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Qingshan Yang Subscribe SciFeed Recommended Articles Related Info Link Google Scholar More by Authors Links on DOAJ Molodov, K. D. Al-Samman, T. Molodov, D. A. on Google Scholar Molodov, K. D. Al-Samman, T. Molodov, D. A. on PubMed Molodov, K. D. Al-Samman, T. Molodov, D. A. /ajax/scifeed/subscribe Article Views Citations - Table of Contents Altmetric share Share announcement Help format_quote Cite question_answer Discuss in SciProfiles thumb_up ... Endorse textsms ... Comment Need Help? Support

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Get Information clear JSmol Viewer clear first_page Download PDF settings Order Article Reprints Font Type: Arial Georgia Verdana Font Size: Aa Aa Aa Line Spacing:    Column Width:    Background: Open AccessArticle On the Plasticity and Deformation Mechanisms in Magnesium Crystals by Konstantin D. Molodov

Konstantin D. Molodov SciProfiles Scilit Preprints.org Google Scholar 1, Talal Al-Samman

Talal Al-Samman SciProfiles Scilit Preprints.org Google Scholar 2 ORCID

and Dmitri A. Molodov

Dmitri A. Molodov SciProfiles Scilit Preprints.org Google Scholar 2,3,* 1 Department of Materials and Process Development, Salzgitter Mannesmann Forschung GmbH, Eisenhüttenstr. 99, 38239 Salzgitter, Germany 2 Institute for Physical Metallurgy and Materials Physics, RWTH Aachen University, 52056 Aachen, Germany 3 International Research Organization for Advanced Science and Technology (IROAST), Kumamoto University, Kumamoto 860-8555, Japan * Author to whom correspondence should be addressed. Metals 2023, 13(4), 640; https://doi.org/10.3390/met13040640 Submission received: 12 February 2023 / Revised: 9 March 2023 / Accepted: 21 March 2023 / Published: 23 March 2023 (This article belongs to the Special Issue Microstructure Evolution and Mechanical Properties of Magnesium Alloys) Download keyboard_arrow_down Download PDF Download PDF with Cover Download XML Download Epub Browse Figures Versions Notes

Abstract: This work presents an overview of the mechanical response and microstructure evolution of specifically oriented pure magnesium single crystals under plane strain compression at room temperature. Crystals of ‘hard’ orientations compressed along the c-axis exhibited limited room temperature ductility, although pyramidal 〈c + a〉 slip was readily activated, fracturing along crystallographic 11 2 ¯ 4 planes as a result of highly localized shear. Profuse 10 1 ¯ 2 extension twinning was the primary mode of incipient deformation in the case of orientations favorably aligned for c-axis extension. In both cases of compression along 〈 11 2 ¯ 0 〉 and 〈 10 1 ¯ 0 〉 directions, 10 1 ¯ 2 extension twins completely converted the starting orientations into twin orientations; the subsequent deformation behavior of the differently oriented crystals, however, was remarkably different. The formation of 10 1 ¯ 2 extension twins could not be prevented by the channel-die constraints when c-axis extension was confined. The presence of high angle grain boundaries and, in particular, 10 1 ¯ 2 twin boundaries was found to be a prerequisite for the activation of 10 1 ¯ 1 contraction twinning by providing nucleation sites for the latter. Prismatic slip was not found to operate at room temperature in the case of starting orientations most favorably aligned for prismatic slip; instead, cooperative 10 1 ¯ 2 extension and 10 1 ¯ 1 contraction twinning was activated. A two-stage work hardening behavior was observed in ‘soft’ Mg crystals aligned for single or coplanar basal slip. The higher work hardening in the second stage was attributed to changes in the microstructure rather than the interaction of primary dislocations with forest dislocations. Keywords: magnesium; single crystal; deformation twinning; plasticity 1. IntroductionMagnesium is the lightest structural metal, about one-third lighter than aluminum, and it has less than one-quarter of the density of ferrous alloys. Moreover, magnesium has a high specific strength and stiffness as well as high damping properties compared to other metals. It is abundantly available in the earth’s crust and seawater and has a high recyclability. Magnesium alloys are also superior at electromagnetic interference shielding, which makes them well suited for housings of portable electronic consumer goods. However, while cast magnesium is widely utilized, there are a number of technical difficulties hampering the large-scale use of wrought magnesium alloys.Generally, magnesium sheets have a poor formability at room temperature, which is improved markedly above 175 °C [1]. During the primary fabrication process, e.g., rolling, grains rotate to unfavorable orientations (strong texture) for isotropic plasticity [2]. Further processing steps must therefore be carried out at elevated temperature to thermally activate additional modes of deformation, which leads to high production costs. Further issues that limit the applicability of wrought magnesium are a strong yield asymmetry [3] due to deformation twinning, edge-cracking in rolling and a low ductility at high strength [4] as well as poor corrosion resistance [2].No matter whether the difficulties associated with the poor ductility of wrought magnesium at ambient temperature will be overcome by innovative processing, alloy and/or microstructure design, understanding the underlying physical mechanisms and microstructural changes in the material during processing is paramount. A lot of research effort was dedicated to uncovering the operating modes of deformation in magnesium. Well-known studies on magnesium single crystals were performed by Wonziewicz and Backofen [5] as well as by Kelly and Hosford [6] in the 1960s. Both studies included plane strain compression tests on differently oriented magnesium single crystals, demonstrating their complex deformation behavior, which is characterized by profuse deformation twinning and the occurrence of recrystallization. However, the investigations mentioned above were limited to rather small strains as their focus was on the temperature dependence of slip and twinning activation. Following these classical works, there was surprisingly little incentive to characterize the deformation behavior of Mg single crystals, despite the apparent need for reliable modeling based on real physical mechanisms. Hence, the lack of a holistic understanding of the deformation behavior, microstructure and texture evolution in Mg single crystals calls for a reexamination of the fundamental underlying mechanisms, especially considering the advent of more advanced characterization techniques and general progress in research.Compared to investigations of polycrystals, studies on single crystals claim several advantages. Single crystals can be precisely oriented with respect to the loading axes, which allows one to target and isolate specific modes of deformation, especially in conjunction with imposed constraints. Furthermore, model case experiments on single crystals provide an opportunity to examine the deformation behavior, microstructure and texture evolution depending on the initial orientation of the crystal. The operating slip modes and the critical resolved shear stress (CRSS) can be readily determined. Multiple generations of deformation twins can be easily distinguished even when the typical twin morphology is lost, which is rarely possible in polycrystals. The emergence of new grains and their orientation relationship to the surrounding matrix can be readily traced, which can grant valuable insight in the mechanisms of recrystallization [7]. In general, the use of single crystals permits a much clearer and forthright analysis of the observed microstructural changes compared to conventional studies on polycrystals.In light of these advantages of studies employing single crystals, it is imperative to note that, in general, individual grains in a polycrystal do not deform as single crystals. This is especially true for metals with crystal structures of high symmetry where multiple slip is prevalent. Single crystal studies therefore aim at exposing the underlying physical mechanisms, rather than claiming to directly predict the deformation behavior of polycrystals. That said, in the case of magnesium, very strong basal textures are predominant; hence, individual grains often undergo a similar deformation compared to what is tested in single crystal studies, especially when constraints are used. Furthermore, the small number of independent deformation modes, i.e., the dominance of basal slip and extension twinning at room temperature, entails a deformation behavior of the polycrystal that can be approximated by a Sachs-type model [8,9,10], in which the operation of only one deformation mode per grain is assumed. To this end, understanding the deformation behavior of single crystals is indispensable. Deformation of Mg at room temperature is particularly arduous owing to only a limited number of slip and twinning systems that can be activated to accommodate the imposed deformation. Plastic anisotropy is greatly pronounced, i.e., crystals of varying orientations are expected to exhibit a drastically different mechanical response (see, for example [11,12]). Performing model case experiments on specifically oriented single crystals allows to isolate and to identify the mechanisms of deformation, which in turn yields valuable insight in the deformation behavior of polycrystalline Mg.The aim of the present work is to provide an overview of the mechanical response and microstructure evolution of specially oriented pure Mg single crystals during plane strain compression at room temperature in order to expose the mechanism involved. Special emphasis is placed on the investigation of crystal orientations in which basal slip is inhibited initially since that provides an opportunity to focus on non-basal slip systems and deformation twinning more exclusively. Furthermore, the critical resolved shear stress of slip and twinning, the work hardening and the fracture behavior are characterized. 2. Specimens and Applied Methods to Study Deformation of Mg Single Crystals 2.1. Single Crystal GrowthMg of commercial purity (99.98%) was used as the starting material for pure Mg crystals. Cylindrical crystal seeds ( ∅ 5 mm × 25 mm ) with the c-axis aligned parallel to the axis of symmetry (maximum angular deviation of 0.2°) were cut from a randomly oriented single crystal by electrical discharge machining (EDM) [13] in order to avoid any plastic deformation of the material. EDM was performed using a wire with a diameter of 0.25 mm. The crystal seeds were etched in a 5% solution of nitric acid and boiled in distilled water for several minutes in order to form a stable oxide layer on the surface. The same procedure was done for a polycrystalline blank. The oxide layer acted as a barrier between the molten Mg and the inside of the mold, preventing diffusion and adherence. After preparation, the crystal seed and the blank were placed in a cylindrical steel mold. While the top part of the mold was made from stainless steel, the steel used for the bottom part that was in contact with the molten Mg did not contain any Ni to prevent contamination. A graphite-based die coating (Acheson Hydrokollag IP 5) was used to cover the inside of the mold.For the growth of oriented single crystals by direct solidification the Bridgman method [14] in a vertical configuration was utilized. The resulting crystals had a conical shape with a base diameter of 34 mm, a length of 56 mm and an opening angle of 4° (Figure 1a). A brass goniometer mounting was adhered to the back end of each crystal using glue mixed with copper powder to ensure conductivity during EDM (when the crystal was mounted onto a goniometer). Each grown single crystal was attached to a 3-axis goniometer (Figure 1b) and aligned according to the desired final orientation of the specimens using the Laue X-ray back-diffraction method detailed in Ref. [15] (Figure 1c). After alignment, specimens were cut for mechanical testing. 2.2. Channel-Die Plane Strain Compression TestsMg single crystals were subjected to plane strain compression (PSC) testing using a channel-die as shown in Figure 2a. For the PSC tests, cuboid specimens with dimensions of 14 mm × 10 mm × 6 mm were fabricated by means of EDM from the specially oriented single crystals. A schematic illustration of the channel-die device with respect to the specimen axes is shown in Figure 2b. Extension was limited to the longitudinal direction (LD) since extension in the transverse direction (TD) was suppressed by the channel-die walls. Mechanical tests were carried out at room temperature and a constant strain rate of 10−3 s−1 by means of a conventional screw-driven ZWICK 1484 testing machine. The specimens were strained up to various logarithmic (true) strains, defined by ε t = ln 1 + ε , where ε is the engineering strain. The applied force and displacement in the CD were both monitored and automatically controlled by a computer with a data acquisition system. The preloading force was 50 N, which translates into 0.36 MPa. To reduce friction, hydraulic oil was used for lubrication. Single crystals with different ideal orientations, including compression along (orientations A and B) and perpendicular to the c-axis (orientations C–F) as well as orientations with the c-axis inclined at an angle of 45° to the compression direction (orientations G and H), were tested in plane strain compression. The mismatch between the crystallographic directions and the specimen axes, i.e., compression (CD), longitudinal (LD) and transverse direction (TD) of the channel die, was less than 1°. An overview of the ideal specimen orientations is given in Table 1 expressed as ( φ 1 , Φ , φ 2 ) in Euler space. 2.3. Sample PreparationMetallographic sample preparation included soft grinding on grinding discs with 1200, 2400 and 4000 grit SiC-paper. Ethanol was used for lubrication/cooling in the final grinding step instead of water to avoid oxidation. Grinding was followed by 3 and 1 µm diamond polishing on Struers MD-Nap cloths at a very low rotation speed using an oil-based and water-free coolant until scratches could not be recognized visually. After diamond polishing, electropolishing in a 5:3 solution of ethanol and H3PO4 at 2 V was performed. Average electropolishing times were about 1 h. In some cases, e.g., channel-die specimens of orientation F, electropolishing times of up to 4 h were required to yield satisfying results, i.e., a deformation and twin free surface. Comparatively long electropolishing times were necessary since single crystalline Mg samples were particularly prone to deformation twinning during sample preparation (i.e., mechanical grinding and polishing). Electropolishing was performed at all times, irrespective of the characterization technique used, e.g., optical microscopy or electron backscatter diffraction. Extremely careful handling of specimens was crucial to avoid deformation of the soft single crystals. The deformed specimens after PSC tests were etched in a 5% solution of nitric acid, cut at the mid-surface of the LD-TD plane by EDM and prepared for further characterization. Chemical color-etching with a freshly prepared 1:1:7 solution of water, acetic acid and picral (4% picric acid in solution with ethanol) was performed for optical microscopy investigations using polarized light. 2.4. Characterization TechniquesMicrostructure characterization using polarized light was performed on a Zeiss Axio Imager A1m optical microscope. Furthermore, C-DIC (circular polarized light-differential interference contrast) [16] was applied to characterize the microstructure and surface topology. Micrographs (between 50 and 300 images per sample) covering the whole specimen surface were acquired manually at low magnifications and stitched together using a specially developed in-house software (utilizing the scale-invariant feature transform algorithm by David Lowe [17]) to obtain a complete high-resolution macroscopic image of the investigated plane. Supplementary stereomicroscopy was performed on a Zeiss Discovery V12 optical microscope fitted with variable LED illumination to provide orientation-dependent information on the surface topology at low magnifications, albeit without offering color contrast. The Laue (back-reflection) technique was used to measure the orientations of single crystals (see Section 3.1) by placing a film between the X-ray source (W-tube operating at 45 kV and 30 mA) and the measured specimen. The positions of Laue-reflections (spots) on the film were determined using magnified scans of the Laue micrographs and indexed using the software ‘OrientExpress 3.4′ by J. Laugier (Laboratoire des Materiaux et du Génie Physique de l’Ecole Supérieure de Physique de Grenoble).Electron backscatter diffraction (EBSD) measurements were performed using a scanning electron microscope (SEM) equipped with a field emission gun (Leo Gemini 1530 with a LaB6 filament) and an HKL-Nordlys II EBSD detector to characterize the microstructure and microtexture. The samples were mounted onto custom machined brass holders to ensure a perfect alignment and tilted to 70° (from the horizontal) towards the EBSD detector to increase the backscattering yield. An acceleration voltage of 20 kV was used for all measurements. The Matlab toolbox MTEX [18,19] was utilized to analyze and visualize the EBSD data, including correction of the EBSD raw data (filtering based on a maximum mean angular deviation and nearest neighbor extrapolation) and grain reconstruction. 3. Plastic Response and Microstructure of Differently Oriented Mg Single Crystals 3.1. Flow Behavior of Mg Single CrystalsAs shown in Figure 3, the investigated pure Mg single crystal specimens with different initial orientations exhibited very different mechanical behavior during ambient temperature PSC deformation. The four orientations A, C, E and G (Figure 3a) shared a 30° rotation around their respective c-axes with the orientations B, D, F and H (Figure 3b). For instance, in the case of orientations A and B, the imposed deformation yielded contraction along the c-axis while extension was confined to the crystallographic 〈 11 2 ¯ 0 〉 and 〈 10 1 ¯ 0 〉 directions, respectively. Orientations C and D represented the case of c-axis extension, whereas the c-axis extension was confined for the orientations E and F (c-axis parallel to TD). Orientations G and H were the only ones that were aligned for basal slip. For all other orientations, the basal plane was either parallel or perpendicular to the compression direction, i.e., basal slip was suppressed.Specimens of orientations A (Figure 3a) and B (Figure 3b), i.e., the case of c-axis contraction, displayed the least ductility at room temperature, reaching a maximum true strain of −7.3% and −6.6%, respectively, at fracture. Crystals of other orientations were significantly more ductile, reaching a true strain of −1, with the exception of orientation D, the flow behavior of which was more similar to that of orientation B (Figure 3b) than its counterpart orientation C (Figure 3a).Apart from crystals of orientations C and D that showed a highly diverse deformation behavior in c-axis extension, the stress-strain curves for the other orientations were qualitatively similar when compared to their counterparts (rotated by 30°). However, quantitatively, significant differences were apparent. In other words, no single crystal tested displayed the exact same deformation behavior as another one with a different orientation, demonstrating the great extent of plastic anisotropy in Mg. 3.2. Contraction along the c-Axis (Orientations A and B)The flow curves for specimens of orientations A and B are shown in Figure 4. Basal slip was suppressed in both cases. Initial strain hardening rates were among the highest compared to other orientations. The hardening rates d σ t / d ε t for specimens of orientation A and B were determined to be 6.3 GPa and 6.9 GPa, respectively, in the initial linear region of the flow curves. However, the hardening rates decreased at higher strains, most notably seen in the case of orientation A (Figure 4). The hardening rates d σ t / d ε t measured in the linear regions of the flow curves prior to fracture were 3.2 GPa for orientation A and 5.5 GPa for orientation B. For both orientations A and B, fracture always occurred along 11 2 ¯ 4 planes repeatedly, as seen in Figure 5. The specimens typically cracked in a rather violent event that caused pieces to break off as shown in Figure 5b. The maximum true stress attained was 320 MPa and 390 MPa for orientation A and B, respectively. Right before fracture occurred, the flow curves exhibited an inflection point, followed by a continuous brief decrease in flow stress. Specimens of orientation A that were polished prior to testing and strained without exceeding the rupture stress revealed the existence of slip traces. Slip traces on the CD-TD plane (the only unconstrained plane in channel-die compression) are shown in Figure 6. Fine horizontal traces (Figure 6a) as well as fine diagonal lines at an angle of 54–57° to TD (Figure 6b) and coarser diagonal bands (Figure 6c) in a diamond-like pattern were present. These slip lines were found to be nearly parallel to traces of 11 2 ¯ 2 planes (Figure 6). It is worth noting that the horizontal lines depicted in Figure 6a could have also originated from slip on the basal plane; however, basal slip could not account for the appearance of diagonal lines on the CD-TD surface of A-orientated crystals (Figure 6b).The CD-LD surface of specimens with orientation B was significantly more rutted compared to those with orientation A. Like in the case of orientation A, horizontal slip lines were apparent. Non-rectilinear diagonal slip traces were observed as well at inclinations in a brought range of 43–58° to TD (Figure 7). These traces could not be clearly assigned to a particular slip plane; however, they were not a result of basal slip. Besides slip lines, numerous 10 1 ¯ 2 extension twins were frequently observed on the CD-LD surface of specimen A and B despite the imposed c-axis contraction. However, their appearance must be attributed to inhomogeneous flow during unloading and the corresponding residual stresses [5,20].Detailed examination of the failed samples (orientation B) revealed the presence of narrow recrystallized bands nearly parallel to the trace of 11 2 ¯ 4 planes (planes of fracture) on the TD-LD surface (Figure 8). The matrix orientations adjacent to the band in Figure 8 (left and right) were rotated away from the initial orientation of the single crystal in different directions about an axis parallel to the band, leading to a misorientation angle of up to 8° between them, i.e., local splitting of the initial basal texture occurred. At other bands (not shown here) splitting of up to 30° was observed. Cracks were either directly evident in such bands, mostly along grain boundaries (Figure 8a), or these bands continued into cracked regions (Figure 8b). An orientation map of the recrystallized band in Figure 8b is shown in Figure 8c. A larger number of 10 1 ¯ 2 extension twins were often found in the vicinity of such cracked/recrystallized bands (blue in Figure 8c); however, 10 1 ¯ 1 or 11 2 ¯ 4 contraction twins were not observed. Moreover, the grains in the band interior did not harbor any indication that the band was a former 11 2 ¯ 4 twin, although this is suggested by the 11 2 ¯ 4 trace of the band. None of the grains had a misorientation (with respect to the matrix) close to that of a 11 2 ¯ 4 twin, which would be 78.14° around 〈 1 ¯ 100 〉. 3.3. Extension along the c-Axis (Orientations C and D)As mentioned above (in Section 3.1), the flow behavior of C- and D-oriented crystals in c-axis extension was distinctively different. The respective stress–strain curves at room temperature are depicted in Figure 9. Specimens of orientation D failed at a true strain of −14.7%, whereas samples with the starting orientation C showed extraordinary high room temperature ductility and deformed up to a true strain of −100%. Analogous to the case of orientation A and B, basal slip was suppressed for both starting orientations C and D due to the alignment of the basal planes with respect to the compression direction. The incipient stage of deformation was characterized by a low strain hardening rate ( d σ t / d ε t ) of 544 MPa and 297 MPa for orientations C and D, respectively. However, at strains above −5%, the initial region of low work hardening was followed by a linear hardening region with significantly higher strain hardening rates of 4.6 GPa and 5.6 GPa for orientation C and D, respectively. The rapid hardening region was interrupted by a sudden continuous drop in flow stress in the case of orientation C at a true stress of 220 MPa and a true strain of −10.5%, after which the flow stress remained steady up to a true strain of −20%. No such drop in flow stress was witnessed for orientation D. Instead, the rapid work hardening continued until fracture. It is worth noting that the fracture stress and hardening rate in the rapid work hardening regime of orientation D was nearly identical to that of orientation B (c-axis contraction) (Figure 3).With respect to the microstructure evolution, profuse 10 1 ¯ 2 extension twinning was observed during early stages of deformation (Figure 10). 10 1 ¯ 2 extension twins consumed about half of the matrix at a true strain of −3% for both C- (Figure 10a) and D-oriented specimens (Figure 10b). In addition to those twins, secondary 10 1 ¯ 2 extension twins were also found inside the primary ones in C-oriented crystals (inset in Figure 10a).A major difference in the behavior of specimens with the initial orientations C and D at this point was in the selected twin variants. For orientation C, the TD–LD (primary) twin traces were at an angle of ±65° to LD (Figure 10a), while for orientation D, the traces of the twinning planes were parallel to TD, as evident from the twin trace in Figure 10b. The twin orientations that were obtained by EBSD measurements are illustrated in Figure 11. In the case of orientation C (Figure 11a), the basal poles of the 10 1 ¯ 2 extension twin orientations were about 30° away from the center of the basal pole figure (compression direction). By contrast, the twin orientations that originated from orientation D were located right in the center of the basal pole figure (Figure 11b).At slightly higher strains, the respective primary twins completely consumed the initial single crystal matrix of orientation C (Figure 12a) and D (Figure 12b), i.e., the starting orientations C and D were completely converted into twin orientations. At the peak stress on the flow curve for orientation C, secondary 10 1 ¯ 2 extension and, more importantly, 10 1 ¯ 1 contraction twins were evident. (For an in-depth description of the deformation behavior of C-oriented crystals at higher strains, the reader is referred to Ref. [21].) No 10 1 ¯ 1 contraction twins were seen to form in the case of orientation D. Specimens of orientation D effectively turned into single crystals after primary twinning (Figure 12b). The orientation of the new matrix (Figure 11b and Figure 13) was nearly identical to that of orientation B (c-axis contraction), i.e., the c-axis was parallel to CD and a crystallographic 〈 10 1 ¯ 0 〉 direction was parallel to LD. However, strictly speaking, the specimens were not fully single-crystalline but comprised low angle (

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