Twinning effects in deformed and annealed magnesium–neodymium alloys

Author’s Accepted Manuscript Twinning Effects in deformed and annealed magnesium-neodymium Alloys C. Drouven, I. Basu, T. Al-Samman, S. KorteKerzel www.elsevier.com

PII: DOI: Reference:

S0921-5093(15)30338-5 http://dx.doi.org/10.1016/j.msea.2015.08.090 MSA32724

To appear in: Materials Science & Engineering A Received date: 22 June 2015 Revised date: 27 August 2015 Accepted date: 28 August 2015 Cite this article as: C. Drouven, I. Basu, T. Al-Samman and S. Korte-Kerzel, Twinning Effects in deformed and annealed magnesium-neodymium Alloys, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2015.08.090 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.

Twinning Effects in Deformed and Annealed Magnesium-Neodymium Alloys C. Drouven, I. Basu*, T. Al-Samman, S. Korte-Kerzel Institut für Metallkunde und Metallphysik, RWTH Aachen University, 52056 Aachen, Germany *Corresponding author. Tel.: +49 241 80 26892; Fax: +49 241 80 22871; Email: [email protected]

Abstract Hot rolled Mg-1wt.% Nd alloy with a weak {0001} < 1120 > basal texture was subjected to plane strain compression (PSC) at room temperature. Compressive loads were imposed parallel to the in-plane (IPC) and through-thickness (TTC) directions of the rolled plate in order to engage different types of deformation twins. PSC was conducted up to 6% and 12%, and the deformed specimens were subsequently annealed at different elevated temperatures to investigate recovery and recrystallization in the formerly developed twins. Deformation and annealing textures and microstructures were characterized using X-ray diffraction, optical microscopy and electron backscatter diffraction techniques. Deformation under IPC and TTC displayed competitive activation of extension and contraction twinning, supplemented by homogenous dislocation slip. Deformation in IPC up to 6% witnessed equal activation of both aforementioned twinning modes that gave rise to soft and hard texture components. TTC up to 12% exhibited a less regular basal-type deformation texture, characterized by a large outspread in the rolling direction, and an asymmetrical splitting of the basal poles around the transverse direction by about -5° and +25°. Texture evolution during annealing was associated with general texture weakening, preserving the primary deformation texture components at relatively low annealing temperatures with no recrystallization (300°C). Presence of strong impurity drag and particle pinning delayed the onset of recrystallization till 325°C for TTC and 350°C for IPC specimens, which took place preferably inside compression and double twins. Related recrystallized grains possessed a high Schmid factor (SF) for basal slip, and were larger and more prevalent on average than competing low SF grains related to extension twins or the matrix. Re-dissolution of precipitates at 450 °C triggered grain growth that led to strong evolution of orientations associated with compression and double twins as a result of a grain size advantage that emerges during recrystallization. Keywords: Deformation twinning, texture, dynamic recovery, particle pinning, rare-earth elements

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1 Introduction Understanding rare-earth (RE) related texture modification in magnesium alloys is a topic, which despite extensive studies still lacks considerable clarity. From our previous investigations on rolled and annealed binary magnesium alloys with single RE additions, it was observed that the extent of texture modification is strongly dependent upon i) the processing parameters and ii) the choice of RE-addition [1-3]. The modified textures were either characterized by a decrease in texture intensity or by evolution of new orientations viz. an ‘RE-texture’, unlike conventional Mg sheet textures [3-6]. While the processing conditions play a fundamental role in facilitating those recrystallization (RX) and grain growth mechanisms that are more likely to render texture modification, such as RX nucleation in shear bands [2, 5, 6], in deformation twins [4, 7-9] or near particles [10-13]; it was seen that the choice of RE-addition strongly alters the conventional deformation and annealing behavior, on the basis of its interaction with the matrix i.e. as solutes or second phases [2, 14]. Hence, in order to effectively correlate the mechanisms underlying deformation, recrystallization and grain growth with the chemical composition, it becomes imperative to characterize the behavior of individual RE-elements as an alloying addition with respect to material processing. Twinning in magnesium and its alloys is quintessential due to a lack of slip modes that can accommodate strains along the c-axis. However, the complexity of deformation twinning arises from the associated large lattice rotations from small strain increments, which significantly influence the subsequent local strain compatibility and overall hardening behavior [15-19]. Furthermore, the rapid reorientation of the twinned matrix becomes instrumental in defining the deformation texture evolution. The most commonly observed twinning modes in magnesium alloys are primary {1012} < 1011 >tension twins involving a lattice rotation of ~86° about the < 1120 > axis and {1011} < 1012 >compression twins leading to a misorientation of ~56° about < 1120 > [15, 16]. Additionally, second generation {1011} − {1012} double twins, wherein the compression twinned regions undergo tension twinning have also been reported frequently [20]. Unfortunately, deformation twinning in magnesium has often been designated as one of the prime causes behind the material’s poor formability due to its detrimental effects on the subsequent deformation behavior. While nucleation and growth of tension twins is strongly favored in conventional magnesium alloys, owing to its lower critical resolved shear stress, the associated lattice reorientation often leads to a basal type sheet texture, which is difficult to deform further [21]. -2-

On the other hand, the growth of nucleated compression twins being energetically unfavorable in most cases results in large stress accumulation at the twin matrix interface, thereby causing material cleavage [15, 22-24]. With respect to twinning induced recrystallization, several studies indicate that nucleation inside compression and double twins can give rise to soft orientations, giving it a significant role in RX texture modification [7, 25, 26]. The process of recrystallization inside twins is dominated by local recovery, either slip assisted or thermally activated, leading to formation of sub-grains resembling twin orientations [12, 27, 28]. Further thermal activation can trigger growth of such orientations, due to the associated high driving pressures. Even though neodymium (Nd) is one of the more commonly used RE-addition in commercial magnesium alloys (e.g. as Nd-based mischmetal), its role in influencing deformation and recrystallization behavior in magnesium is little understood. First investigations on Nd-containing magnesium alloys were reported in the early 1980s, wherein wrought Mg-Nd alloys tested in tension, displayed an increase in the maximum elongation with increasing Nd content [29]. The tendency inverted once the Nd concentration exceeded 4.2 wt. %. This ductility enhancement at low Nd additions was associated with grain refinement. The authors further postulated that even low amounts of Nd can lead to formation of stable secondary phases with the matrix impurities, thereby purifying the grain boundaries and resulting in enhanced plasticity. In the past two decades, studies on twinning behavior in magnesium-neodymium alloys have been relatively few in number, with no conclusive findings. Investigations by Hantzsche and co-workers [4], showed that when the Nd content was increased from 0.01 at.% (~ 0.05 wt.%) to 0.04 at.% (~ 0.22 wt.%) the tendency of twinning inclined more towards nucleation of compression and double twins as compared to expected tension twins. The twins appeared in the form of deformation bands, producing orientations with off-basal character. Quite contrarily, a separate study on Mg-Mn alloys containing 0.5 wt. % and 1wt.% Nd reported anomalously high activity of tension twinning at low temperatures with increasing RE amounts and was attributed to an increased critical resolved shear stress for basal slip, arising from interaction between Mg3Nd intermetallic prismatic plates and basal dislocations [30]. In another study on Mg-Nd alloys [31], it was reported that the recrystallization related texture weakening was primarily due to Zener pinning. For low Nd content alloys, grain growth led to strengthening of the texture, whereas in high Nd content alloys, annealing at higher temperatures resulted in particle coarsening and abnormal grain growth. -3-

The current work focuses on deformation twinning in a rolled Mg-1wt. %Nd alloy and the related annealing response at different temperatures. It will be shown that twinning in this alloy is characterized by equitable activation of compression and tension twins, regardless of the starting orientation, and that annealing primarily results in texture weakening, with Nd playing a dual role of solute as well as precipitate. Moreover, the underlying mechanisms behind the deformation and recrystallization behavior will be discussed with respect to the chemical composition.

2 Experimental procedure Fig. 1 represents a schematic of the processing route implemented in the current work. Binary Mg-1wt. % Nd alloy was produced by means of induction melting under a protective gas atmosphere of Ar/CO2. The cast billets were subjected to solution annealing at 450°C for 20 h in order to homogenize the material. The measured average grain size in the homogenized as-cast state was ~767 μm ± 4.8% relative standard error (R.S.E). Rolling blocks of dimensions 80 mm x 40 mm x 40 mm, obtained from the as-cast billet, were hot rolled at 450°C (nominal furnace temperature) in multiple passes, applying 10% thickness reduction per pass. After each pass the rolled material was returned to the furnace and held at 450°C for 6 min. The final thickness reduction achieved was 70%, amounting to an overall true strain of ε = -1.2 and a final material thickness of 12 mm. After rolling, the material was subjected to a short recrystallization annealing treatment at 400°C for 12 min, in order to obtain a fully recrystallized and stress relieved microstructure, yet preserving the rolling texture for subsequent loading parallel and perpendicular to the sheet normal. Deformation experiments were performed under plane strain compression (PSC) using a channel die set-up in a conventional screw driven universal testing machine, ZWICK-1484. The deformation was imposed under a constant strain rate of 10-2 s-1 at ambient temperature. Friction effects were minimised by using hexagonal boron nitride as lubricant between specimen and the die walls. The PSC specimens were machined from the rolled material in two different orientations corresponding to (i) through thickness compression (TTC), wherein the compression axis was parallel to the c-axes of most grains and (ii) in-plane compression (IPC), with the compression axis being perpendicular to the c-axes in the material (cf. Fig. 1d). The specimen dimensions used for the channel-die tests were 12 mm along the longitudinal direction (LD), 10 mm along the transverse direction (TD) and 4 mm along the compression direction (CD). While {1012} tension twinning is typically known to exhaust after an imposed strain of ~6.5%, the nucleation of {1011} compression twins have been -4-

observed close to the peak flow stress values [15, 16]. On the basis of this information, deformation was conducted up to a selected strain of 6% for the IPC mode and 12% (corresponding to the recorded peak stress) for the TTC mode. In order to ensure reproducibility, each test was performed 3 times and the average stress-strain curves were presented. Post PSC deformation, 60 min isochronal annealing treatments were conducted at temperatures between 300°C and 450°C. The annealed specimens were subsequently water quenched. X-ray diffraction (XRD), optical microscopy and electron backscatter diffraction (EBSD) measurements were performed on the deformed and annealed specimens. Specimens for XRD measurements were ground with SiC paper (1200 grit) to the sample mid-plane to obtain the bulk texture and avoid surface effects. Subsequently they were polished with 3 and 1 µm alcohol based diamond paste, and mirror-finished with 0.25 μm colloidal silica. X-ray pole figure measurements were performed using Bruker D8-Advance diffractometer operating at 30 kV and 25 mA and equipped with a high resolution area detector. The MTEX MATLAB toolbox [32] was employed to calculate the orientation distribution functions (ODF) and full pole figures from the collected raw diffraction data using spherical harmonic functions. For microstructure characterization, mechanically polished samples were additionally electropolished using Struers AC2 solution. The electro-polishing parameters are given in Table 1. For polarized light microscopy, respective specimens were etched in acetic picral to reveal the grains and grain boundaries. Automated EBSD measurements were conducted on electropolished samples using a LEO-1530 scanning electron microscope equipped with a field emission gun operating at 20 kV and an HKL-Nordlys II EBSD detector with a collection step size of about 0.5 µm. The resulting EBSD indexing of the current specimens was above 80%. The acquired raw EBSD data was analyzed using the HKL Channel 5 and MTEX software packages. Chemical characterization of secondary phases was performed on a FEI Helios 600i dual-beam microscope equipped with EDAX energy-dispersive X-ray spectroscopy (EDX) system. Grain size measurements were performed by means of conventional lineal intercept techniques. Table 1: Electro-polishing parameters of Mg-1Nd specimens

Temperature Voltage Polishing time

-20°C to -25°C 34 V 35 s

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3 Results 3.1 Characterisation of the material prior to PSC deformation The applied thermo-mechanical treatment prior to the deformation experiments led to the emergence of a relatively weak texture, shown in Fig. 2b and c for the TTC and IPC loading orientations, respectively. Even though the short duration anneals successfully retained the basal characteristic of the rolling texture, strong texture weakening was on the other hand inevitable, with the maximum basal pole intensity decreasing to ~ 3 MRD (Multiples of random distribution). It is noted that the effect of different loading schemes in TTC and IPC would be less dramatic in case of a weak sheet basal texture but in the present case we were still able to parse relevant conclusions from the observations obtained in both loading modes. The observed texture weakening during rolling and annealing seems to be a distinctive feature of rolled Mg-REs subjected to recrystallization treatments, since in conventional Mg alloys the basal rolling texture is normally retained after recovery and RX annealing [33], which renders Mg alloys similar to body centred cubic (bcc) alloys in this regard. During grain growth annealing conventional Mg alloys can even experience strengthening of the ‘already strong’ deformation texture [34]. The optical microstructure upon the short RX treatment displayed a fully recrystallized structure with equiaxed grain morphology (Fig. 2d). The measured average grain size was 11.6 ± 1.5 μm. 3.2 PSC deformation – Flow behavior and microstructural evolution Fig. 2a displays the plane strain compression flow curves for loading in IPC and TTC at ambient temperature. The observed initial non-linearity in the flow curves is possibly a result of compliance effects between compression plates and the specimen. Deformation in the IPC mode revealed a slightly higher hardening rate at lower strains (< 5%). For strains higher than 5%, the hardening rate showed a sudden decrease, possibly indicating the onset of deformation twinning. The peak stress (corresponding to a logarithmic true strain of 12%) for the TTC mode, was measured as ~ 380 MPa. In both loading types, the material did not show any signs of failure after the selected strains were reached. The TTC bulk texture revealed a prominent basal pole splitting with a maximum intensity of 7 MRD (c.f. Fig. 3a). The related TTC microstructure featured heavily-banded structures, possibly arising from sequential twinning events. The twins seemed to assume a fine and needle-like morphology, visible only at higher magnifications (500X). Additionally, a few recrystallized regions comprising fine equiaxed grains were also identified. For deformation in IPC, the primary texture components were concentrated in the vicinity of LD (cf. Fig. 3b), with a new, yet faint texture component -6-

emerging in the centre of the pole figure. The texture intensity in the deformed state was ~ 5 MRD; lower than that measured for the TTC mode. The deformation microstructure corresponding to the IPC mode, revealed characteristics quite unlike the expected behavior. Deformation in IPC mode would be expected to invoke c-axis extension, and in that respect, trigger {1012} extension twinning, which typically shows a thick wavy morphology [24]. However, the micrographs shown in Fig. 3b evince a microstructure consisting of fine lamellar twin clusters forming banded regions along with large wavy lenticular twins distributed homogeneously in the microstructure. Fine equiaxed grains in the vicinity of banded twins were additionally observed. In both cases, the as-deformed state revealed presence of extremely fine precipitate clusters (average size < 1 μm) distributed homogeneously throughout the microstructure. EDX analysis of these precipitates revealed a chemical composition of Mg12Nd, which is known to be a magnesium rich metastable phase [29]. 3.3 Isochronal annealing treatments – XRD textures and microstructural evolution Figs. 3a and b display the corresponding texture and microstructural evolution during static annealing treatments for the TTC and the IPC modes. In case of TTC, annealing at 300°C led to transformation of the split basal texture, seen in the as-deformed state, into more or less a single-peak basal texture, with the peak component being slightly shifted from the CD (cf. Fig. 3a). The texture intensity revealed negligible changes. Annealing at 350° C, retained the basal characteristic but resulted in a reduction of the texture intensity from ~7 MRD down to ~5 MRD. For annealing treatments at even higher temperatures (viz. 400°C and 450°C) the texture features and intensity remained more or less invariant (cf. Fig. 3a). The IPC deformed specimens (cf. Fig, 3b), when annealed at 300°C, revealed a spread of orientations between LD and CD associated with slight texture weakening relative to the deformed state. Annealing at 350°C and higher temperatures led to a complete eclipsing of the basal orientation. At 400°C, a well-defined texture component is observed to appear in between LD and CD, with the c-axes tilted ~45° (with a spread of ±15°) away from the LD (cf. Fig. 3b). At 450°C, annealing resulted in an increased spread of the orientations (i.e. randomisation) lying between LD and CD. With respect to the influence of annealing temperature upon the texture intensity, heat treatments above 300° C led to an interesting nonmonotonic behavior (cf. Fig. 3b), wherein the texture intensity showed an anomalous maximum at 400°C (~ 9MRD), followed by significant drop for annealing treatment at 450°C (~5MRD). -7-

Microstructurally, the onset of recrystallization in the TTC mode was not witnessed until 300°C, evident from the still present twinned structure (cf. Fig.3a). Annealing at 350°C and 400°C revealed a fully recrystallized grain structure with limited grain growth. Increasing the temperature to 450°C caused a transition from a relatively fine to a coarse grained microstructure, indicating a sudden acceleration in grain growth kinetics. With respect to IPC, recrystallization seemed to initiate at 350°C. A partially recrystallized microstructure, with nucleation of extremely fine grains inside the banded twin regions, is observed in Fig. 3b. Annealing at 400°C produced a bimodal microstructure comprising recrystallized bands interspersed with unrecrystallized twinned regions (cf. Fig 3b). Further increase in temperature marked the onset of rapid grain coarsening, as seen in the TTC experiments. The density of the fine precipitates, observed in the as-deformed microstructure, displayed a considerable decrease with increasing annealing temperature. 3.4 EBSD analysis In order to gain further insight into the recrystallization and grain growth related microstructural evolution, EBSD measurements were performed on select annealed states. Figs. 4a and b represent typical EBSD microstructures for the investigated Mg-1Nd alloy after TTC and annealing at 325°C, and after IPC and annealing at 350°C, respectively. The annealing time for both samples was 60 min. Image quality (IQ) indexing is used as the background here and the boundaries of the different twin types are highlighted using different colours. Since the TTC microstructure remained fully deformed after the 300°C anneal and was completely recrystallized after 350°C, an intermediate temperature was selected to obtain a partially recrystallized microstructure. The initial indexing rates were 77% and 50%, for the IPC and TTC specimens, respectively, which were further cleaned by standard noise reduction up to a minimum of 5 indexed neighbours. The remaining zero solutions are shown here in black. The twin boundaries were identified by means of their characteristic minimum axis/angle relationships (Table 2), allowing a maximum spread of ±6°. High angle grain boundaries (> 10°) have been highlighted in black. Table 2 Twin-matrix misorientation relationships for different twin types in Mg

Twin type Tension twin (TT) Compression twin (CT) Double twin (DT)

Twinning plane {1012} {1011} {1011} − {1012}

Misorientation axis 〈1120〉 〈1120〉 〈1120〉

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Misorientation angle 86.3° 56.2° 37.5°

The deformed microstructural regions in the TTC specimen, shown in gray (cf. Fig. 4a), are characterized by long elongated grains along with numerous thin twin clusters (indicated by the systematically non-indexed regions). The identified twins in different colours revealed {1012} extension twins (highlighted in red), {1011} compression twins (blue), as well as {1011} − {1012} double twins (yellow). The EBSD microstructure of the IPC specimen (Fig. 4b) revealed a peculiar feature of much higher density of compression and double twin boundaries than expected for this loading mode. The twins appeared, similar to Fig. 4a, in sequential microstructural pattern, giving rise to banded regions in the microstructure. The recrystallized grains, highlighted in green (cf. Figs. 4a and b), were demarcated from the still present deformation microstructure on the basis of a critical grain orientation spread (GOS) of 1° [2]. As evident from the EBSD maps, most of the recrystallized regions were closely correlated with the twin clusters. The average recrystallized grain size measured from Fig. 4 for both specimens was ~ 2 µm. Fig. 4c and d show the measured EBSD textures in terms of single orientation scatter and recalculated basal pole figures of the recrystallized and deformed fractions of the annealed TTC microstructure at 325° C. While the deformation texture showed strong clustering of basal orientations, and other mild off-basal components, the recrystallization texture comprised a large spread of basal and off-basal orientations. The texture intensity of the recrystallized fraction, displayed in the recalculated pole figures was much weaker than the intensity of the deformation texture, marked by an intense basal orientation. Analogously, Fig. 4e and f depict the EBSD textures of the IPC microstructure after annealing at 350°C. The pole figures associated with the recrystallization microstructure displayed a continuous span of orientations between LD and CD forming a recrystallization fiber. On the other hand, the pole figures associated with the deformed structure showed a concentration of orientations in the vicinity of LD, matching the deformation texture prior to annealing. Noteworthy is the complete absence of basal orientations in both recrystallized and deformed grains despite loading in IPC. In both cases (TTC and IPC) the overall EBSD texture, which is a superposition of the deformed and recrystallized data subsets, corroborated well with the XRD bulk texture given in Fig. 3. Fig. 5 shows the EBSD microstructures and textures for the IPC and TTC specimens after annealing at 450°C. The original indexing rate for both measurements was above 90%. Figs. 5a and d represent the inverse pole figure (IPF) mappings for IPC mode and the TTC mode respectively. The illustrated microstructures were obviously completely recrystallized and underwent subsequent grain growth. The measured average grain size was about 54 ± 0.8 μm for both specimens. The corresponding EBSD textures are shown in Figs. 5b and e for the IPC -9-

and TTC specimens, respectively. Multiple mappings acquired at similar magnifications were merged in order to obtain statistically sufficient local texture data. The resultant textures hence displayed are superposition of EBSD datasets, accounting for 1085 and 1220 grains for the measured IPC and TTC specimens, respectively. The pole figures corresponding to the annealed IPC sample, showed a concentration of orientations between LD and CD, with the peak pole density residing approximately at 35° away from LD towards CD. No basal component was visible in the texture data. In case of the annealed TTC sample, the pole figure data indicated a large spread of orientations about the CD with comparable texture strength relative to the IPC sample. The EBSD pole figures showed good agreement with the XRD macro-texture data. Fig. 5c and f are a reproduction of the same IPC and TTC microstructures in Fig. 5a and d represented in terms of Schmid factor (SF) maps for basal slip. In this representation, a ‘red’ grain with an SF value close to 0.5 experiences the highest resolved shear stress on its (0001) basal plane and along its < 1120 > slip direction, making it a soft grain in terms of ease of slip. Analogously, ‘blue’ grains with very low SF values are considered ‘hard’ grains in this context. The relatively high average SF values (above 0.3) after both IPC and TTC annealing indicate a dominance of soft orientations, described by a high density of non-basal orientations (given by a large texture spread) in TTC and a welldefined texture component between LD and CD in case of IPC. 4 Discussion The experiments conducted in the present work were designed to obtain a clear correlation between extension/compression twinning propensity and the loading of the sheet (in-plane vs. through thickness) in the presence of a rare earth element (Nd) that is known to produce both solute and precipitate effects. The obtained results indicated that the observed twinning behavior and the texture evolution during annealing were quite different from what is so far known from conventional magnesium alloys subjected to similar treatments. The forthcoming sections hence discuss in greater detail the underlying mechanisms triggering such distinctive deformation (cf. Sections 4.1 and 4.2) and annealing (cf. Section 4.3) characteristics in the investigated Mg-1Nd alloy. 4.1 Twinning behavior during TTC and IPC loading The starting texture prior to loading in the TTC mode of deformation (cf. Fig. 2b) displayed a weak basal texture with significant angular spread of basal poles about the CD in both LD and TD directions. In pure Mg and also conventional Mg alloys processed by rolling, significant twin formation is only seen when the sheets are subjected to IPC, in which - 10 -

compression is applied in the rolling or transverse directions of the sheet, i.e. normal to the crystallographic c-axis. This type of loading will preferentially nucleate {1012} extension twins that are typically easy to form due to their very low activation stress unlike {1011} and {1013} compression twins [16, 24]. Interestingly, compression of the current Mg-1Nd alloy sheet in the through thickness direction resulted in a high density of {1011} compression twins and {1011} − {1012} double twins, with negligible traces of {1012} extension twinning. In this respect, Fig. 6a shows an IQ mapping for the as-deformed state in TTC mode, with extension, compression and double twin boundaries highlighted in red, blue and yellow, respectively. Fig. 6b represents the same region in IPF colouring, highlighting low (> 2°) and high (> 10°) angle grain boundaries in white and black, respectively. Owing to the weak, initial texture with a broad orientation spectrum, deformation twinning produced various grain orientations, some of which were almost basal and others off-basal. The magnified region shown in Fig. 6b displays a parent grain ‘P1’ with off-basal double twin (daughter) orientation ‘I’, and another parent grain ‘P2’ with basal double twin orientation ‘II’ (cf. Fig. 6d). Schmid factor EBSD map for basal slip assuming compression along the CD (through thickness) is shown in Fig. 6c. It is noted that an additional, lateral stress component originating from the channel die walls (back stress): $%& = 0.5$'& was considered here for the SF calculations. A high average SF value in most of the deformed regions indicates strong favorability toward basal slip. The scatter and recalculated pole figures for the as-deformed condition are displayed in Fig. 6d, indicating a mild deformation texture with both basal and off-basal components. On the basis of the results shown, the following deformation characteristics can be suggested for the TTC mode: a) Deformation twinning was dominated by compression and double twins, which were observed to form inhomogeneous deformation zones that trigger shear band formation (cf. Fig. 3 (RT) and Fig. 6). The relative ease of compression and double twinning, seen in this alloy, is in accordance with other studies on Mg-RE alloys [26] that reported promoted compression and secondary twin formation compared to conventional Mg-Al based alloys, where the activation stresses required to engage caxis compression are considerably much higher than those required for c-axis extension. b) The observed compression and double twinning in this alloy was likely accompanied by homogeneous dislocation slip activity in terms of basal and non-basal slip, both inside and in the vicinity of twins. This was evident by the massive low angle boundary formation (cf. Fig. 6b) and large orientation spreads (cf. Fig. 6d) indicated in - 11 -

the scatter data. By contrast, the high local stress concentrations in conventional Mg alloys loaded in compression along the c-axis are normally relaxed by the formation of few compression and double twins that subsequently trigger shear failure. The underlying reason for this is the rapid increase of geometrically necessary dislocation (GND) density at the twin-matrix interface, as a result of favored basal slip in the compression and double twins. An excess of GNDs produces high local stresses and subsequent failure. Deformation under IPC mode was aimed at engaging c-axis extension, and thereby activating {1012} extension twinning in the deformed specimens for subsequent annealing experiments. However, previously shown results in Figs. 3b and 4b reveal an unexpected twinning behavior, characterized by predominance of compression and double twins, similar to what was observed in the TTC mode. The occurrence of these twins during IPC loading can be explained by the weak starting texture and the large orientation spread (Fig. 2c). In addition, considering non-zero lateral stress components, grain orientations deviating from the ideal <0001> || LD orientation will experience a compressive stress component along their caxes [35]. Finally, internal stresses that are induced to maintain compatibility can lead to localized build-up of triaxial stress states that would also impose compression stress components in the c-axis. Similar to the TTC mode, intense slip activity was also evident in this deformation mode (cf. Fig. 4b). Fig. 7 shows a representative area of interest of one deformed specimen in IPC up to 6%. Fig. 7a is represented in terms of IPF colouring; Fig. 7b and c in terms of unit cell orientation of individual grains and grain orientation spread (GOS), respectively. The thick ‘purple’ band in Fig. 7a, labelled as region ‘II’ was in part indexed as a {1011} compression twin, originating from an adjoining parent orientation ‘I’ and ‘III’. A high density of low angle boundaries observed within and in close proximity of this band can be interpreted as an indication of dislocation slip activity. The high GOS values observed for region ‘II’ (cf. Fig. 7c) further corroborate the presence of slip activity. Fig. 7d shows the misorientation profile across regions ‘I’, ‘II’ and ‘III’ (see white dashed arrow in Fig. 7a). The measured misorientation angle at the boundary between regions ‘I’/‘II’ and ‘II’/‘III’ shows good agreement with the theoretical misorientation angle known for {1011} compression twinning i.e. ~ 56°. Fig. 7e depicts the scatter and recalculated pole figures of the sampled area in Fig. 7a. Evidently, the twin band of region ‘II’ displays a considerable scatter around the position of an ideal {1011} twin orientation, most likely arising from lattice rotations due to - 12 -

dislocation slip activity inside the twinned region. Ease of deformation slip is an important criterion when considering the twinning behavior, since twin formation is always accompanied by accommodation stresses at the twin tips and regions of intersection with the grain boundaries [19, 36]. Homogeneous deformation (activation of non-basal slip) would significantly reduce the accommodation work during nucleation of compression and double twins, hence facilitating their formation. From the perspective of material behavior, the influence of RE-alloying is also a major factor in the context of this seemingly ‘anomalous’ twinning behavior exhibited by the current alloy. The aspects associated with the influence of RE-addition on intrinsic and extrinsic stacking fault energies [37, 38], and interaction between solute RE atoms and dislocations [1, 39, 40] are particularly important. Additionally, there is the issue of particle twin interaction that has been readily investigated for {1012} extension twins but not yet for {1011} contraction and double twins [26]. Generally, if second phase particles cannot be sheared by the moving twinning ledge (during twin propagation), there will be some back-stress opposing the applied stress propelling twin growth. The level of this back-stress can vary depending on the plastic relaxation occurring within the twin as a result of dislocation slip. The latter is in turn determined by the effectiveness of particles in inhibiting slip, which greatly depends on the particle morphology and orientation. In the current work it was not possible to quantify the magnitude of the aforementioned factors, which we aim to pursue in more detail in a separate study. 4.2 Slip-induced twin fragmentation during cold deformation Fig. 7 shows an encircled region of interest in the EBSD map comprising a fragmented part of a former compression twin band that underwent secondary extension twinning and dynamic recovery giving rise to new recrystallized grains with high angle boundaries. It is also evident that some subgrains, with at least one low angle boundary, still remain in that area because they have not yet accumulated a sufficiently high misorientation relative to their deformed surrounding microstructure. The GOS map in Fig. 7c clearly shows that the average value of orientation spread (misorientation angle between all the points in a grain) in the fragmented band area is significantly lower than in the surrounding deformed bands displaying GOS values up to 4°. This difference in GOS values is important for partitioning the deformed and recovered/recrystallized regions (GOS ≤ 1°) in the cold deformed samples. As discussed earlier, deformation during plane strain compression in both TTC and IPC modes ensued by compression and double twinning accompanied by homogeneous slip. The - 13 -

latter was particularly prevalent in the soft-oriented bands conjugated with twins, which explains the preferential subgrain formation inside these bands (Fig. 7a), on the basis of dynamic recovery and dislocation rearrangement. Fig. 8 shows EBSD data of a selected, twinned area of one IPC specimen in the deformed state. Grains with an average GOS value ≤ 1° are displayed in green. Other grains corresponding to higher GOS values constitute the rest of the microstructure, which is represented in terms of IQ (grayscale shading) as the background. The regions partitioned in this way were further analyzed for basal slip SF values of deformed and recrystallized grains (Fig. 8b). It is evident that the majority of green-highlighted grains are associated with compression and double twins, which in terms of GOS directly corresponds to a low level of deformation and dislocation content in these areas. The outlined region in Fig. 8a is reproduced separately in terms of IPF map to illustrate a double twin (DT2) inside a ‘parent’ double twin (DT1). Such multiple-generation twinning events were seen to generate new, ideal twin orientations that reside on a well-defined fiber of basal poles in the (0002) pole figure shown in Fig. 8c with SF colour coding. The calculated high SF compression and double twin theoretical orientations (indicated by square and star symbols) originating from a parent matrix described by <0001> || LD, nicely overlap with the twin orientations determined experimentally from EBSD. From the SF map in Fig. 8b it is obvious that almost all of the compression and double twinned regions were highlighted in red, inferring a very high SF for the activation of basal slip in their interior. This is further quantified in Figs. 8c and d for deformed and recrystallized grains, respectively. A class width of 0.05 was used to plot the respective histograms. Apparently, the frequency peak in the deformed grains corresponded to an average SF value of 0.15, which is considered sufficient for activation of some basal slip in the matrix grains. However, in terms of massive dislocation activity, prolific basal glide was only possible in the confined compression and double twin bands that were found to dynamically fragment and recrystallize, giving rise to high fraction of new grains of soft orientations (Fig. 8d). The process of twin fragmentation, which can be regarded as a precursor stage to twin dynamic recrystallization, has been investigated elsewhere [27]. However, it is important to mention at this point that the fragmentation process of a twin band is realized by prismatic, and not by basal glide. This can be easily conceived from the more or less same, high basal SF value for all fragmented twin bands. Basal slip would be expected to cause a rotation of the basal plane, and thus, the c-axis toward the compression direction, - 14 -

which would give rise to local SF variations within the bands. Additionally, it would also result in a gradual decrease in the average SF value so that basal slip would eventually cease to operate. On the other hand, prismatic slip entails a rotation of the lattice around the c-axis, which does not affect the SF for basal slip, in accordance to what has been observed from the EBSD measurements (Fig. 8e). The magnified section in Fig. 8e further corroborates the suggested twin fragmentation mechanism, showing compression and double twin orientations forming a fiber about their c-axes. The additional points residing on the twin orientation fiber in the (0002) pole figure of recrystallized grains are associated with the new DRX grains arising from twin fragmentation. 4.3 Effect of TTC and IPC loading on the annealing microstructure development The onset of static recrystallization for the TTC (ε = 12%) and IPC (ε = 6%) specimens was observed at 325° C and 350° C, respectively, whereas in comparison, recrystallization in conventional Mg-Al-Zn alloys occurs at temperatures in the range of 200°C-250°C [8]. Such delayed recrystallization kinetics of the investigated Mg-1Nd alloy has already been attributed to the addition of RE elements, and is therefore not surprising. REs in the form of solutes and precipitates are known to pin dislocations and grain boundaries, thereby retarding recovery and recrystallization kinetics [41, 42]. Upon annealing, static recrystallization in the current work was observed to initiate inside compression and double twins that were not yet fragmented during deformation. Unlike extension twins, compression and double twins contribute to dislocation slip by producing crystallographically soft orientations that will be prone to maximum plastic deformation, and hence would be more liable to recrystallization than matrix grains and other extension twins of a hard orientation for slip. Additionally, the thin morphology of compression and double twins, owing to their energetic resistance to growth, results in smaller mean free paths for dislocation slip inside the twinned regions. Therefore, with increasing strain, dislocation generation leads to rapid accumulation in the dislocation density inside the twins and twin-matrix interface. Higher dislocation density, supplemented with homogeneous slip deformation in these regions propels the onset of local dynamic recovery at ambient temperatures and could promote recrystallization under thermal activation. In the present study, samples loaded in TTC exhibited fully recrystallized microstructures at 350°C, whereas the IPC mode displayed an even more sluggish microstructure transition with complete recrystallization attained only after annealing at 450°C. Such a large difference

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in recrystallization kinetics as a function of loading condition could be attributed to the following aspects: a) The imposed strains in TTC were higher than those in IPC, which implies a higher available driving force for recrystallization under TTC deformation. b) The TTC mode primarily engaging c-axis compression would be expected to nucleate relatively higher fractions of compression and double twins, leading to a higher density of recrystallization nucleation sites available during annealing. A greater nucleation density would give rise to earlier impingement of recrystallizing grains, at which recrystallization transitions into grain growth. Prior to full recrystallization at 450°C in the annealed IPC specimens, annealing at 400° C produced a bimodal microstructure exhibiting regions with massive, unrecrystallized twinning (A) and largely recrystallized zones (B). These regions are roughly demarcated by yellow dashed lines in Fig. 9a. In addition, selected areas ‘1’ and ‘2’ for EBSD analysis are outlined in dashed rectangles. Fig. 9b, d and f correspond to the recrystallized region ‘1’, whereas Fig. 9c, e and g represent the heavily twinned region ‘2’. The IPF mapping shown in Fig. 9b reveals a primarily recrystallized microstructure, with few deformed grains still withstanding recrystallization (cf. Fig. 9f, GOS map). The related scatter and recalculated pole figures (cf. Fig. 9d) indicate a predominance of texture components between LD and CD. The recrystallization texture corresponding to EBSD correlates well with the XRD bulk texture (cf. Fig. 3b). The IPF mapping of region ‘2’ (Fig. 9c) shows a deformed microstructure with highly deformed, twinned regions and few recovered/recrystallized grains of large size (~ 100 µm) (cf. Fig. 9g). The deformed twins corresponded to {1012} extension twins with 86° reorientation from the prismatic orientation of parent matrix grains (cf. Figs. 9c and e), assuming a hard orientation for easy dislocation slip, and thus, an unfavorable prerequisite for recrystallization. Annealing at 450°C resulted in significant grain growth in both TTC and IPC annealed specimens. While upon annealing of TTC specimens the occurrence of grain growth was evident by a transition from a fine into a coarse grained microstructure, annealing of IPC specimens displayed an abrupt change from a partially recrystallized state into a coarse microstructure, highlighted by strong grain growth. Fig. 10a and c represent a grain size analysis on the basis of basal SF values for both loading modes, after annealing at 450° C. For the TTC mode, the relative volume fractions corresponding to the high and low SF grains - 16 -

were 78% and 22%, respectively. The measured average grain sizes at 450° C, were 47.8 ± 0.46 μm for low SF grains vis-à-vis 54.1 ± 0.52 μm for the high SF ones. Correlating with the single orientation scatter and recalculated pole figures (cf. Fig. 10b), the low SF grains seem to correspond primarily to basal orientations with maximum scatter of 20° from CD, whereas the dominant high SF grains form a ring of off-basal orientations lying between 30° and 60° from CD. Similar analysis for the IPC mode also revealed a higher fraction of high SF grains (~70%) in comparison to the low SF one (~30%). Furthermore, like the TTC mode the average grain sizes for the low SF and high SF grains after annealing at 450°C, were measured as 47.6 ± 0.52 μm and 55.9 ± 0.48 μm, respectively (cf. Fig. 10c). The related EBSD textures in Fig. 10d indicate that the high SF regions primarily corresponded to the compression and double twin orientations, whereas the low SF ones mostly reflected the initial parent orientations. A basal orientation, corresponding to {1012} extension twins was diminished after this particular annealing treatment (cf. Fig. 9c and e, annealing at 400°C). The recrystallization and grain growth kinetics observed in the current study are both marked by a significantly delayed onset. In order to understand the overall reluctant nature of recrystallization and grain growth in the present alloy, solute effects of Nd on static recrystallization during annealing need to be examined. The maximum solid solubility of Nd in magnesium at 548°C is approx. 3.62 wt. % (0.63 at. %) [29]. With decreasing temperature the solubility limit drops sharply to 0.7 wt. % at 400°C, reaching almost zero at ambient temperature [29]. This behavior of drastic variance in solubility limit with respect to temperature suggests that Nd would be expected to display both solute and precipitate effects. Microstructural characterization in the deformed state (both TTC and IPC) indicated dense distribution of fine particles that were in part lying in long stringers (Fig. 3), probably related to the processing history of the material. Fig. 11a represents the variation of grain size with respect to the annealing temperature. At temperatures between 300°C and 400°C, the maximum solubility of Nd in solid Mg is lower than 1 wt. %, whereby Nd can exist both in the form of solute and in precipitates. In such case, factors affecting grain boundary mobility, such as solute drag, grain boundary segregation and second-phase Zener pinning could be expected, thus explaining the delayed onset of recrystallization. At 400°C, the measured average recrystallized grain sizes for specimens subjected to TTC and IPC were 6 μm (Standard Error (SE) = 0.1 μm) and 12 μm (SE = 0.6 μm), respectively, indicating strong pinning behavior. Secondary electron images obtained after annealing at 400° C for 60 minutes revealed a distribution of very fine, round - 17 -

particles (< 1 μm) along the grain boundaries, ideal for pinning grain boundary migration (Fig. 11b). The somewhat smaller grain size in the TTC specimen can be attributed to the previously mentioned difference in the nucleation density of recrystallization, causing earlier impingement and ensuing grain growth. On the other hand, the IPC specimen, still displaying a partially recrystallized microstructure at 400°C, remains by the stage of nuclei growth by means of the migration of high angle grain boundaries driven by the stored energy of deformation, which is orders of magnitude higher than the curvature-driven grain growth. From Fig. 3 and Fig. 9d the TTC annealing texture at 400°C indicated a greater density of orientations with off-basal than basal character. In terms of preferential recrystallization, it can be argued that the high dislocation densities in the vicinity of contraction and double twins provide a higher driving force for the growth of off-basal nuclei as compared to basal recrystallizing grains nucleating from extension twins or the matrix. Bearing a high misorientation with the deformation microstructure, the off-basal nuclei are also believed to possess some advantage in overcoming pinning, and hence, grow more preferably in comparison to the basal orientations. The pinning effect would, however, still be strong enough to restrict free growth of recrystallized grains to large sizes (cf. Fig. 11a). After annealing at 450° C, the obtained average grain sizes indicate a fivefold to eightfold increase for both IPC and TTC specimens reaching ~ 55 µm (Fig. 11a). The maximum solid solubility of Nd in magnesium increases to ~ 1.2 wt. % at 450°C [29], implying sufficient driving force to prompt re-dissolution of the fine Mg12Nd precipitates decorating the grain boundaries as evident from Fig. 11c, serving as a representative example for the whole specimen. The thermal activation at 450°C would also empower the impeded grain boundary to break free of the existing impurity drag. Reduced influence of particle pinning and solute drag would hence promote unhindered grain growth. This is displayed in the respective TTC specimen by almost balanced growth of all nucleated orientations resulting in the appearance of both basal and off-basal texture components (cf. Fig. 3a and 5e). In the IPC specimen, recrystallized grains nucleated from compression and double twins seemed to have an obvious growth advantage over other orientations, which get consumed in the process (cf. Fig. 5a and 10d).

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5 Conclusions The deformation and recrystallization behavior of rolled Mg-1Nd alloy can be summarized in terms of the following key characteristics: a) Deformation in Mg-1Nd displayed competitive activation of both {1012} extension and {1011} contraction (compression) twinning that gave rise to soft and hard texture components. b) The predominance of compression and {1011}-{1012} double twins in the deformed IPC microstructure, and their dynamic fragmentation into fine grains by means of progressive basal and prismatic slip led to the formation of a prismatic texture fiber of recrystallized grains. c) The deformed TTC microstructure exhibited a less common basal-type texture, characterized by a significant LD spread, and an asymmetrical basal poles splitting around the TD by about -5° and +25°. d) Texture evolution during annealing was associated with general texture weakening. 300°C annealing retained the primary deformation texture components since no recrystallization took place. The existence of strong impurity drag and Mg12Nd particle pinning delayed the onset of recrystallization till 325°C for TTC and 350°C for IPC specimens. Recrystallization was observed to nucleate preferably in compression and double twins. e) Related recrystallized grains possessed a high Schmid factor (SF) for basal slip, and were more frequent and larger on average than competing low SF grains related to extension twins or the matrix. f) During the 450°C anneal, the pinning force on the grain boundaries was much reduced due to precipitate re-dissolution and higher thermal activation for grain boundary mobility leading to unrestricted grain growth in both TTC and IPC specimens. While the TTC microstructure revealed almost unequivocal growth of basal and off-basal orientations, grain orientations related to former compression and double twins dominated the growth kinetics in the annealed IPC microstructure, as a result of a grain size advantage that emerges during recrystallization. Acknowledgement The financial support of the Deutsche Forschungsgemeinschaft (DFG), Grant No. AL 1343/12 is gratefully acknowledged. - 19 -

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Fig. 1 Processing pathway employed in the present work, steps (a) – (f) Fig. 2 True stress-strain curves corresponding to TTC (in blue) and IPC (in green) deformation of Mg-1Nd (a); Starting (0002) pole figures for TTC deformation (b) and IPC deformation (c); Pre-PSC Optical microstructure obtained after rolling and recrystallization annealing at 400° C for 12 min Fig. 3 XRD textures in terms of (0002) pole figures and optical microstructures of Mg-1Nd for deformed and annealed states corresponding to TTC (a) and IPC (b) modes

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Fig. 4 EBSD image quality (IQ) maps of Mg-1Nd after TTC deformation and annealing at 325° C (a); and IPC deformation and annealing at 350° C (b); {1012} extension twin boundaries in red, {1011} compression twin boundaries in blue and {1011}-{1012} double twin boundaries in yellow are further highlighted in Figs. 4a and b. (0002) single orientation scatter data and recalculated pole figures of recrystallized and deformed subsets for TTC (c)-(d) and IPC modes (e)-(f) Fig. 5 EBSD inverse pole figure (IPF) mapping of Mg-1Nd after TTC deformation and annealing at 450° C (a), and IPC deformation and annealing at 450° C (d); (0002) single orientation scatter and recalculated pole figures for IPC (b) and TTC (e) modes. Basal schmid factor (SF) maps corresponding to IPC (c) and TTC (f) modes Fig. 6 EBSD IQ map for Mg-1Nd after TTC deformation, showing {1012} extension twin boundaries in red, {1011} compression twin boundaries in blue and {1011}-{1012} double twin boundaries in yellow; corresponding IPF mapping for Mg-1Nd after TTC deformation (b), magnified region depicts matrix reorientation after double twinning leading to basal and off-basal orientations; Basal SF maps (c) corresponding to the microstructure shown in Fig. 6a; (0002) single orientation scatter and recalculated pole figures for Mg1Nd after TTC Fig. 7 EBSD IPF (a), unit cell (b) and GOS (c) mappings showing low angle grain boundary formation inside compression twin during IPC deformation, encircled area indicates evidence of slip aided dynamic recovery seen by appearance of grains with low GOS values (Fig. 7c); misorientation profile across the compression twin ‘II’ bounded by parents ‘I’ and ‘III’ (d); (0002) single orientation scatter and recalculated pole figures corresponding to the region in Fig. 7a Fig. 8 EBSD IQ (a) and Basal (SF) mappings for IPC deformed Mg-1Nd, showing sequential double and compression twins; grains from DRX highlighted in green in Fig. 8a; (0002) single orientation scatter overlaid with theoretically possible compression(with calculated SF values) and double twin variants and basal SF distribution for deformed (c) and recrystallized (d) subsets; unit cell representation of Fig.8a in IPF coloring indicating rotation of twin orientations about their c-axes due to active prismatic slip (e) Fig. 9 Optical microstructure of Mg-1Nd at 400° C after IPC deformation (a); (b) and (c) show EBSD IPF mappings corresponding to regions ‘1’ and ‘2’, respectively; EBSD texture and GOS mappings corresponding to region ‘1’ ((d) and (f)) and region ‘2’( (e) and (g)) Fig. 10 Grain size distribution for high and low basal SF grains after annealing at 450° C for TTC (a) and IPC (c) modes; EBSD texture data for high and low basal SF grains for TTC (b) and IPC (d) modes Fig. 11 Grain size variation with annealing temperature for both TTC and IPC modes (a); SE images after 400° C annealing showing fine precipitates (<1 µm) decorating the grain boundaries (b); SE images after 450°C annealing indicating disappearance of particles at the grain boundary (c)

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