Deformation and recrystallization mechanisms of Mg-Al-Nd and Mg-Zn-Nd alloys

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Deformation and recrystallization mechanisms of Mg-Al-Nd and Mg-Zn-Nd alloys

vorgelegt von M. Eng.

Xun Zeng

DQ der Fakultät III – Prozesswissenschaften der Technischen Universität Berlin zur Erlangung des akademischen Grades

Doktor der Ingenieurwissenschaften -Dr.-Ing. -

genehmigte Dissertation

Promotionsausschuss:

Vorsitzender: Dr. Sören Müller Gutachter: Prof. Dr. Claudia Fleck Gutachter: Prof. Dr. Karl Ulrich Kainer

Tag der wissenschaftlichen Aussprache: 26. JulL 2021

Berlin 2021

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Abstract

In the present study, the influence of the rare earth (RE) elements addition on the deformation and recrystallization mechanisms of Mg alloys were investigated using Mg-Nd based alloys. The main focus of the study was given to unveil the interrelationship between the alloy compositions, microstructure and texture developments during deformation and recrystallization annealing, and the resulting mechanical properties. To reach the complex goals, the quasi in-situ tension and annealing experiments were carried out in combination with the electron backscatter diffraction (EBSD) measurements on the Mg-Nd based alloys with the addition of Al (AN12) or Zn (ZN12). The contributions of different deformation modes could be determined by analyzing the slip traces and the twin variant selections occurred during the uniaxial tension. The microstructure and texture evolution during recrystallization annealing were investigated with respect to the distinct nucleation and grain growth processes between the examined alloys.

The main conclusions are summarized as follows: the ZN12 alloy shows better formability with a higher fracture strain and strain hardening exponent than those of the AN12 alloy. When comparing the grains with similar sizes and Schmid factor (SF) values, higher activities of dislocation slip and twinning are observed in the ZN12 than the AN12. Basal slip is the dominant slip mode in both samples. Non-basal slip and tensile twinning are already activated in the early deformation stage (ε = 0.02) in the ZN12 sample, whereas they are rarely observed in the AN12 sample. It is supposed that the ZN12 sample has a lower CRSS for basal slip and CRSS ratio of non-basal to basal slip compared with the AN12 sample. The strain accommodation is more difficult in the AN12 sample due to insufficient deformation modes, leading to premature cracks.

During the recrystallization annealing, the strong basal-type texture is retained in the hot-rolled AN12 sample, while the RD spread texture component becomes less distinct in the ZN12 sample. Secondary twins act as the preferred sites of the nucleation in the AN12, and the nuclei grow rapidly by consuming the twins.

The recrystallization nuclei show the orientation corresponding to the RD spread texture, which is similar to the secondary twin orientation. Apart from the twin induced nucleation, the strain induced grain boundary migration is also observed in the AN12 sample which strengthens the deformation texture. On the contrary, the ZN12 sample shows a retarded recrystallization process due to the precipitation at the grain and twin boundaries. The precipitates at twin boundaries restrict the growth of recrystallization nuclei within the secondary twins. Apart from secondary twins, ternary twins are found in the ZN12 sample, which result in the recrystallization nuclei with scatter basal poles and a weak texture. In the early annealing stage, the formation of dislocation cells is also found in the ZN12 sample. The merging of the dislocation cells is difficult owing to the low dislocation density within the cells.

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Zusammenfassung

In der vorliegenden Arbeit wurde der Einfluss des Zusatzes von Seltenen Erden (RE) auf die Verformungs- und Rekristallisationsmechanismen von Mg-Legierungen anhand von Mg-Nd-Basislegierungen untersucht.

Der Fokus dieser Studie lag auf der Untersuchung der Zusammenhänge zwischen den Legierungszusammensetzungen, der Mikrostruktur- und Texturentwicklung während der Verformungs- und Rekristallisationsglühung und den daraus resultierenden mechanischen Eigenschaften. Um die komplexen Ziele zu erreichen, wurden die quasi in-situ Zugversuche und Wärmebehandlungen in Kombination mit den Elektronenrückstreuungsbeugungsmessungen (EBSD) an den Mg-Nd-Basislegierungen mit dem Zusatz von Al (AN12) oder Zn (ZN12) durchgeführt. Die Beiträge der verschiedenen Verformungsmodi konnten durch die Analyse der Gleitlinien und der verschiedenen Zwillinge, die während der einachsigen Zugversuche auftraten, bestimmt werden. Das Mikrostruktur und die Texturentwicklung während des Rekristallisationsglühens wurden im Hinblick auf die unterschiedlichen Keimbildungs- und Kornwachstumsprozesse zwischen den untersuchten Legierungen analysiert.

Die wichtigsten Schlussfolgerungen sind: Die ZN12-Legierung zeigte eine verbesserte Umformbarkeit mit einer höheren Bruchdehnung und einem höheren Verfestigungsexponenten als die AN12-Legierung. Selbst beim Vergleich der Körner mit ähnlichen Größen und Schmid-Faktor (SF)-Werten wurden höhere Aktivitäten des Versetzungsgleiten und der Zwillinge in der ZN12 als in der AN12 beobachtet, während basales Gleiten der dominante Gleitmodus in beiden Proben ist. Nichtbasales Gleiten und Zugzwillinge werden in der ZN12-Probe bereits in der frühen Verformungsphase (ε = 0,02) aktiviert, während sie in der AN12-Probe nur selten beobachtet werden. Es wird vermutet, dass die ZN12-Probe im Vergleich zur AN12- Probe ein geringeres CRSS des basalen Gleiten und ein geringeres CRSS-Verhältnis von nichtbasalem zu basalem Gleiten aufweist. Die Dehnungsaufnahme ist in der AN12-Probe aufgrund unzureichender Verformungsmodi schwieriger, was zu den vorzeitigen Bruch führt. Während des Rekristallisationsglühens wird die starke basale Textur in der warmgewalzten AN12-Probe beibehalten, während die in Walzrichtung gekippte Texturkomponente in der ZN12-Probe weniger deutlich wird. Sekundäre Zwillinge fungieren als die bevorzugten Orte der Keimbildung im AN12, und die Keime wachsen schnell, indem sie die Zwillinge aufbrauchen. Die Rekristallisationskeime zeigen die der in Walzrichtung gekippte Textur entsprechende Orientierung, die der Orientierung der sekundären Zwillinge ähnlich ist. Abgesehen von der zwillingsinduzierten Keimbildung wurde in der AN12-Probe auch die dehnungsinduzierte Korngrenzenwanderung beobachtet, die die Verformungstextur verstärkte. Im Gegensatz dazu zeigte die ZN12-Probe einen verzögerten Rekristallisationsprozess aufgrund von Ausscheidungen an Korn- und Zwillingsgrenzen. Die Ausscheidungen an den Zwillingsgrenzen schränken das Wachstum der Rekristallisationskeime innerhalb der sekundären Zwillinge ein. Neben den sekundären Zwillingen wurden

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in der ZN12-Probe auch ternäre Zwillinge gefunden, die zu Rekristallisationskeimen mit zerstreuten Basalpolen und einer schwachen Textur führten. In der frühen Glühphase wurde die Bildung von Versetzungszellen auch in der ZN12-Probe gefunden. Die Verschmelzung der Versetzungszellen ist aufgrund der geringen Versetzungsdichte innerhalb der Zellen schwierig.

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Acknowledgements

I gratefully acknowledge the financial support of the China Scholarship Council (CSC).

I would like to express my gratitude to everyone who encouraged and supported me throughout my doctoral studies. This thesis would not be finished without their guidance and support.

Most of all, I am deeply grateful to my supervisor Prof. Dr. Karl Ulrich Kainer for his inspiring and considerate guidance, patience and support for my research. I also gratefully thank Prof. Dr. Claudia Fleck for giving me the opportunity to carry out my doctoral research under her supervision in Technische Universität Berlin. I do appreciate Dr. Sören Müller and Prof. Walter Reimers for the support of defense and doctoral procedures.

My sincerely gratitude also goes to Dr. Dietmar Letzig for his enlightening guidance and fruitful discussions during this work. I would like to thank Dr. Sangbong Yi for his supports and advices that improve my understanding of the deformation and recrystallization behaviors of Mg alloys. I would like to thank Dr. Jan Bohlen, Dr. Gerrit Kurz and Dr. Jose Victoria-Hernandez for the interesting discussions and the help with lab facilities, experiments.

I appreciate the help of colleagues with whom I have been working during the past four years at Helmholtz- Zentrum Geesthacht: Dr. Sumi Jo, Dr. Sangkyu Woo, Mr. Alexander Reichart, Mr. Stefan Koch, Mr.

Guadalupe Cano, Miss. Maria Nienaber, Mr. Changwan Ha, Miss Jing Li, Ms Yukyung Shin. Thank all of you for sharing your knowledge and helpful discussions.

Finally, I would like to thank my parents for their love and support.

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List of abbreviations

Al Aluminum

AN12 Mg-1.0Al-1.7Nd

CRSS Critical resolved shear stress

CDRX Continuous dynamic recrystallization DDRX Discontinuous dynamic recrystallization

DRX Dynamic recrystallization

DTIN Deformation twin induced nucleation

EBSD Electron backscatter diffraction

ED Extrusion direction

FIB Focused ion beam

GND Geometrically necessary dislocations

GOS Grain orientation spread

HAGB High angle grain boundary

HCP Hexagonal close-packed

HRDIC High resolution digital image correlation

LAGB Low angle grain boundary

Mg Magnesium

ND Normal direction

PSN Particle stimulated nucleation

RD Rolling direction

RE Rare earth

SBIN Shear bands induced nucleation

SEM Scanning electron microscopy

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SIBM strain induced grain boundaries migration

SRX Static recrystallization

TEM Transmission electron microscopy

TD Transverse direction

TDRX Twinning dynamic recrystallization

TKD Transmission Kikuchi diffraction

SF Schmid factor

UTS Ultimate tensile strength

VPSC Visco-plasitic Self Consistant

Zn Zinc

ZN12 Mg-1.0Zn-1.7Nd

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List of symbols

b Burgers vector

K Twinning invariants shear plane

η Twinning invariant shear direction

σ True stress

σy Yield strength

ε True strain

f Number fraction of twins

Fx Area fraction of the recrystallized grains

m Schmid factor, see Eq. 1-1

φ Angle between the stress and the normal of glide plane, see Eq. 1-1 λ Angle between the stress and the normal of glide direction, see Eq. 1-1

m’ Geometrical compatibility factor, see Eq. 2-1

ϕ Angle between two slip directions of adjacent grain pair, see Eq. 2-1 κ Angle between slip plane normal directions of adjacent grain, see Eq. 2-1

g growth rate, see Eq. 2-2

G Grain size of recrystallized grain, see Eq. 2-2

θ Strain hardening rate, see Eq. 3-1

K Strength coefficient, see Eq. 3-2

n Strain hardening exponent, see Eq. 3-2

τbasal CRSS values of basal slip, see Eq. 4-1

mbasal Average SF values of basal slip, see Eq. 4-1

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1 Introduction

1.1 Magnesium and its alloys

Due to the increasing demands of weight savings and environmental-friendly material, Magnesium (Mg) and its alloys are regarded as promising candidates for this century [1]. The application of Mg in automobile, aerospace industry is extending owing to its excellent properties of low density, high specific strength and good castability [2]. Besides, researchers investigate the usage of Mg alloys for biomedical purpose like biodegradable implants as they show reasonable corrosion behavior and low toxicity [3].

To modify the mechanical properties, like strength, ductility, creep resistance and elastic properties, of Mg alloys, diverse alloy elements are usually introduced for industrial applications. Aluminum (Al) can enhance the strength by the solid solution strengthening effect. Zinc (Zn) element increases fluidity in casting, corrosion resistance and room temperature strength. The AZ series, a combination of Al and Zn into Mg, are well studied. The AZ91 is the most commonly used die casting Mg alloy with good castability, high specific strength [4]. Zirconium (Zr) is known as an excellent grain refiner, which is beneficial to mechanical properties. Lithium (Li) addition introduces the phase transition to ductile body-centered cubic Li-rich phase [5]. Besides, Li decreases the c/a ratio of Mg alloys, leading to an enhanced non-basal slip activity [6]. A small amount of Calcium (Ca) improves the creep resistance [7] and rollability of Mg alloy sheets.

Compared with the conventional Mg alloys, Rare earth (RE) elements containing Mg alloys weaken the crystallographic texture and improve the mechanical performance [8].

According to the manufacturing methods, Mg alloys are usually classified as casting Mg alloys and wrought Mg alloys. More than 95% of Mg alloy products are produced by die casting, while the application of wrought Mg alloys is very limited due to the low formability at room temperature [4]. Casting Mg alloys normally have better productivity and save the production cost, but they show lower mechanical properties compared to the wrought counterpart. The demand for high-performance structural materials promotes the development of wrought Mg alloys. The poor formability of wrought Mg alloys, resulting from the intrinsic feather of hexagonal close-packed (HCP) crystallographic structure, restricts their application. The HCP structure leads to a strong anisotropy [9]. The tendency of forming a strong texture of Mg alloys after fabrication processes, like rolling and extrusion, deteriorates the ductility of semi-finished products [10].

On the other hand, the high critical resolved shear stress (CRSS) of non-basal slip in Mg alloys at room temperature restricts the dislocations glide dominantly on basal plane [11]. Basal slip has only two independent slip systems, and fails to fulfill the von Mises criterion that requires five independent slip systems to undergo a homogeneous strain [12].

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An effective approach to enhancing the formability of wrought Mg alloys is to reduce the anisotropy by altering the basal-type texture to a more randomized texture. The crystallographic orientation has a key impact on the activation of deformation modes. A randomized texture facilitates different deformation modes in each grain to accommodate homogeneously the strain. The texture randomization can be achieved by applying specific processing routes such as cross rolling [13], differential speed rolling [14] and twin roll casting [15]. However, these methods of severe plastic deformation require special processing facilities and processing routes, which takes a high cost and time-consuming processes. The addition of rare earth elements is well acknowledged as an efficient way of weakening the texture of Mg alloys. Researchers showed great interests to develop Mg-RE alloys for extrusion and sheet applications [8, 9, 16, 17].

1.2 Deformation mechanisms of Mg alloys

Deformation of materials includes elastic deformation and plastic deformation. The elastic deformation will be restored when the applied force is removed while the plastic deformation is irreversible. Deformation modes such as dislocation slide and twinning are responsible for the plastic deformation. The activation of certain deformation modes is strongly dependent on the temperature, load, microstructure and crystallographic texture. The understanding of the deformation mechanisms is important to improve the strength, ductility and toughness of the materials. The early studies of the deformation mechanisms focused on single crystal of pure Mg [18]. However, the mechanical behavior of polycrystal Mg alloys varies away from the single crystal, for example, prismatic slip trace was observed near the grain boundary in deformed polycrystal Mg samples [19] though the study of Mg single crystal showed the critical resolved shear stress (CRSS) of prismatic slip is 50-100 times higher than basal slip [20]. In polycrystals, the slip plane and slip direction of dislocations vary between each grain in dependence on its crystallographic orientation. When the dislocations reach grain boundary, they cannot easily transmit to the neighboring grain because of the mismatched lattice. These dislocations accumulate near the grain boundary, leading to a localized stress.

The localized stress, higher and differs from the applied stress, is responsible for the activation of difficult slip systems like prismatic and pyramidal slip. With the development of experimental technology like electron backscattered diffraction (EBSD), in-situ neutron diffraction and high resolution digital image correlation (HRDIC), researchers are able to reveal the underlying deformation mechanisms of Mg alloys.

1.2.1 Crystallography of Mg

Mg and its alloys have a HCP crystal structure, with the space group of P63/mmc. The crystal structure is schematically shown in Figure.1.2.1. The HCP structure corresponds to an ABAB stacking of the atomic layers. The unit cell of HCP structure consists of three axis, a1, a2 and c. The length of the axis is a1 = a2 ≠ c, while the angles between the axis are α = β = 90 °, γ = 120 °. In pure Mg, the lattice parameters are a =

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0.32092 nm and c = 0.52105 nm [21]. The c/a ratio in pure Mg is 1.623. The Li addition leads to a reduction of the c/a ratio, resulting in the change of close packed plane from basal to prismatic plane [22]. The atomic positions projection to the basal plane is illustrated in Figure. 1.2.1 (b).

Figure.1.2.1 Schematics illustrations of: (a) hexagonal close-packed structure, (b) atom projection onto the basal planes.

1.2.2 Slip systems in Mg alloys

In materials science, plastic deformation is usually achieved by the motion of dislocations, namely slip. Slip systems describe the set of slip planes and slip directions for which dislocation motion easily occurs. The Burgers vector, denoted as b, represents the direction and magnitude of lattice distortion caused by a dislocation. Dislocations usually slide on the close-packed planes and in close-packed directions because of a small magnitude of Burgers vector. The smaller magnitude of Burgers vector indicates a smaller energy barrier for atoms to move from one position to the next position. In Mg alloys, (0001) plane and <112�0>

direction are the close-packed plane and direction, also called basal plane and <a> direction. Apart from that, slip can occur on the {101�0} prismatic planes which are perpendicular to the basal planes, or the inclined pyramidal {112�1} and {112�2} planes. The <c+a> Burgers vector of pyramidal slip can accommodate strains along [0001] axis, c-axis, which is the very important deformation mode at high temperature. Slip systems in Mg alloys are normally classified into three categories, basal <a>, prismatic

<a> and, pyramidal slip <c+a> slip. The schematics of different slip systems are shown in Figure. 1.2.2.

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Table. 1.2.1 Three main slip systems in Mg alloys.

Slip systems Slip planes Slip directions Burgers vector Magnitude

Number of independent

slip system

Basal (0001) <112�0> 𝑎𝑎⃑ |𝑎𝑎| 2

prismatic {101�0} <112�0> 𝑎𝑎⃑ |𝑎𝑎| 2

{101�1} <112�0> 𝑎𝑎⃑ |𝑎𝑎| 4

pyramidal {112�1}

<112�3> 𝑐𝑐⃑+𝑎𝑎⃑ �|𝑎𝑎|²+ |𝑐𝑐|² 5 {112�2}

Figure. 1.2.2 Slip systems in Mg alloys: (a) <a> slip on basal, prismatic and pyramidal planes, (b) <c+a>

slip on pyramidal planes. Slip planes and slip directions are marked as transparent frames and thick black arrows, respectively.

Basal slip is the most common slip system at room temperature due to the lowest CRSS compared with other slip systems. This slip system has a Burgers of <112�0> vectors and glides on (0001) basal planes, which are the close-packed direction and planes in pure Mg. There are three <a> Burgers vector on the basal plane, however, only two of them are independent. According to the Von Mises Criterion, five independent slip systems are required for uniform polycrystalline deformation. This is why basal slip is usually discussed with other deformation modes like prismatic slip and twinning.

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At room temperature, it is believed that prismatic slip can hardly occur in Mg alloys because its CRSS was 50-100 times higher than basal slip [23]. However, a considerable amount of prismatic slip traces were observed at room temperature deformation by Hauser [22]. The Burgers vector of prismatic slip is <112�0>

vector which is the same with basal slip. It provides another two independent slip systems. However, both basal slip and prismatic slip with <a> Burgers vectors cannot accommodate the strain along c-axis.

Pyramidal slip has a Burgers vector of <c+a> either on 1^st order {101�1} or 2^nd order {112�2} planes. It is believed to play a minor role in the deformation process due to its large Burgers vector. In fact, pyramidal slip was present in all Mg alloys with various deformation temperature [24]. Pyramidal slip tends to dissociate into partial dislocations to reduce the high energy. Pyramidal slip is the only slip system that can accommodate the strain along c-axis, and thus is important for the homogeneous deformation [25].

Temperature is one of the key factors that affect the activation of deformation modes. In the crystal lattice, the defects like precipitates and solute atoms hinder the dislocation motion by producing short-range stress.

When the temperature is increased, the dislocations within the material have sufficient energy to overcome these short-range stresses, in other words the CRSS is reduced. Compared with other deformation modes, basal slip is less sensitive to the temperature. The difference of CRSS between basal slip and other deformation modes rapidly dropped with increasing temperature, as is shown in Figure. 1.2.3 [26]. This phenomenon can be explained with cross slip from basal to non-basal planes and will be discussed later.

The improved ductility of Mg alloys derives from the cooperation of various deformation modes at elevated temperature.

Figure. 1.2.3 The relationship between CRSS values of different deformation modes and temperature in literature [26].

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The activation of slip systems is influenced by many factors such as solid solution, precipitations and texture.

The solution atoms introduce mismatch and stress field into the lattice because of the difference in atomic radius. This mismatch will hinder the movement of dislocations. A high level of stress is necessary for dislocations to overcome this obstacle, so called solid solution strengthening mechanism. The experimental work shows that rare earth elements improve the creep resistance by strengthening the basal slip [27]. The solution softening of prismatic slip was firstly reported by Dorn [22]. They associated the improved ductility by Li addition with the enhanced activity of prismatic slip. The reduced c/a ratio [28] and cross slip [29]

mechanisms are the most likely reasons for the phenomenon.

The geometric relationship between precipitate and basal plane plays an important role in the aging hardening behaviors. Nie [30] proposed that the plate shape precipitates on basal planes, which are the most common precipitates like Mg17Al12 in Mg alloys, are the least effective precipitate geometry for pinning basal dislocations. The precipitates in Mg-RE alloys, such as WE54 and WE43, are plate shape and located on prismatic planes. These precipitates are the most effective to impede basal slip. As prismatic and pyramidal slip are considered as complement deformation modes to basal slip, it is difficult to exclude the effect of precipitates on non-basal slip from basal slip.

The empirical Hall-Petch relationship, smaller grain size better mechanical properties, has been developed for Mg and its alloys [31]. Decreasing the grain size will decrease the dislocation concentration at the grain boundaries, and increase the applied stress necessary to move a dislocation across a grain boundary. There are controversial views on the grain size effect on slip systems. One is that smaller grains suppress basal slip more than prismatic and pyramidal slip [32]. Prismatic slip was mostly observed within some micrometers to the grain boundaries due to the highly localized stress in boundary region. Giving a certain dimension of a material, the one with a smaller grain size microstructure has a higher volume fraction of preferred sites for prismatic and pyramidal slip. In another world, the activity of prismatic and pyramidal slip will be increased. However, the study by Jain [33] showed a similar level of basal vs non-basal slip activity in different grain size samples.

Due to the relatively low basal stacking fault energy, the basal dislocation tends to dissociate into partial dislocations. Thus, the basal to non-basal planes cross slip, which requires the constriction of partial dislocations is less likely to occur in Mg alloys. Nevertheless, the activation of cross slip was reported to play an important role in the deformation behaviors [34]. Yoshinaga [35] proposed the jog-pair mechanism to explain the cross slip in Mg alloys, as is shown in Figure. 1.2.4. Screw <a> dislocations on primary basal plane firstly constricted in some atomic length scale. It can cross slip onto prismatic plane and dissociate in another basal plane.

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Figure. 1.2.4 Schematic illustrations of jog-pair mechanism: (a) dissociate basal screw slip constrict and cross glide onto prismatic plane, (b) the non-basal part of the glide dislocation cross slip onto basal planes

again, (c) the basal screw dislocation entirely climb to the neighboring basal plane [35].

1.2.3 Twinning systems in Mg alloys

Apart from dislocation slips, twinning also plays a critical role in the deformation behaviors of Mg alloys, for example, the tension compression anisotropy. Unlike dislocations which can glide along both directions of Burgers vector, twinning is polar. Twinning initiates with the formation of a nucleus at grain boundaries or twin boundaries (e.g., secondary twinning). The initial rapid growth of nucleus derives from the high stress concentration at grain boundaries. After that, the propagation of twins is controlled by the twinning dislocations in the twin interface. Normally, the stress for twin nucleation is higher than that required for twin growth [36]. The crystallography of twinning is explained with twinning elements, as is shown in Figure. 1.2.5 (a). K1 is the first invariants shear plane, containing the twinning shear direction η1. K2 and η2

refer to the second invariants. Figure. 1.2.5 (b) exhibits the most common {101�2} tensile twinning relationship with a rotation of 86° around <112�0> axis in Mg alloys.

Figure. 1.2.5 Schematic illustrations of (a) twinning elements, (b) tensile twinning elements in Mg alloys [37].

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As is reported by Roberts [38], deformation twinning in Mg alloys can occur on various planes like {101�1}, {101�2}, {112�1}, {112�2}. Depending on how the stress is applied to the c-axis (or a-axis) of the crystal, primary twinning systems are classified into tension and compression twinning. When the c-axis is stretched with a tension stress or the a-axis is compressed, the tensile twinning can occur. On the contrary, compression twins are favored when the c-axis is compressed or a-axis is stretched. Among those, the primary {101�2} and {101�1} twinning, with a rotation of 86.3° and 56° around <112�0> axis, are the most common tension and compression twinning modes in Mg alloys, as is shown in Figure. 1.2.6. The {101�1}- {101�2} secondary twinning modes describes the {101�2} tensile twinning within the primary {101�1}

compression twins. This secondary twinning mode is often observed in hot rolled Mg alloy sheets, although secondary twinning has many twinning variants.

Table. 1.2.2 Different twinning systems in Mg alloys.

Twin type K1 K2 η1 η2 S

Tension twin {101�2} {1�012} <101�1�> <1�011�> -0.129

Compression twin {101�1} {101�3�} <101�2�> <303�2> 0.138

Secondary twin (101�1) [1�012]

Figure. 1.2.6 Schematic illustrations of different twinning modes in Mg alloys: (a) tension twins, (b) compression twins and (c) secondary twins. The twinning planes are marked with the transparent frame in

(a) and (b). Black lines and blue lines indicate HAGBs and basal planes in (c).

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Twin nucleation and growth are related with the glide of twin dislocations on twinning planes. Wang [39]

applied atomic simulations to study the (101�2) twinning nucleation. He proposed that nucleation mechanisms require simultaneous nucleation of multiple twinning dislocations. Capolungo [40] investigated the dissociation of dislocations, and he assumed that the pile-up dislocations can stimulate the dissociation of leading dislocations. High stress level produces high dislocation density, and of course high potential for nucleation. The high stress and density of defect structures in grain boundaries regions facilitate twin nucleation [41]. In the classic twin growth mechanism, twin growth is achieved by continuous generation of twinning dislocation loops in the twin planes [42], which may not be the case for Mg alloys. The work by Serra suggested that the interaction of lattice dislocations and twin interface plays an important role in twin growth [43]. Wang [44] investigated the dislocation slip-induced twin growth with in-situ tension experiments. He proposed that twinning disconnections from dislocation transformation at twin boundaries promotes twin boundaries migration.

Twinning has a significant effect on the yielding asymmetry of Mg alloys. As mentioned above, the activation of tension or compression twinning depends on the stress component. The conventional processing methods like rolling and extrusion result in a strong texture. Researchers used to attribute the yield asymmetry to the impact of texture. However, the work by Agnew showed that this yield asymmetry also existed even in the random textured sample [28].

In terms of strain hardening, twinning can increase the strain hardening rate in Mg alloys. This firstly comes from its effect on the microstructure. Profuse twins can divide the large parent grain into several small grains [45]. A high strain hardening rate is expected because of the Hall-Petch strengthening. Twin boundaries act as barriers for dislocation movement. Moreover, tensile twinning can reorient the grain to a hard orientation for basal slip in Mg alloy. The neutron diffraction work by Brown [46] revealed that the internal stress increased rapidly in the tension twinned grains due to their hard orientation. On the contrary, secondary twinning was also reported to produce a favorable orientation for slip, leading to work softening [47].

Twinning as a polar deformation mode has a close relationship to the texture. For example, the extrude Mg alloy with a fiber texture of <112�0> parallel to the extrusion direction (ED). When the material is compressed along ED, it has a tendency of tensile twinning because the a-axis undergoes compression stress.

Schmid factor (SF) is a good tool to evaluate the orientation impact on twinning behavior. The concept of SF will be given in the end of this section. Twinning modes with high SF values are usually preferred at low stress levels. However, some researchers have also reported the non-Schmid factor twinning phenomenon [48]. This can be explained by the local stress that deviates from the applied stress. Jonas [49] reported that the activated low SF twin variants can help accommodate the local strain. Guan [50] attributed the basal slip

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in adjoining grains can stimulate the low SF twin variant. However, although low SF twins can form due to the local stress, the high SF twins variants are more likely to grow than low SF ones [48].

Apart from the texture effect, the impact of CRSS on twinning should be taken into consideration. Generally speaking, the CRSS for tensile twinning is slightly higher than that of basal slip, but much lower than prismatic and pyramidal slip at room temperature [51]. As is known that basal slip with <a> Burgers vector does not have a component along c-axis, tensile twinning with a relatively low CRSS is another important deformation mode to accommodate the strain. The yield stress of single crystal with an ideal orientation for slip/twinning (with a corresponding SF of 0.5) is regarded as the slip/twinning CRSS. It is reported that the CRSS ratio of tensile twinning to basal slip in pure Mg at room temperature falls in the range of 1:2-4 [51- 53]. The measurement of twinning CRSS in polycrystalline material is more complicated. Gharghouri [54]

regarded the stress when texture change was firstly detected as the twinning CRSS during in-situ neutron diffraction experiments. Brown [46] reported a CRSS for tensile twinning in AZ31 ranging from 25 to 35 MPa, in good agreement with previous studies [26, 55-57]. The CRSS for compression twinning is higher compared to tensile twinning. The CRSS for compression twinning in Mg single crystal at room temperature was reported to be 110 MPa by Yoshinaga [58], in the range of 70-150 MPa by Wonsiewicz [52]. Barnett [59] estimated the compression twinning CRSS in AZ31 to be 110 to 125 MPa.

In the solid solution Mg alloys, the solution atoms introduce misfit into the lattice. The misfit can impede the dislocation slide and twinning. Kelley and Hosford [51] found the addition of Li or Th hardened the normal stress for twinning in the grains with favorable orientation. Nie [60] investigated the periodic segregation of Gd and Zn in twin boundaries of Mg alloys using transmission electron microscopy (TEM) and density functional theory computations. He proposed that the Gd atoms segregating on tension twin boundaries during annealing strengthened the materials.

Precipitates can also influence the twinning behaviors of Mg alloys during deformation. Clark [61] proposed that rod-shaped precipitates parallel to the c-axis in Mg-Zn alloy have a minor effect on tensile twinning while they can obviously pin the dislocations. Stanford [62] reported that the fine and uniformly distributed precipitates could promote twin nucleation but hinder twin growth. Gharghouri [63] studied the effect of precipitates on twinning in Mg-Al alloy. He found that twins can be held up by particles, engulf or bypass particles. Considering the high yield stress of the particles, he concluded that the engulfed particles were inclined rather than sheared by the twins. Robson [64] suggested the non-shearable precipitates produced a back stresses onto the twins which was relaxed by slip in the twinned volume.

A high strain rate slightly facilitates the formation of twins in Mg alloys compared to its obvious promotion to dislocation slip[31]. This can be explained by the insensitivity of twinning dislocation to the obstacles [65]. The CRSS for tensile twinning is quite stable in the temperature range from room temperature to

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300 °C [53], while that of compression twinning decreased with increasing the temperature [52].

Considering the fact that few twins are observed in AZ31 after high temperature deformation, the softening at elevated temperature is mainly attributed to the decrease of non-basal slip CRSS other than the change of twinning CRSS [66]. It is generally accepted that large grain size is beneficial to twinning. The dominant deformation mode changes from slip to twinning with increasing the grain size and decreasing the temperature, as shown in Figure. 1.2.7 [31]. Meyers suggested that larger grain size resulted in a lower CRSS for twinning based on Hall-Petch relationship [67].

Figure. 1.2.7 Effect of grain size on σ0.002 and σ0.2 of AZ31 alloy in compression test at 150 °C [31].

Schmid factor

It is generally accepted that the HCP structure is accountable for the deformation anisotropy. That is the geometrical relationship of grain orientation with the applied stress plays an important role in the mechanical response. Schmid factor (SF) is often used as an indicator to explain whether the deformation modes are easy to be activated. The SF is expressed as [68]:

m = cos(φ) cos(λ) (Eq. 1-1)

where φ and λ are the angles between the applied stress and the vector normal of glide plane and glide direction, respectively. The SF ranges from 0 to 0.5. A max value of 0.5 can be reached when φ = λ = 45°, indicating an ideal orientation for the deformation mode.

The SF law works well with uniaxial stress state, like uniaxial tension and compression. The normal deformation process like rolling or extrusion has a more complicated stress state, and thus a stress tensor is used to calculate the SF [69]. There are some differences in the SF calculation of dislocation slip and twinning. Dislocation glide is bidirectional while twinning is polar. Take the (0001)<112�0> basal slip as an

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example, basal dislocations can glide on both [112�0] and [1�1�20] directions. The tensile twinning can occur with the (101�2)[1�011] mode but not with the (101�2)[101�1�] mode because tensile twinning can only occur when tension stress is applied on the c-axis (or compression stress on a-axis). As a result, the SF value of dislocation glide is always positive while that of twinning can be negative. The low and negative twinning were reported because of local stress deviating from the applied stress.

1.2.4 Effect of RE on the deformation texture

Typical deformation textures can be formed in HCP structure metals during thermomechanical process, e.g., the basal texture from rolling and fiber texture from extrusion. This phenomenon can be explained by the active deformation modes. Take the rolling process as an example, the grains tend to rotate the c-axis parallel to ND by basal slip and tensile twinning, leading to a basal type texture. The basal texture restricts the application of wrought Mg alloy. The room temperature formability is poor because the easily active deformation modes, basal slip and tensile twinning, are not favored with this texture. The addition of RE into Mg alloys changes the active deformation modes, leading to a weak deformation texture. Agnew [28]

applied the Visco-plasitic Self Consistant (VPSC) simulation and attributed the RD split component to the increased activity of non-basal slip in Mg-Y alloy. Sandlöbes [70] proposed that Y can decrease the basal stacking fault energy, facilitating the non-basal slip. Besides, Y addition promotes the formation of compression and secondary twins. The diverse active deformation modes produce a weak texture. Li [71, 72] studied the grain interaction model including twinning of AZ31 and ME20 alloys during plain strain compression. He pointed that the pyramidal <a> slip in ME20 alloy has an advantage over prismatic slip and plays an important role in the formation of RE texture. Barnett [73] reported Ce containing Mg alloy show a more homogeneous deformed microstructure and weaker basal texture than pure Mg. The ductility is enhanced in Mg-Ce alloy owing to the smaller frequency of shear bands where the cracks nucleate. The texture weakening effect of RE addition can be magnified by the extra alloying elements like Zn and Zr.

The RE elements usually have larger atomic radius and thus introduce positive misfit when they are soluted into the matrix. The non-RE elements like Zn have smaller atomic radius which can compensate the misfit of RE by co-segregations. The co-segregations act as effective obstacles to dislocation motion, leading to a RE texture.

1.3 Recrystallization mechanisms of Mg alloys

Recrystallization is a process that the newly formed defect-free grains replace the deformed grains. The material in deformed state contains a high density of defects like dislocations and interfaces. It has a higher free energy than the perfect crystal state, and thus tends to reduce the energy by recrystallization. However, this process is too slow to be observed at room temperature. When the material is subjected to a high

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temperature, the atoms are more thermally activated. Thus, recrystallization can easily occur to eliminate the defects.

Recrystallization annealing is a widely used heat treatment method. In the early stage of annealing, the annihilation of point defects and rearrangement of dislocations take place without obvious movement of grain boundaries. This process is called recovery during which some physical properties of the material can be partially restored. For example, the electrical conductivity can be improved because the density of defects decreases. Recovery can release the internal residual stress, beneficial to the stress corrosion properties.

Dynamic recovery during deformation has a significant effect on the creep behavior [74]. The general concept of recovery can be explained in the following steps. Dislocations in the deformed condition exist as tangles. At elevated temperature, the tangled dislocation can climb in company with the annihilation of opposite sign dislocations. The movement of dislocations forms the dislocation cells that continuously absorb the neighboring dislocations. Then the dislocation cells transfer into the subgrain structure by the rearrangement of dislocations.

1.3.1 Definition of recrystallization

The properties can only be partially restored since the deformed structure and majority of dislocations still remain after recovery. Recrystallization can fully restore the properties by replacing the deformed grains with the newly formed defect-free grains. Recrystallization involves the nucleation and grain growth process.

The recrystallization, which occurs in a heterogeneous manner that distinct stages of nucleation and growth can be observed, is called discontinuous recrystallization. On the other hand, continuous recrystallization involves the simultaneous activities of nucleation and growth. Jazaeri [75] investigated the effect of parameters on the transition from discontinuous to continuous recrystallization in Al alloys. He pointed that continuous recrystallization is favored with smaller grain size, larger precipitates and strains.

Depending on whether the recrystallization occurs during deformation, recrystallization process is classified into dynamic recrystallization (DRX) and static recrystallization (SRX), respectively. DRX normally occurs at elevated temperature and large strains and is often observed in extrusion process. This thesis mostly focuses on the rolled Mg sheets in which DRX is not the dominant recrystallization mechanism. A basic background of DRX is introduced here. More detailed information can be found in the reference [76]. The thermomechanical processing parameters play an important role in the microstructure and texture evolution of DRX. The ductility at elevated temperature is enhanced because of the activated non-basal slip and softening mechanisms like dynamic recovery. There are three main DRX mechanisms in Mg alloys [77] : continuous DRX (CDRX) [76, 78], discontinuous DRX (DDRX) [79] and DRX mechanism associated with twinning (TDRX) [80]. The schematic of different DRX mechanisms is shown in Figure. 1.3.1 [81].The CDRX (route B1 to B4) is regarded as a strong recovery process. The dislocation plie-ups at grain

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boundaries form the dislocation arrays (LAGBs) which later transferred into HAGBs. On the contrary, DDRX starts with the process of bulging interiors by basal slip. This bulged area is then cut from the parent grain by the LAGBs that come from the operation of non-basal. After that, it evolved into recrystallization nuclei by absorbing mobile dislocations. TDRX is more related to the nucleation within the twins or twin intersections. The nucleation can be triggered in primary twins, twin and twin intersections, twin and grain boundaries intersections. The twin boundaries act as barriers for dislocation slip. A higher dislocation density in the twin boundaries region would facilitate the nucleation process. The necklace structure is the most common feature of TDRX.

Figure. 1.3.1 Schematic of different DRX mechanisms: TDRX (A1-A4), CDRX (B1-B4), DDRX (C1-C4) [81].

Compared with DRX, SRX is more important in the recrystallization of Mg alloy sheets. SRX is the post- deformation annealing process. In the thermomechanical process, SRX is wildly used to restore the ductility of materials, ensuring that it can undergo further deformation without cracking. For example, the intermediate annealing between rolling passes can release the residual stress by SRX. The alloying elements play an important role in the kinetics of SRX. On one hand, it affects the active deformation modes and the stored deformation energy. Stanford [82] proposed a higher stored energy in the deformed grains of Mg-Gd alloy than AZ31 alloy. On the other hand, the solution atoms and precipitates hinder the migration of grain boundaries during grain growth [76].

Effect of SRX on texture evolution in Mg alloys

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SRX is helpful to optimize the microstructure and texture. Unlike Al and Cu alloys in which recrystallization will significantly alter the deformed texture [83, 84], It is reported that the basal texture in pure Mg or commercial Mg alloys like AZ31 will be maintained or even strengthened during SRX [85]. Steiner [86]

attributed the basal texture to the grain size advantage of basal oriented grains over other grains during annealing. Bhattacharyya [87] studied the boundary energy and pointed that the low energy and mobility of LAGBs was responsible for the basal texture. This sharp basal texture is harmful to the ductility due to the strong anisotropy of Mg alloys because the normal deformation modes at room temperature, basal slip and twinning, are not favored with the basal texture.

RE addition is reported to weaken the basal texture and improve the formability of Mg alloys [88]. Bohlen [16] reported the addition of Y into Mg-Zn based alloys can develop TD spread texture component, reduce the anisotropy and improve the elongation. Hantzsche [89] proposed that the texture weakening mechanism of RE requires sufficient chemistry content. The weakening mechanisms of RE addition, in terms of deformation aspect, are explained by the reduced stacking fault energy [70], the decreased CRSS ratio of non-basal to basal slip [25], suppression of shear bands [73]. The details of RE addition on the deformation texture can be found in the deformation section. RE addition can also influence the recrystallization texture.

Several mechanisms are proposed regarding the effect of RE on texture evolution during annealing, like particle stimulated nucleation (PSN) [88, 90, 91], shear bands induced nucleation (SBIN) [92, 93], and deformation twin induced nucleation (DTIN).

1.3.2 Nucleation in Mg alloys

Burke and Turnbull [94] used to explain the nucleation of recrystallization with the phase transformations nucleation theory that nucleation is the result of random atomic fluctuations. However, this theory is not so convincing because the required energy to form the LAGBs is very high [95] while the driving force is low.

It is generally accepted that the nuclei come from the small volumes in the deformed condition. Some nucleation mechanisms of recrystallization are termed as strain-induced grain boundaries migration (SIBM) [96, 97], LAGBs migration [98, 99] and subgrains coalescence [100, 101]. In SIBM mode, the different dislocation densities between neighboring grains are the driving force of grain boundaries motion. As is shown in Figure. 1.3.2 (a), the thick black line separates the less deformed grain, grain A, and more deformed grains, grain B. During annealing, the grain boundaries can bow towards the interior of grains B due to the tendency of reducing the energy [102]. LAGBs migration mode describes the growth of subgrain by absorbing the neighboring LAGBs in Figure. 1.3.2 (b). Subgrain coalescence mechanism explains the formation of a large subgrain from small subgrain pair. As is shown in Figure. 1.3.2 (c), the rotation of small subgrains builds the coincide lattices, leading to the coalescence of subgrains [101].

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Figure. 1.3.2 Three different nucleation mechanisms: (a) SIBM, (b) LAGBs migration, (c) subgrains coalescence [103].

Effects of particles on the recrystallization process have been extensively investigated in recent years. Some theories are well established in Al alloys [104]. First, particles whether can be sheared or not will introduce the stress and strain field and deformation energy in the neighboring regions, which means the driving force for recrystallization is enhanced. Another advantage of the particles is that they provide extra surface for nucleation, decreasing the energy requirement of nucleation. However, the uniformly distributed particles with small spacing can pin the grain boundaries, which retard the grain growth process. There are some arguments when we correlate the weakened texture in Mg-RE alloys to PSN mechanisms. Robson [105]

reported a higher misorientation gradient in the vicinity of particles than the matrix away from particles, resulting in randomly oriented recrystallization nuclei, as is shown in Figure. 1.3.3. Talal [106] proposed the growth advantage of these PSN grains over other grains is responsible for the weakened texture.

However, Stanford [107] reported the texture was weakened in ME10 rather than AZ31 despite the fact that both alloys have comparable size and distribution of the particles. Besides, the dilute Mg-Zn-RE alloys with few precipitates exhibited a weak texture [16]. Thus, PSN may not be the major mechanisms for texture weakening.

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Figure. 1.3.3 EBSD maps of: (a) IPF map of M03 alloy in the region of particles, (b) corresponding line misorientation plot from (a), (c) recrystallized grains around particles in M16 alloy, (d) scatter orientation

of recrystallized grains in (0001) pole figure [105].

Shear bands are the severely deformed band structure with large internal misorientation. The origin of shear bands is not clear. Couling [108] suggested that the shear bands in Mg-RE alloys came from double twinning.

Sandlöbes [70] pointed that shear bands consist of narrow secondary twins and lamellar matrix. The lattice continuously rotated around <112�0> axis in the shear bands and the misorientation of shear bands to the surrounding matrix can reach 15°. The addition of Y modified the distribution of shear bands and decreased the potential of failure within shear bands. Stanford [92] found the recrystallized grains showed a random texture within shear bands while those formed at grain boundaries maintained the deformed texture. Basu [109] reported the grain growth advantage of non-basal grains within shear bands over basal grains. Guan [110] studied the shear bands related recrystallization during quasi in-situ annealing of WE43 alloy. He proposed that the basal texture was weakened rather than replaced by the RE texture. The scattered weak texture of recrystallization nuclei was preserved during grain growth. But there are some doubts regarding the explanation of texture weakening mechanisms of RE addition by SBIN. For example, the shear bands were observed in many commercial Mg alloys like AZ31 [73] and Mg-1Zn [93], which exhibited strong basal textures after rolling. What’s more, the texture weakening effect of RE addition was also reported even without the formation of shear band by many researchers [17, 111].

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Figure. 1.3.4 EBSD maps of Mg-Gd extruded at 415 °C: (a) full ebsd map, (b) recrystallized grains at grain boundaries and corresponding inverse pole figure, (c) recrystallized grains within shear bands and

corresponding inverse pole figure [92].

Apart from the significant impacts on the mechanical properties, deformation twins are also the favored sites for recrystallization. It should be noted that recrystallization in Mg alloys normally triggers within compression [112] and secondary twins [113] rather than tension twins. Unlike the stable compression or secondary twin boundaries, tension twin boundaries with a high mobility can release the strain accumulation during twin propagation [66, 114]. As a result, the driving force for recrystallization is much higher in compression and secondary twins than in tension twins. Li [112] and Martin [115] suggested that even though the orientations of nuclei within twins deviated from that of the matrix, their contribution to the recrystallization texture is minor due to the low volume fraction of the twins. Besides, the twin boundaries restricted the growth of nuclei into the matrix. Guan [113, 116] traced the whole recrystallization nucleation and grain growth of WE43 alloys during quasi in-situ annealing. He reported DTIN was the dominant recrystallization mechanism as the recrystallized grains originating from secondary twins account for a volume fraction of 69.9% in the fully recrystallized condition. The RE texture formed during nucleation was maintained afterwards while dynamic precipitation at twin and grain boundaries restricted the formation of basal oriented grains by SIBM.

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Figure. 1.3.5 EBSD maps of WE43 alloy during quasi in-situ annealing: (a-e) tension twins (f-k) primary tension twins and compression-tension twins, (l-q) compression-tension twins [116].

1.3.3 Grain Growth in Mg alloys

The grain growth is termed as the process that recrystallized nuclei grow at the expense of deformed matrix or other nuclei. Stored deformation energy and grain boundaries curvature are the driving force for the grain growth process [76]. Compared to the deformed matrix with a higher dislocation density, recrystallized nuclei are defect-free. They tend to grow into the deformed matrix to decrease the system energy, which is also the case for the grain boundaries with sharp curvature. Some researchers proposed that the grain growth had a greater effect on the texture evolution than the nucleation process. The preferential grain growth mechanism was applied to explain the texture weakening mechanism of Mg-RE alloys.

Some specific grain boundaries with a high mobility are accountable for the preferential grain growth. In the rolled Mg sheets, basal texture was remained after fully recrystallization [117]. A special feature of this recrystallization process is the 30° <0001> misorientation between the recrystallized grain and the deformed matrix. This phenomenon was explained by the high mobility of these specific grain boundaries [118]. The basal texture could be strengthened during the growth of basal oriented nuclei with 30° <0001>

misorientation grain boundaries because their c-axis was maintained parallel to ND. The RE addition can

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hinder the formation of this special misorientation grain boundaries [119], contributing to a weakened texture in Mg-RE alloys. Barrett [120] calculated the energy distribution of c-axis tilt boundaries for Mg-Y and pure Mg. He found that the introduction of Y would reduce the energy gap of certain misorientation grain boundaries, thus stabilizing nuclei with a wide variety of orientations and giving them nearly equal opportunity for growth. The difference in the deformation energy can also promote preferential grain growth.

Imandoust [121] reported the recrystallized grains with high SF value underwent more deformation and higher energy were consumed by low SF grains, contributing to the formation of RE texture.

Figure. 1.3.6 Preferential grain growth mechanisms: (a) relationship between grain energy and tilt angle of adjacent grains [120], (b) IPF maps of Mg-Gd alloy annealed for 3 h at 450 °C, (c) corresponding SF for

basal slip[121].

1.4 Objective of the study

This study mainly focuses on the deformation and recrystallization mechanisms of Mg-RE alloys. Our objective is to understand the origin and development of RE texture which are important to develop high performance Mg alloys.

The working alloy is Mg-Nd based alloys. Nd is the chosen RE element with moderate solubility in Mg.

Compared to the well-studied RE element Y, Nd had a more pronounced texture weakening effect even at a very low concentration of 0.04 at %. The Mg3Nd precipitates on the prismatic plane increase the CRSS for basal slip [122]. This enhances the activities of non-basal slip, contributing to a weak deformation texture.

The precipitates and solute Nd atoms can modify the recrystallization texture by pinning the grain boundaries. Apart from the Nd, Zn and Al are also introduced because they can magnify the texture weakening effect of RE. Zn and Al are the most common alloying elements in Mg alloys. Both of them have smaller atomic radius than Mg and Nd. As a result, they can decrease the lattice misfit of Nd addition.

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Besides, Al addition increases the strength of Mg alloys due to its strong affinity with Mg atoms. Zn decreases the stacking fault energy that promotes the twinning activities.

Advanced techniques like quasi in-situ tensile and annealing, transmission Kikuchi diffraction (TKD) are applied to reveal the underlying deformation and recrystallization mechanisms of Mg-Nd based alloys. In previous studies, experiments with different samples at various strain or annealing time were used to explain the microstructure and texture evolution. The deviation of the samples is usually regarded to have a negligible effect on the result, which in fact is not the case. The quasi in-situ experiments have the advantage of tracing the entire process of deformation and recrystallization of the same samples, which are more convincing and easier to understand. TKD characterization is an ideal tool. It can be easily evaluated like EBSD, and it has a high resolution comparable to TEM. TKD samples are thin slices prepared by focused ion beam (FIB). Another good advantage of TKD over TEM is that we can prepare the slice exactly at the regions of interest by marking the region with a deposit layer.

The evaluation of the results is based on the EBSD analysis software and MTEX Matlab toolbox. The EBSD software like TSL OIM analysis (EDAX system) is a well-developed commercial software to evaluate the local texture and grain boundaries structure of polycrystalline materials. With OIM software, we can obtain the primary result like the pole figure, misorientation plot, twin boundaries fraction. However, more detailed analysis is not available in OIM software. For example, the SF calculation in OIM only shows the max value even though one slip system has many equivalent variants. The calculation of SF for twinning is not trustable because the software also regards twinning as bidirectional deformation mode which is not the case. The non-SF twinning cannot be evaluated and it is impossible to index different twin variants in OIM software.

MTEX is used to supplement these functions. It is very helpful to quantitatively evaluate the large data of quasi in-situ experiments with MTEX. The details of MTEX analysis will be presented in the experiment chapter.

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2 Experiment Procedure

2.1 Materials and sample preparation

The alloys used in this study were prepared by direct-chill casting with pure Mg, Nd and Al or Zn. The chemical compositions characterized by X-ray fluorescence method are shown in Table. 2.1.1. The Mg-Nd ternary alloys with the addition of Al or Zn are termed as AN12 and ZN12, respectively. The alloy melt was stirred at 720 °C under an atmosphere of Ar and 3 Vol.% SF6 for 30 min before pouring. After that, the melt was cast into a steel mold, which was pre-heated to 680 °C and coated with boron nitride. Water quenching of the melt was applied by sinking the filled mold slowly into a water bath. Because of the lower melting temperature with the alloying of Zn, homogenization treatment at 450 °C to the AN12 alloys and 420 °C to the ZN12 alloy were applied respectively, following the water quenching.

Table. 2.1.1 Chemical composition (wt.%) of the examined Mg alloys.

Alloy Nd Al Zn Mg

Mg-1.0Al-1.5Nd (AN12) 1.77 1.18 - Balanced

Mg-1.0Zn-1.5Nd (ZN12) 1.71 - 1.07 Balanced

The homogenized ingots were machined into slabs with the dimension of 100 (width) × 50 (length) × 10 mm (thickness). The slabs were hot rolled at 420 °C with a pass reduction degree of 10-15 %. The intermediate annealing of the rolled slabs was carried out at the rolling temperature for 10 min after each rolling step, to release the residual stress and avoid the cracking due to the temperature drop. The sheet was rolled to the final thickness of 2.3 mm after 11 passes with a total strain of 0.77.

2.2 Microstructure characterization

Samples for the microstructure characterization were cut from the rolled sheets with a band saw. The cut samples were mounted with a cold cure embedding resin. The samples were then ground with emery papers from 800 to 2500 grit. After that, they were mechanically polished with oxide polishing suspension and etched with a picric acid-based etchant [123]. The microstructure of the samples was observed on the Leica DM4 P metallurgical microscope. The observation plane was set to the ND-TD plane because it exhibits the material flow along the main deformation axis, such as the rolling direction (RD) and thickness direction (ND).

The sample preparation for EBSD measurement is very similar to that for the optical microscopy. After mechanical polishing, electrolytic polishing was used to remove the tiny scratches and deformation layer

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including deformation twins occurred by mechanical polishing. Electrolytic polishing was conducted by using a commercially available AC2 electrolyte and Lectropol-5 machine at -20 °C and 30 V. The EBSD measurements were carried out on a field emission gun scanning electron microscope (FEG-SEM), Zeiss Ultra 55, working at 15 kV. The step size of EBSD measurement was 0.4 μm. The TSL OIM analysis 7 was used to evaluate the EBSD data.

2.3 Quasi in-situ EBSD measurements during tension and annealing

Quasi in-situ experiments were carried out by EBSD measurements repeatedly taken at the same sample position, before and after each interruption of experiments, like annealing and tensile loading. Quasi in-situ experiments are very useful to study the deformation and recrystallization process because they allow to trace the microstructural development without considering the inhomogeneity of the samples. The details of quasi in-situ experiments are given below.

The samples for the quasi in-situ measurements were prepared from the as-rolled sheets of the examined alloys by using the method described in the above section. An EBSD measurement was conducted at the central area of the sample, termed as the as-rolled condition. After EBSD measurement, the sample fixed on the sample holder was placed on a pre-heated stage at the annealing temperature. They were inserted into the furnace under the Ar protection atmosphere for the annealing process. After the annealing for the targeting time, the samples were immediately quenched in deionized water. It should be mentioned that only the sample holder was merged into the water to cool down the sample, the measurement surface was not contaminated by water. To remove the thin surface oxide layer, the sample was cleaned with citric acid solution for 1-2 seconds. Since the cleaning time was short, its impact on the thickness reduction is negligible. The sample was then put back into SEM for the second EBSD measurement. Same operations were repeated for the third or fourth EBSD measurements. Thus, the microstructure and texture evolutions during annealing could be successfully traced at the target area.

Figure. 2.3.1 (a) Zeiss Ultra 55 FEG-SEM, (b) the setup of quasi in-situ annealing.

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The quasi in-situ tensile tests were performed on the fully recrystallized samples to exclude the effect of twins and residual stress. Dog-bone samples with a gauge length of 18 mm parallel to the RD were machined from the fully recrystallized sheets. The samples were ground, mechanical and electrolytic polished. The first EBSD measurement was conducted to record the texture and microstructure of the initial tensile sample.

The sample was then fixed on the Kammrath-Weiss universal testing module (5KN) for the uniaxial tensile test. The constant strain rate of 10^-3 s^-1 was used in this study. The tensile test was stopped at true strains of ε = 0.02, 0.05 and 0.08. Besides the EBSD measurements, SEM images were taken on the sample at different strains to obtain the microstructure development and slip trace analysis.

Figure. 2.3.2 (a) Kammrath-Weiss tensile machine, (b) dimensions of the tensile sample for the quasi in- situ tensile tests.

Apart from the EBSD data at different strains, another important result of the quasi in-situ tensile experiment is direct evidence of the dislocation movements, namely slip trace analysis. In Mg alloys, basal slip with the lowest CRSS value is regarded as the dominant deformation mode. The understanding of dislocation slip will help explain the mechanical response during deformation. The activity of dislocation slip is normally evaluated with TEM, which enables direct imaging of the dislocation lines. However, the small volume of transparent area to the electron beam and poor statistics make it difficult to explain the macro deformation behavior. Numerical methods, like Visco-plastic Self Consistent (VPSC) model and simulation, are also used to predict the activities of different slip systems and twinning modes by fitting the stress-strain curve and texture evolution. The weakness of simulation is no direct evidence of the dislocations. Besides, it is not easy to get a metal-physical explanation of deformation mechanisms since many sets of input parameters can be fitted well to the texture evolution and stress-strain curve. EBSD assisted slip trace analysis, combining good statistical and direct evidence of dislocation slip and twinning, is adequate to study the deformation mechanisms.

The slip traces, e.g., on the basal plane, is schematically shown in Figure. 2.3.3 (a). Numerous dislocations glide on the slip plane and, as a result, a small step of displacement will appear at the intersection of slip

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plane and sample surface. This step is weak and usually comes in a straight line shape in the SEM images.

An example of the slip trace in the deformed AN12 alloys is shown in Figure. 2.3.3 (b). Based on the theoretical slip plane trace calculated from the grain orientation, Figure. 2.3.3 (c), the slip lines appeared in Figure 2.3.3 (b) are indexed as the basal slip plane trace (black line). Thus, it can be concluded that basal slip is the dominant deformation mode in this grain.

Figure. 2.3.3 (a) schematic of slip trace, (b) observed basal slip plane trace, (c) theoretical slip traces of different slip systems of the grain in (b).

2.4 Transmission Kikuchi diffraction

Transmission Kikuchi diffraction (TKD) is a powerful technique to study metallic materials at the nano- scale. It provides a higher spatial resolution than conventional EBSD measurement because the interaction volumes of TKD is smaller and more sensitive, as is shown in Figure.2.4.1 [124]. The signal is formed by collecting Kikuchi-scattered transmitted electrons in TKD while conventional EBSD uses the backscattered electrons.

The preparation of a thin slice TKD sample includes several steps, as is shown in Figure. 2.4.2. First, a Pt layer with a dimension of 25 × 1 µm with the thickness of approximately 1.5 μm was deposited on the electrolytic polished surface. It is used to mark the region of interest, and simultaneously to protect from the ion milling. The gallium ion beam was used to remove the materials adjacent to the marked area. After that, the thin film was cut and welded on the micromanipulator. It was then attached on the TEM grid using platinum deposition and the weld to the micromanipulator is cut loose.

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Figure. 2.4.1 Comparison of signal collection between conventional EBSD and TKD [124].

Figure. 2.4.2 Preparation route of TKD samples: (a) Pt deposition, (b) bulk-out, (c) U-cut, (d) lift out, (e) mounting, (f) thinning.

Deformation and recrystallization mechanisms of Mg-Al-Nd and Mg-Zn-Nd alloys