgy approach is powder consolidation which can deteriorate the magnetic ⁎
properties of the powder. It should also be noted that the permanent magnets commonly are demanded in bulk form so techniques based on bulk deformation have been investigated in various researches. Ohtani et al. [7] developed an anisotropic permanent magnet in Mn-Al-C system by warm-plastic deformation of the material in τ-phase. Valiev et al. [21] studied the effect of prior hot working on the workability of the same alloy and concluded that the initial τ-phase structure has an important influence on the workability which can be modified by the primary hot working of the parent ε-phase. Recently, Bittner et al. [22,23] have performed detailed investigations on the structural defects of τ-phase in Mn-Al-C system in the un-deformed state as transformed from high-temperature ε-phase and in the deformed state. Also, several studies have been carried out recently concerning the effect of microstructural features on magnetic properties of the alloy developed by hot deformation [24–26]. Although it has been shown that during high-temperature deformation the process parameters such as the amount of strain, the strain rate, the temperature, the state of stress and strain, influence greatly on microstructure refinement and consequently the magnetic properties of the final product, a comprehensive research study on the deformation parameters is lacking. Therefore, the aim of the present study is to focus on elaborating the hot-deformation behavior and final
Corresponding author. E-mail address: [email protected] (S.A.S. Ebrahimi).
https://doi.org/10.1016/j.msea.2019.02.082 Received 13 January 2019; Received in revised form 23 February 2019; Accepted 23 February 2019 Available online 25 February 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
Materials Science & Engineering A 751 (2019) 271–282
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Each sample was dwelled about 10 min at the deformation temperature before applying the load to ensure about the uniformity of the temperature throughout the sample. The obtained flow curves were corrected in terms of the effect of friction by the well-accepted formulation [27]. The deformed samples were quenched right after hot compression test and then they were cut along their longitudinal axis. In order to monitor the microstructural evolution using an Olympus optical microscope, the cross-sectioned samples were prepared with common grinding and polishing routes and were etched with a chemical solution of 6% HCl, 3% HNO3, 1% HF, and 90% H2O (volume fraction). The microstructural characterization was carried out using the FEI Quanta 450 scanning electron microscopy equipped with a Bruker (XFlash® 6|10) detector. Also, electron backscattering diffraction analysis (EBSD) was performed to study the grain structure and crystallographic texture with more details, using a field emission scanning electron microscope (FEGSEM, JEOL 7600 F, Japan) equipped with an EBSD detector and analysis system (HKL Nordlys and CHANNEL 5) operating at 10 keV. The EBSD analysis was accomplished near to the center of the compressed sample by considering a quadrant cross-section part. The step size for electron beam scanning was equal to 0.2 µm and an area of 100 µm× 120 µm was indexed. Tango software was utilized to plot the orientation, grain boundary, phase, and pattern quality maps from the EBSD analysis data. Mambo software was used for plotting (inverse) pole figures from EBSD orientation and misorientation data while Salsa software was employed to plot the 2D and 3D orientation distribution functions (ODF), as well. The thermal response of the homogenized alloy during heat treatment was investigated by a simultaneous thermal analyzer (NETZSCH STA 409 PC/PG, Germany) with Al2O3 crucible under argon atmosphere, in the temperature range of RT up to 1200 °C. To identify the corresponded phase states, four homogenized samples were annealed with a heating rate same as the mentioned differential thermal analysis (10 °C/min). These samples were heated in a muffle furnace up to selected temperatures based on the DTA result and immediately were quenched in water. X-ray diffraction measurements were performed on the annealed samples and the deformed samples by the Rigaku Ultima IV diffractometer (Rigaku, Japan) using Cu-Kα radiation (λ = 1.54060 Å).
Fig. 1. Schematic view of the experimental procedure.
microstructure of the bulk ferromagnetic Mn-Al-C alloy using hot compression testing at different temperature and strain rates. 2. Material and methods An Mn-Al alloy with a chemical composition of Mn51Al47C2 was prepared through a controlled atmosphere induction melting process followed by casting into a water-cooled copper mold under pure argon. It is worth noting that after vacuuming and argon purging, titanium getter system was used to purify the chamber from residual oxygen. In order to ensure the homogeneity of the casted alloy and to adjust the composition, the re-melting process was carried out two times. The composition of the alloy was measured using the Inductively Coupled Plasma-Optical Emission Spectroscopy (ICP-OES, 730-ES, Varian, USA) and carbon content was measured using the LECO CS-244 carbon determinator (ASTM E1019). The final ingot was encapsulated in an evacuated quartz tube and was homogenized at 1100 °C for 12 h, and then it was detained at the temperature of 900 °C for 30 min to reduce thermal shock during a consequent quench of the encapsulated ingot in water. To perform hot compression tests (ASTM E209), cylindrical samples with 6 mm diameter and 9 mm height were prepared from the homogenized ingot using electrical discharge machining (EDM) technique. The compression tests were performed at the temperature range of 600–800 °C and strain rates of 10−3 to 1 s−1 up to 50% of initial height by a universal testing machine (INSTRON 8502, USA) equipped with a fully digital and computer-controlled furnace. The schematic view of the experimental procedure is shown in Fig. 1.
3. Results and discussion The flow stress curves obtained from compression tests at various temperatures and strain rates are presented in Fig. 2. According to this
Fig. 2. Flow curves obtained from compression tests at various temperatures and strain rates.
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Fig. 3. Microstructures of (a) the homogenized sample and the samples deformed up to the strain of 0.7 at the temperatures and the strain rates of (b) 600 °C-0.1 s−1, (c) 650 °C-0.001 s−1, (d) 650 °C-0.1 s−1, (e) 700 °C-0.001 s−1, and (f) 700 °C-0.1 s−1.
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Fig. 4. EBSD analysis results for the deformed sample at 650 °C with the strain rate of 0.001 s−1: (a) Grain boundary map, (b) Channeling contrast image, (c) Orientation map. In this map high and low angle boundaries are highlighted (HAGBE (above 15°) are red and LAGBE (3–15°) are white), and (d) inverse pole-figure map.
Fig. 5. Optical, SEM (secondary electron), and EBSD micrographs of the sample deformed at 650 °C with the strain rate of 0.001 s−1.
figure, by increasing the deformation temperature the flow stress is gradually decreased. Moreover, there is a peak stress for each curve which can be attributed to the occurrence of dynamic recrystallization (DRX) and/or dynamic phase transformation and/or flow localization [28]. In other words, during deformation, the flow stress rises up to peak stress and thereafter it decreases to the steady-state stress. According to these flow curves, it seems that by increasing the
deformation temperature and reducing the strain rate, the peak stress and the peak strain are decreased. It is worth mentioning that for two temperatures of 600 °C and 650 °C at the strain rate of 1 s−1, the flow behavior of material was close to brittle and it was not possible to record the repeatable data for these conditions during hot compression testing. This is the main reason that the data for these two testing conditions are not be presented in Fig. 2.
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Fig. 6. Crystallographic micro-texture analysis results related to the EBSD investigation in Fig. 4 as 2D plots of (a) IPF and (b) PF as well as the 3D plots of ODF in (c) color-counter intensity and (d) inside component plots.
Microstructural investigations on the deformed sample at temperatures of 600 °C, 650 °C, and 700 °C reveal that the deformed microstructure is contained of DRXed area while the microstructure of homogenized sample showed rough and highly twinned surface of the τ-phase as presented in Fig. 3. According to Fig. 3(b to f) new recrystallized grains appeared around the previous grains. It has been stated that the dynamic recrystallization generally starts at the old grain boundaries where the higher stored energy, i.e., the dislocation density is provided [28]. Then, new grains nucleate at the boundaries of the growing grains and result in the formation of recrystallized grains band around the old grains (necklace structure). The same mechanism was reported by Bittner et al. [22]. Therefore, flow stress decreases as DRX happens. The deformation at higher strain rates yielded to the lower DRXed grain size; however, as it can be seen from Fig. 3(b, e, and f), some strain localized zones are formed at the DRXed area, i.e., around the old grains, where the flow stress is lower than the other parts of the microstructure. The flow localization can also be evidenced by the higher rate of softening in flow stress curve of the samples deformed at the strain rate of 0.1 s−1 with respect to the lower rates. It is known that, during deformation at high strain rates, a large part of the irreversible plastic work contributes to the heat generation, while the rest is stored as strain energy in the form of internal defects [27]. Hence, in addition to flow localization at these samples, adiabatic heating during deformation leads to a higher reduction in the flow stress.
Low magnification general view from the presence of deformed coarse and recrystallized fine grains besides of each other is presented in EBSD analysis maps of Fig. 4. In these grain boundary and crystallographic orientation plots, the previously claimed necklace mechanism for the occurrence of partial DRX phenomenon during hot compression process can be realistically highlighted. Fig. 5 represents these DRXed grains at the necklace structure in the sample deformed at the temperature of 650 °C with the strain rate of 0.001 s−1 by illustrating the high magnification optical, SEM and EBSD micrographs. The occurrence of pitting during etching of this alloy makes it difficult to reveal the DRXed grains by optical microscopy. Meanwhile, EBSD analysis worked well in revealing such fine-grains. The crystallographic microtextural analysis results which are related to the EBSD analysis of the deformed sample at the temperature of 650 °C under the strain rate of 0.001 s−1 are illustrated in Fig. 6. The {100}, {110}, and {111} inverse pole figure (IPF) and pole figure (PF) plots as the average 2D maps for the scanned area are shown in Fig. 6a and b, respectively. In order to determine the preferred crystallographic orientations, the corresponding 3D orientation distribution function (ODF) tomography is calculated in terms of color-counter intensity and featured components plots and it is demonstrated in Fig. 6c and d, respectively. Based on these estimations regarding crystallographic texture from the deformed and recrystallized regions, a quite strong textural component of {011} < 01¯1> is noted as dominant (with the maximum intensity of about 30). It could be explained by the action of
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Fig. 7. Microstructures of the samples deformed up to the strain of 0.7 at the temperatures and the strain rates of (a) 750 °C-0.1 s−1, (b) 750 °C-0.001 s−1, (c) 800 °C0.1 s−1, and (d) 800 °C-0.001 s−1.
shear banding and strain localization during hot compression testing without any strain-induced phase transformation for such sample as it will be discussed in the following next sections. However, such flow pattern of deformation aided by the operative DRX phenomenon yields to the generation of mentioned shear-like crystallographic textural component in the angle close to 45° for the evaluated location (as addressed in the experimental procedure section with more details) [29]. Deformation at the higher temperatures, i.e., > 750 °C, resulted in the formation of the homogenous multi-phase microstructure (Fig. 7). According to Fig. 7, deformation at higher strain rates and lower temperatures led to the formation of a finer microstructure. As it was shown in Fig. 2, all of the curves have a peak which may indicate the dynamic recrystallization. However, having the higher softening rate in the samples with higher strain rates is in contradiction with DRX characteristics. In other words, by increasing the strain rate, DRX changes to
DRV which leads to the disappearance of peak stress [28]. Therefore, the softening which occurs beyond the peak in these curves may be attributed to adiabatic deformation heating (at strain rates higher than 0.1 s−1) and/or phase transformation. In order to find the phase transformation temperatures, differential thermal analysis (DTA) was employed. Also to recognize the transformations, some samples were heated with the same rate as DTA and they were just quenched in the water. As it can be seen from Fig. 8, there are two endothermic peaks and one exothermic peak which represent the phase transformations at temperatures of 817, 884 and 656 °C respectively. It was stated that in the similar alloy the exothermic peak is related to the formation of the ferromagnetic τ-phase (tetragonal) from the parent as-quenched ε-phase [30]. But here, because of the different producing route (quench with intermediate rate); τ-phase is the starting phase for the samples and as it can be seen in Fig. 8, XRD patterns
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Fig. 8. DTA graph with the rate of 10 °C/min and the corresponded X-ray diffraction patterns along with the related phase diagram.
reflect the existence of τ-phase before and after the broad low-intensity exothermic peak. Moreover, the occurrence temperature of the peak is about 170 °C higher and intensity is much lower than the similar mentioned case. Therefore, more studies are required to reveal the source causing this peak. This phase of τ is metastable and will decompose to the equilibrium phases γ2 (Al8Mn5) and β (cubic) by maintaining at the elevated temperatures [31]. Based on the Al-Mn equilibrium phase diagram, γ2 (Al8Mn5) and some β phase will trans(hexagonal) phase form to γ and then at higher temperature + transformation will occur. X-ray diffraction analysis on the sample heated up to the temperature of 750 °C indicates that τ-phase is still stable. According to the study by Huang and Kuo [30], the addition of C reduces the temperature of + phase transformation and for 0.6 wt% C and more, the + two-phase area will be removed and as a result, the 2 + transformation will occur. They also suggested that, by doping the high carbon content (0.6 wt%), τ-phase can be stabilized up to the transformation temperature for 2 + which results in a high-intensity endothermic peak in DTA graph [30]. X-ray diffraction analysis on the sample heated up to temperature of 850 °C, indicates that the structure is contained of 2 + + phases. Hence, the high intensity endothermic peak is caused due to the decomposition of τ ( 2 + ) and + phase transformation. Moreover, considering the alloy 2 composition (Mn content equal to51 at%) and Al-Mn equilibrium phase diagram (Fig. 8), it seems that the second endothermic peak in the present study is related to the 2 ( ) phase transformation. The Xray diffraction pattern of the sample heated up to the temperature of 950 °C reflecting the existence of the ε phase, confirms this important matter. Considering these results, it is expected that the τ single phase remains stable through the structure during deformation at the mentioned temperature range, up to the temperature of 800 °C. However, it is an expectation and must be checked at follow with experimental testing.
XRD analyses of the deformed samples at the temperatures of 650 °C up to 800 °C with the strain rates in the range of 0.001 s−1 to 1 s−1 are presented in Fig. 9. As it can be seen, although at the temperature of 700 °C τ-phase is still stable, at the temperature of 750 °C, in all strain rates in the range of 0.001–1 s−1, the microstructure is contained of + phases. Although according to the static thermal analysis 2 + results reported in the previous section the τ-phase was found stable up to the temperature of 800 °C, however, based on these new results it seems that the deformation has led to a decrease in the 2 + phase transformation temperature from 800 °C down to a temperature between 700 °C and 750 °C. Even at extremely high strain rate of 1 s−1, there is not considerable chance for occurrence of strain-induced phase transformation under temperature of 700 °C. Fig. 7a and b, which represent the microstructures of the deformed samples at the temperature 750 °C with the strain rates of 0.1 and 0.001 s−1, are covered by three phases that support again the decrease in the 2 + phase transformation temperature. Similarly, Fig. 7c and d show that the microstructures of the deformed samples at the temperature of 800 °C with strain rates of 0.1 s−1 and 0.001 s−1 is contained of three phases, but the β-phase fraction has been reduced extremely. In accordance with XRD results, the backscattered electron micrograph associated with the EDS line-scan analysis results indicates that the deformed samples at temperatures higher than 700 °C are contained of three phases with different contents of Al and Mn elements (Fig. 10). According to this figure, the black area which has the minimum value of Al element must be β-phase and light gray which has the higher value of Al could be γ2-phase. While the rest one (dark gray) can be τ single phase for the deformed sample at the temperature of 700 °C and ε-phase for the samples deformed at temperatures of 750 °C and 800 °C. It is well-established that the diffusivity can be enhanced by inducing the strain in terms of increasing the dislocation density and number of shear bands that can leads to increasing the number of
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Fig. 9. XRD analysis of the samples deformed at the temperatures of (a-c) 650 °C, (d-f) 700 °C, (g, h) 750 °C, and (i, j) 800 °C with the strain rates of (a, d, g, i) 0.001 s−1, (b, e) 0.1 s−1, and (c, f, h, j) 1 s−1.
different strain rates (0.001 s−1 and 0.1 s−1) are shown in Fig. 11. As it can be seen in this figure, a sharp stress drop occurs between the deformation temperatures of 700 °C and 750 °C. Considering the straininduced transformation temperature, i.e., between 700 °C and 750 °C, this fact implies that the transformed phases are softer than the transforming phase. The appearance of the inflection point in the work hardening rate versus stress plot (θ-σ) can be considered as an applicable method to find whether the phase transformation occurrence proceeds dynamically or not during deformation [33,34]. In the other words, considering the fact that the transformed phases are softer than the transforming phase and transformation mechanism involves nucleation and growth, the curve of work-hardening as a function of stress must have a zero point and an inflection point, as follow.
nucleation sites as well as elevating the driving force for phase transformation [28,32]. They promote in turn the nucleation phenomenon and increase the nucleation rate. Based on the fact that the diffusion rate is influenced by dislocation density, nuclei will grow rapidly and this will make the new grains of different phases to impinge on each other. By increasing the deformation strain rate higher density of defects, i.e., dislocation density and shear bands, can be formed. Therefore, the higher nucleation rate is accomplished with the higher rate of growth and will result in finer transformed phases. This fact can also be deduced from comparing the grain size of the samples deformed at the temperature of 800 °C with the strain rates of 0.1 s−1 and 0.001 s−1 (Fig. 7). As it can be seen from this figure, the deformed sample at higher strain rate is contained of both finer γ2- and ε-phases. Variations of the normalized change in stress at the strain of 0.7, i.e., ( T1 T2 )/ T2 , with respect to the deformation temperature at two
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Fig. 10. Backscattered electron SEM micrographs of the samples deformed at (a) 700 °C, (b) 750 °C, and (C) 800 °C with the strain rate of 0.001 s−1 associated with line scan EDS result for Al (gray) and Mn (red).
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power per unit volume (P) absorbed by the workpiece during plastic flow can be divided into two different parts, the power dissipated by the plastic work which mostly converts into heat (G), and the power dissipated due to the occurrence of the metallurgical mechanisms such as DRX, DRV and phase transformation (J) [35]. In order to find the power dissipation capacity of the material, a dimensionless parameter which is called the efficiency of power dissipation (η) has been introduced, i.e., = 2m / m + 1. It was stated that the deformation processing should be focused on the regions of maximum efficiency of the power dissipation, unless structural instabilities would occur, for example, flow localization. It has been stated that the material can be subjected to the instabilities during flow behavior when the rate of entropy production by the microstructural evolution in the material is lower than the applied rate of entropy. Based on this fact, the following criterion, ξ, was proposed to find the instability during deformation [36]:
=
Fig. 12. Variation of the work hardening rate with the stress.
= 0 And
=
=0 =p
(2)
The negative value of the instability coefficient indicates the flow instability and the location of the microstructural instability region can be found by mapping the instability parameter versus temperature and strain rate for a certain strain. The relationship between ln and ln is obtained by cubic curve fitting, and then the strain rate sensitivity exponent m can be obtained based on m ( ) = ln / ln T, equation for each certain temperature and strain as the function of strain rate. Fig. 13 represents the power dissipation contour map obtained through plotting the parameter η against the temperature and strain rate. As it was mentioned, the maximum η represents the optimal processing window. The higher value of η means that more power is dissipated due to the occurrence of dynamic microstructural evolution. As it can be seen from Fig. 13, η parameter is increased by increasing the temperature which is maximum for temperatures higher than 750 °C with strain rates of lower than 0.01 s−1. As it is shown on the map, this area (represented in dark blue) is the deformation conditions that lead to the dynamic phase transformation. In other words, the power was dissipated mainly through the phase transformation. The instability map obtained by plotting the instability coefficient against the temperature and strain rate is super-imposed on the more power dissipation map. The negative values of the instability coefficient and red area indicate the flow instability region which is located at the lower temperature region and the higher strain rates. As it was mentioned previously, the flow localization can occur in this area as the result of DRX around the grains. Based on the XRD results and the microstructural investigations, the blue areas reflecting the dynamic phase transformation condition are super-imposed on the instability map in order to find the optimal area for deformation. According to the microstructures presented in Figs. 10 and 13, compressive deformation at temperatures over than 700 °C leads to starting the τ to γ2 + β and subsequently to ε phase transformation. Considering this map, the optimal working area, in which the instability and the phase transformation are avoided is a narrow window at the temperature range between 650 °C and 700 °C and the strain rate lower than 0.1 s−1.
Fig. 11. Variation of the normalized change in saturation stress for change in the deformation temperature.
= c
ln(m /(m + 1)) +m ln
(1)
Fig. 12 shows the variations of work hardening rate regarding the stress for different strain rates. Based on Fig. 12, it can be found that for each curve, all the inflection points lay on a single straight line. It indicates that the rate of microstructure evolution is constant at the critical state. According to the presented evidence, it seems that the dynamic softening observed at temperatures of 750 °C and 800 °C is attributed to the dynamic phase transformation mechanisms. This phenomenon can be associated with adiabatic heating effect for high strain rates which accelerates the rate of softening during deformation. It has been shown that, during high temperature deformation, the
4. Conclusions The Mn-Al alloy with the chemical composition of Mn51Al47C2 was subjected to the compressive deformation at elevated temperatures with the wide range of strain rates and the following results were obtained:
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Fig. 13. Processing map for the strain of 0.7 with the corresponded optical microstructures, noting the red and the blue areas are corresponded to the instability zone and the phase transformation, respectively.
1) It was observed that during deformation, τ-phase dynamically transformed to γ2+ β + ε at a temperature lower than 750 °C. 2) Deformation at temperatures lower than 700 °C, where τ-phase was stable, led to the dynamic recrystallization around the initial grains which in turn resulted in the flow curve softening. 3) Based on the lower strength of the produced phases, the dynamic strain-induced transformation caused a flow softening in the stressstrain curves. 4) As a main finding after elaboration of different temperatures and strain rates during hot compression testing, the flow localization phenomenon and dynamic phase transformation can be prohibited in the ranges of 650–700 °C for temperature and 0.001 s−1 to 0.1 s−1 for strain rate.
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Acknowledgement This work was also supported by Slovak Foundation VEGA grant 2_0158_16, and by grant APVV-14–0936. The authors wish to thank Dr. Lubomir Orovcik (Institute of Materials and Machine Mechanics, Slovak Academy of Sciences, Dúbravská cesta 9/6319, 845 13, Bratislava, Slovak Republic) for helping in RBEI imaging. Data availability The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study. References [1] J.M.D. Coey, Permanent magnets: plugging the gap, Scr. Mater. 67 (2012) 524–529. [2] J.M.D. Coey, New permanent magnets; manganese compounds, J. Phys. Condens. Matter 26 (2014). [3] M.A. Bohlmann, J.C. Koo, J.H. Wise, Mn-Al-C for permanent magnets (invited), J. Appl. Phys. 52 (1981) 2542–2543. [4] H. Kono, On the ferromagnetic phase in manganese-aluminum system, J. Phys. Soc. Jpn. 13 (1958) 1444–1451. [5] A.J.J. Koch, P. Hokkeling, M.G. v. d. Steeg, K.J. de Vos, New material for permanent magnets on a base of Mn and Al, J. Appl. Phys. 31 (1960) S75–S77. [6] J.H. Park, Y.K. Hong, S. Bae, J.J. Lee, et al., Saturation magnetization and crystalline anisotropy calculations for MnAl permanent magnet, J. Appl. Phys. 107 (2010).
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recrystallization under hot, cold and severe plastic deformation conditions, Prog. Mater. Sci. 60 (2014) 130–207. J. Huang, P. Kuo, Effect of quenching temperature on the formation and magnetic properties of the ferromagnetic r-phase in Mn-Al-C alloys, Mater. Sci. Eng. B 20 (1993) 292–297. W.H. Dreizler, A. Menth, Transformation Kinetics of the Ferromagnetic AI lo, IEEE Trans. Magn. 16 (1980) 534–536. J. Choi, D. Seo, J. Lee, K. Um, et al., Formation of ultrafine ferrite by strain-induced dynamic transformation in plain low carbon steel, ISIJ Int. 43 (2003) 746–754. E.I. Poliak, J.J. Jonass, A one parameter approach to determining the critical
conditions for the initiation of dynamic recrystallization, Acta Mater. 44 (1996) 127–136. [34] H. Mirzadeh, A. Najafizadeh, Prediction of the critical conditions for initiation of dynamic recrystallization, Mater. Des. 31 (2010) 1174–1179. [35] V. Gopala Krishna, Y.V.R.K. Prasad, N.C. Birla, G. Sambasiva Rao, Processing map for the hot working of near-α titanium alloy 685, J. Mater. Process. Technol. 71 (1997) 377–383. [36] Y.V.R.K. Prasad, T. Seshacharyulu, Processing maps for hot working of titanium alloys, Mater. Sci. Eng. A 243 (1998) 82–88.
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Strain-induced phase transformation of an Mn-Al alloy during hot compression
T
H. Dehghana, S.A. Seyyed Ebrahimia, , M. Noskob ⁎
a b
Advanced Magnetic Materials Research Center, School of Metallurgy and Materials, College of Engineering, University of Tehran, Tehran 1439957131, Iran Institute of Materials and Machine Mechanics, Slovak Academy of Sciences, Dúbravská cesta 9/6319, 845 13 Bratislava, Slovak Republic
ARTICLE INFO
ABSTRACT
Keywords: Mn-Al permanent magnet alloy Rare earth free Hot deformation Strain-Induced phase transformation Processing map Bulk Ferro magnetic property
The high-temperature deformation of an Mn-Al alloy with a chemical composition of Mn51Al47C2 was assessed using hot compression testing at the temperature range of 600–800 °C and the strain rates ranging 0.001 s−1 to 1 s−1. Optical microscopy (OM), scanning electron microscopy (SEM), electron backscattering diffraction (EBSD) and X-ray diffraction (XRD) analyses were implemented to characterize the samples. It was found that during deformation, the τ phase dynamically transformed into γ2+ β + ε phases and the transformation temperature reduced from the equilibrium value of about 790 °C down to < 750 °C. In such temperature range, considering the lower strength of the produced phases, the dynamic strain-induced phase transformations caused a flow softening in the stress-strain curves. However, deformation at temperatures lower than 700 °C, in the state of τ phase stability, led to dynamic recrystallization which was responsible for flow softening. The standard approach was applied to calculate the processing maps which represented an optimal working region in the temperature range of 650 °C up to 700 °C with the strain rates lower than 0.1 s−1 in order to avoid flow localization and dynamic phase transformations.
1. Introduction The Mn-Al permanent magnet alloys have been the subject of researches based on their potentially better magnetic properties than hard ferrites and also more affordable prices than rare earth magnets [1–3]. The ferromagnetic L10 phase τ-Mn-Al is the only magnetic phase in the Mn-Al system [4,5]. A high theoretical magneto-crystalline anisotropy field about 38 kOe, saturation magnetisation about 144 emu/g, and (BH) max of 12.64 MGOe have been reported for this alloy [6]. This meta-stable phase can be stabilized by adding small amounts of carbon that leads to increasing its coercivity, as well [3,7]. It has been shown that the magnetic properties strongly depend on the microstructural characteristics, including phase transformations and stability of different phases, grain structure, and crystallographic texture. Considering this matter, several methods have been introduced to produce and/or improve the magnetic properties of Mn-Al permanent magnet alloy, such as mechanical alloying [8–10], mechanical milling [11–15], melt spinning [16,17], gas atomization [17], and consolidation via sintering [18,19] or ECAE [20], etc. However, some disadvantages can restrict the application of these processing technologies. For instance, one of the essential steps of the fabrication process in powder metallurgy approach is powder consolidation which can deteriorate the magnetic ⁎
properties of the powder. It should also be noted that the permanent magnets commonly are demanded in bulk form so techniques based on bulk deformation have been investigated in various researches. Ohtani et al. [7] developed an anisotropic permanent magnet in Mn-Al-C system by warm-plastic deformation of the material in τ-phase. Valiev et al. [21] studied the effect of prior hot working on the workability of the same alloy and concluded that the initial τ-phase structure has an important influence on the workability which can be modified by the primary hot working of the parent ε-phase. Recently, Bittner et al. [22,23] have performed detailed investigations on the structural defects of τ-phase in Mn-Al-C system in the un-deformed state as transformed from high-temperature ε-phase and in the deformed state. Also, several studies have been carried out recently concerning the effect of microstructural features on magnetic properties of the alloy developed by hot deformation [24–26]. Although it has been shown that during high-temperature deformation the process parameters such as the amount of strain, the strain rate, the temperature, the state of stress and strain, influence greatly on microstructure refinement and consequently the magnetic properties of the final product, a comprehensive research study on the deformation parameters is lacking. Therefore, the aim of the present study is to focus on elaborating the hot-deformation behavior and final
Corresponding author. E-mail address: [email protected] (S.A.S. Ebrahimi).
https://doi.org/10.1016/j.msea.2019.02.082 Received 13 January 2019; Received in revised form 23 February 2019; Accepted 23 February 2019 Available online 25 February 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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Each sample was dwelled about 10 min at the deformation temperature before applying the load to ensure about the uniformity of the temperature throughout the sample. The obtained flow curves were corrected in terms of the effect of friction by the well-accepted formulation [27]. The deformed samples were quenched right after hot compression test and then they were cut along their longitudinal axis. In order to monitor the microstructural evolution using an Olympus optical microscope, the cross-sectioned samples were prepared with common grinding and polishing routes and were etched with a chemical solution of 6% HCl, 3% HNO3, 1% HF, and 90% H2O (volume fraction). The microstructural characterization was carried out using the FEI Quanta 450 scanning electron microscopy equipped with a Bruker (XFlash® 6|10) detector. Also, electron backscattering diffraction analysis (EBSD) was performed to study the grain structure and crystallographic texture with more details, using a field emission scanning electron microscope (FEGSEM, JEOL 7600 F, Japan) equipped with an EBSD detector and analysis system (HKL Nordlys and CHANNEL 5) operating at 10 keV. The EBSD analysis was accomplished near to the center of the compressed sample by considering a quadrant cross-section part. The step size for electron beam scanning was equal to 0.2 µm and an area of 100 µm× 120 µm was indexed. Tango software was utilized to plot the orientation, grain boundary, phase, and pattern quality maps from the EBSD analysis data. Mambo software was used for plotting (inverse) pole figures from EBSD orientation and misorientation data while Salsa software was employed to plot the 2D and 3D orientation distribution functions (ODF), as well. The thermal response of the homogenized alloy during heat treatment was investigated by a simultaneous thermal analyzer (NETZSCH STA 409 PC/PG, Germany) with Al2O3 crucible under argon atmosphere, in the temperature range of RT up to 1200 °C. To identify the corresponded phase states, four homogenized samples were annealed with a heating rate same as the mentioned differential thermal analysis (10 °C/min). These samples were heated in a muffle furnace up to selected temperatures based on the DTA result and immediately were quenched in water. X-ray diffraction measurements were performed on the annealed samples and the deformed samples by the Rigaku Ultima IV diffractometer (Rigaku, Japan) using Cu-Kα radiation (λ = 1.54060 Å).
Fig. 1. Schematic view of the experimental procedure.
microstructure of the bulk ferromagnetic Mn-Al-C alloy using hot compression testing at different temperature and strain rates. 2. Material and methods An Mn-Al alloy with a chemical composition of Mn51Al47C2 was prepared through a controlled atmosphere induction melting process followed by casting into a water-cooled copper mold under pure argon. It is worth noting that after vacuuming and argon purging, titanium getter system was used to purify the chamber from residual oxygen. In order to ensure the homogeneity of the casted alloy and to adjust the composition, the re-melting process was carried out two times. The composition of the alloy was measured using the Inductively Coupled Plasma-Optical Emission Spectroscopy (ICP-OES, 730-ES, Varian, USA) and carbon content was measured using the LECO CS-244 carbon determinator (ASTM E1019). The final ingot was encapsulated in an evacuated quartz tube and was homogenized at 1100 °C for 12 h, and then it was detained at the temperature of 900 °C for 30 min to reduce thermal shock during a consequent quench of the encapsulated ingot in water. To perform hot compression tests (ASTM E209), cylindrical samples with 6 mm diameter and 9 mm height were prepared from the homogenized ingot using electrical discharge machining (EDM) technique. The compression tests were performed at the temperature range of 600–800 °C and strain rates of 10−3 to 1 s−1 up to 50% of initial height by a universal testing machine (INSTRON 8502, USA) equipped with a fully digital and computer-controlled furnace. The schematic view of the experimental procedure is shown in Fig. 1.
3. Results and discussion The flow stress curves obtained from compression tests at various temperatures and strain rates are presented in Fig. 2. According to this
Fig. 2. Flow curves obtained from compression tests at various temperatures and strain rates.
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Fig. 3. Microstructures of (a) the homogenized sample and the samples deformed up to the strain of 0.7 at the temperatures and the strain rates of (b) 600 °C-0.1 s−1, (c) 650 °C-0.001 s−1, (d) 650 °C-0.1 s−1, (e) 700 °C-0.001 s−1, and (f) 700 °C-0.1 s−1.
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Fig. 4. EBSD analysis results for the deformed sample at 650 °C with the strain rate of 0.001 s−1: (a) Grain boundary map, (b) Channeling contrast image, (c) Orientation map. In this map high and low angle boundaries are highlighted (HAGBE (above 15°) are red and LAGBE (3–15°) are white), and (d) inverse pole-figure map.
Fig. 5. Optical, SEM (secondary electron), and EBSD micrographs of the sample deformed at 650 °C with the strain rate of 0.001 s−1.
figure, by increasing the deformation temperature the flow stress is gradually decreased. Moreover, there is a peak stress for each curve which can be attributed to the occurrence of dynamic recrystallization (DRX) and/or dynamic phase transformation and/or flow localization [28]. In other words, during deformation, the flow stress rises up to peak stress and thereafter it decreases to the steady-state stress. According to these flow curves, it seems that by increasing the
deformation temperature and reducing the strain rate, the peak stress and the peak strain are decreased. It is worth mentioning that for two temperatures of 600 °C and 650 °C at the strain rate of 1 s−1, the flow behavior of material was close to brittle and it was not possible to record the repeatable data for these conditions during hot compression testing. This is the main reason that the data for these two testing conditions are not be presented in Fig. 2.
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Fig. 6. Crystallographic micro-texture analysis results related to the EBSD investigation in Fig. 4 as 2D plots of (a) IPF and (b) PF as well as the 3D plots of ODF in (c) color-counter intensity and (d) inside component plots.
Microstructural investigations on the deformed sample at temperatures of 600 °C, 650 °C, and 700 °C reveal that the deformed microstructure is contained of DRXed area while the microstructure of homogenized sample showed rough and highly twinned surface of the τ-phase as presented in Fig. 3. According to Fig. 3(b to f) new recrystallized grains appeared around the previous grains. It has been stated that the dynamic recrystallization generally starts at the old grain boundaries where the higher stored energy, i.e., the dislocation density is provided [28]. Then, new grains nucleate at the boundaries of the growing grains and result in the formation of recrystallized grains band around the old grains (necklace structure). The same mechanism was reported by Bittner et al. [22]. Therefore, flow stress decreases as DRX happens. The deformation at higher strain rates yielded to the lower DRXed grain size; however, as it can be seen from Fig. 3(b, e, and f), some strain localized zones are formed at the DRXed area, i.e., around the old grains, where the flow stress is lower than the other parts of the microstructure. The flow localization can also be evidenced by the higher rate of softening in flow stress curve of the samples deformed at the strain rate of 0.1 s−1 with respect to the lower rates. It is known that, during deformation at high strain rates, a large part of the irreversible plastic work contributes to the heat generation, while the rest is stored as strain energy in the form of internal defects [27]. Hence, in addition to flow localization at these samples, adiabatic heating during deformation leads to a higher reduction in the flow stress.
Low magnification general view from the presence of deformed coarse and recrystallized fine grains besides of each other is presented in EBSD analysis maps of Fig. 4. In these grain boundary and crystallographic orientation plots, the previously claimed necklace mechanism for the occurrence of partial DRX phenomenon during hot compression process can be realistically highlighted. Fig. 5 represents these DRXed grains at the necklace structure in the sample deformed at the temperature of 650 °C with the strain rate of 0.001 s−1 by illustrating the high magnification optical, SEM and EBSD micrographs. The occurrence of pitting during etching of this alloy makes it difficult to reveal the DRXed grains by optical microscopy. Meanwhile, EBSD analysis worked well in revealing such fine-grains. The crystallographic microtextural analysis results which are related to the EBSD analysis of the deformed sample at the temperature of 650 °C under the strain rate of 0.001 s−1 are illustrated in Fig. 6. The {100}, {110}, and {111} inverse pole figure (IPF) and pole figure (PF) plots as the average 2D maps for the scanned area are shown in Fig. 6a and b, respectively. In order to determine the preferred crystallographic orientations, the corresponding 3D orientation distribution function (ODF) tomography is calculated in terms of color-counter intensity and featured components plots and it is demonstrated in Fig. 6c and d, respectively. Based on these estimations regarding crystallographic texture from the deformed and recrystallized regions, a quite strong textural component of {011} < 01¯1> is noted as dominant (with the maximum intensity of about 30). It could be explained by the action of
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Fig. 7. Microstructures of the samples deformed up to the strain of 0.7 at the temperatures and the strain rates of (a) 750 °C-0.1 s−1, (b) 750 °C-0.001 s−1, (c) 800 °C0.1 s−1, and (d) 800 °C-0.001 s−1.
shear banding and strain localization during hot compression testing without any strain-induced phase transformation for such sample as it will be discussed in the following next sections. However, such flow pattern of deformation aided by the operative DRX phenomenon yields to the generation of mentioned shear-like crystallographic textural component in the angle close to 45° for the evaluated location (as addressed in the experimental procedure section with more details) [29]. Deformation at the higher temperatures, i.e., > 750 °C, resulted in the formation of the homogenous multi-phase microstructure (Fig. 7). According to Fig. 7, deformation at higher strain rates and lower temperatures led to the formation of a finer microstructure. As it was shown in Fig. 2, all of the curves have a peak which may indicate the dynamic recrystallization. However, having the higher softening rate in the samples with higher strain rates is in contradiction with DRX characteristics. In other words, by increasing the strain rate, DRX changes to
DRV which leads to the disappearance of peak stress [28]. Therefore, the softening which occurs beyond the peak in these curves may be attributed to adiabatic deformation heating (at strain rates higher than 0.1 s−1) and/or phase transformation. In order to find the phase transformation temperatures, differential thermal analysis (DTA) was employed. Also to recognize the transformations, some samples were heated with the same rate as DTA and they were just quenched in the water. As it can be seen from Fig. 8, there are two endothermic peaks and one exothermic peak which represent the phase transformations at temperatures of 817, 884 and 656 °C respectively. It was stated that in the similar alloy the exothermic peak is related to the formation of the ferromagnetic τ-phase (tetragonal) from the parent as-quenched ε-phase [30]. But here, because of the different producing route (quench with intermediate rate); τ-phase is the starting phase for the samples and as it can be seen in Fig. 8, XRD patterns
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Fig. 8. DTA graph with the rate of 10 °C/min and the corresponded X-ray diffraction patterns along with the related phase diagram.
reflect the existence of τ-phase before and after the broad low-intensity exothermic peak. Moreover, the occurrence temperature of the peak is about 170 °C higher and intensity is much lower than the similar mentioned case. Therefore, more studies are required to reveal the source causing this peak. This phase of τ is metastable and will decompose to the equilibrium phases γ2 (Al8Mn5) and β (cubic) by maintaining at the elevated temperatures [31]. Based on the Al-Mn equilibrium phase diagram, γ2 (Al8Mn5) and some β phase will trans(hexagonal) phase form to γ and then at higher temperature + transformation will occur. X-ray diffraction analysis on the sample heated up to the temperature of 750 °C indicates that τ-phase is still stable. According to the study by Huang and Kuo [30], the addition of C reduces the temperature of + phase transformation and for 0.6 wt% C and more, the + two-phase area will be removed and as a result, the 2 + transformation will occur. They also suggested that, by doping the high carbon content (0.6 wt%), τ-phase can be stabilized up to the transformation temperature for 2 + which results in a high-intensity endothermic peak in DTA graph [30]. X-ray diffraction analysis on the sample heated up to temperature of 850 °C, indicates that the structure is contained of 2 + + phases. Hence, the high intensity endothermic peak is caused due to the decomposition of τ ( 2 + ) and + phase transformation. Moreover, considering the alloy 2 composition (Mn content equal to51 at%) and Al-Mn equilibrium phase diagram (Fig. 8), it seems that the second endothermic peak in the present study is related to the 2 ( ) phase transformation. The Xray diffraction pattern of the sample heated up to the temperature of 950 °C reflecting the existence of the ε phase, confirms this important matter. Considering these results, it is expected that the τ single phase remains stable through the structure during deformation at the mentioned temperature range, up to the temperature of 800 °C. However, it is an expectation and must be checked at follow with experimental testing.
XRD analyses of the deformed samples at the temperatures of 650 °C up to 800 °C with the strain rates in the range of 0.001 s−1 to 1 s−1 are presented in Fig. 9. As it can be seen, although at the temperature of 700 °C τ-phase is still stable, at the temperature of 750 °C, in all strain rates in the range of 0.001–1 s−1, the microstructure is contained of + phases. Although according to the static thermal analysis 2 + results reported in the previous section the τ-phase was found stable up to the temperature of 800 °C, however, based on these new results it seems that the deformation has led to a decrease in the 2 + phase transformation temperature from 800 °C down to a temperature between 700 °C and 750 °C. Even at extremely high strain rate of 1 s−1, there is not considerable chance for occurrence of strain-induced phase transformation under temperature of 700 °C. Fig. 7a and b, which represent the microstructures of the deformed samples at the temperature 750 °C with the strain rates of 0.1 and 0.001 s−1, are covered by three phases that support again the decrease in the 2 + phase transformation temperature. Similarly, Fig. 7c and d show that the microstructures of the deformed samples at the temperature of 800 °C with strain rates of 0.1 s−1 and 0.001 s−1 is contained of three phases, but the β-phase fraction has been reduced extremely. In accordance with XRD results, the backscattered electron micrograph associated with the EDS line-scan analysis results indicates that the deformed samples at temperatures higher than 700 °C are contained of three phases with different contents of Al and Mn elements (Fig. 10). According to this figure, the black area which has the minimum value of Al element must be β-phase and light gray which has the higher value of Al could be γ2-phase. While the rest one (dark gray) can be τ single phase for the deformed sample at the temperature of 700 °C and ε-phase for the samples deformed at temperatures of 750 °C and 800 °C. It is well-established that the diffusivity can be enhanced by inducing the strain in terms of increasing the dislocation density and number of shear bands that can leads to increasing the number of
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Fig. 9. XRD analysis of the samples deformed at the temperatures of (a-c) 650 °C, (d-f) 700 °C, (g, h) 750 °C, and (i, j) 800 °C with the strain rates of (a, d, g, i) 0.001 s−1, (b, e) 0.1 s−1, and (c, f, h, j) 1 s−1.
different strain rates (0.001 s−1 and 0.1 s−1) are shown in Fig. 11. As it can be seen in this figure, a sharp stress drop occurs between the deformation temperatures of 700 °C and 750 °C. Considering the straininduced transformation temperature, i.e., between 700 °C and 750 °C, this fact implies that the transformed phases are softer than the transforming phase. The appearance of the inflection point in the work hardening rate versus stress plot (θ-σ) can be considered as an applicable method to find whether the phase transformation occurrence proceeds dynamically or not during deformation [33,34]. In the other words, considering the fact that the transformed phases are softer than the transforming phase and transformation mechanism involves nucleation and growth, the curve of work-hardening as a function of stress must have a zero point and an inflection point, as follow.
nucleation sites as well as elevating the driving force for phase transformation [28,32]. They promote in turn the nucleation phenomenon and increase the nucleation rate. Based on the fact that the diffusion rate is influenced by dislocation density, nuclei will grow rapidly and this will make the new grains of different phases to impinge on each other. By increasing the deformation strain rate higher density of defects, i.e., dislocation density and shear bands, can be formed. Therefore, the higher nucleation rate is accomplished with the higher rate of growth and will result in finer transformed phases. This fact can also be deduced from comparing the grain size of the samples deformed at the temperature of 800 °C with the strain rates of 0.1 s−1 and 0.001 s−1 (Fig. 7). As it can be seen from this figure, the deformed sample at higher strain rate is contained of both finer γ2- and ε-phases. Variations of the normalized change in stress at the strain of 0.7, i.e., ( T1 T2 )/ T2 , with respect to the deformation temperature at two
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Fig. 10. Backscattered electron SEM micrographs of the samples deformed at (a) 700 °C, (b) 750 °C, and (C) 800 °C with the strain rate of 0.001 s−1 associated with line scan EDS result for Al (gray) and Mn (red).
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power per unit volume (P) absorbed by the workpiece during plastic flow can be divided into two different parts, the power dissipated by the plastic work which mostly converts into heat (G), and the power dissipated due to the occurrence of the metallurgical mechanisms such as DRX, DRV and phase transformation (J) [35]. In order to find the power dissipation capacity of the material, a dimensionless parameter which is called the efficiency of power dissipation (η) has been introduced, i.e., = 2m / m + 1. It was stated that the deformation processing should be focused on the regions of maximum efficiency of the power dissipation, unless structural instabilities would occur, for example, flow localization. It has been stated that the material can be subjected to the instabilities during flow behavior when the rate of entropy production by the microstructural evolution in the material is lower than the applied rate of entropy. Based on this fact, the following criterion, ξ, was proposed to find the instability during deformation [36]:
=
Fig. 12. Variation of the work hardening rate with the stress.
= 0 And
=
=0 =p
(2)
The negative value of the instability coefficient indicates the flow instability and the location of the microstructural instability region can be found by mapping the instability parameter versus temperature and strain rate for a certain strain. The relationship between ln and ln is obtained by cubic curve fitting, and then the strain rate sensitivity exponent m can be obtained based on m ( ) = ln / ln T, equation for each certain temperature and strain as the function of strain rate. Fig. 13 represents the power dissipation contour map obtained through plotting the parameter η against the temperature and strain rate. As it was mentioned, the maximum η represents the optimal processing window. The higher value of η means that more power is dissipated due to the occurrence of dynamic microstructural evolution. As it can be seen from Fig. 13, η parameter is increased by increasing the temperature which is maximum for temperatures higher than 750 °C with strain rates of lower than 0.01 s−1. As it is shown on the map, this area (represented in dark blue) is the deformation conditions that lead to the dynamic phase transformation. In other words, the power was dissipated mainly through the phase transformation. The instability map obtained by plotting the instability coefficient against the temperature and strain rate is super-imposed on the more power dissipation map. The negative values of the instability coefficient and red area indicate the flow instability region which is located at the lower temperature region and the higher strain rates. As it was mentioned previously, the flow localization can occur in this area as the result of DRX around the grains. Based on the XRD results and the microstructural investigations, the blue areas reflecting the dynamic phase transformation condition are super-imposed on the instability map in order to find the optimal area for deformation. According to the microstructures presented in Figs. 10 and 13, compressive deformation at temperatures over than 700 °C leads to starting the τ to γ2 + β and subsequently to ε phase transformation. Considering this map, the optimal working area, in which the instability and the phase transformation are avoided is a narrow window at the temperature range between 650 °C and 700 °C and the strain rate lower than 0.1 s−1.
Fig. 11. Variation of the normalized change in saturation stress for change in the deformation temperature.
= c
ln(m /(m + 1)) +m ln
(1)
Fig. 12 shows the variations of work hardening rate regarding the stress for different strain rates. Based on Fig. 12, it can be found that for each curve, all the inflection points lay on a single straight line. It indicates that the rate of microstructure evolution is constant at the critical state. According to the presented evidence, it seems that the dynamic softening observed at temperatures of 750 °C and 800 °C is attributed to the dynamic phase transformation mechanisms. This phenomenon can be associated with adiabatic heating effect for high strain rates which accelerates the rate of softening during deformation. It has been shown that, during high temperature deformation, the
4. Conclusions The Mn-Al alloy with the chemical composition of Mn51Al47C2 was subjected to the compressive deformation at elevated temperatures with the wide range of strain rates and the following results were obtained:
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Fig. 13. Processing map for the strain of 0.7 with the corresponded optical microstructures, noting the red and the blue areas are corresponded to the instability zone and the phase transformation, respectively.
1) It was observed that during deformation, τ-phase dynamically transformed to γ2+ β + ε at a temperature lower than 750 °C. 2) Deformation at temperatures lower than 700 °C, where τ-phase was stable, led to the dynamic recrystallization around the initial grains which in turn resulted in the flow curve softening. 3) Based on the lower strength of the produced phases, the dynamic strain-induced transformation caused a flow softening in the stressstrain curves. 4) As a main finding after elaboration of different temperatures and strain rates during hot compression testing, the flow localization phenomenon and dynamic phase transformation can be prohibited in the ranges of 650–700 °C for temperature and 0.001 s−1 to 0.1 s−1 for strain rate.
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