Multi-dimensional hydride characterization and accommodation behavior in pure titanium

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Solute effect on plastic deformation

There are two types of interstitial sites for solute atoms in HCP structure, octahedral (O) and tetrahedral (T) sites, as shown in Fig. 1.7. Based on the solubility and lattice parameter change, Ehrlich et al. [35] suggested that oxygen atoms occupy O sites in titanium. Due to the limit of solubility, only one half of these sites are occupied. Inelastic neutron scattering measurements have
been used to determine the hydrogen site occupied in α-Zr and α-Ti, the result shows that hydrogens are located in T sites [36]. With the framework of the density functional theory, H occupancy is confirmed in T interstitial in Zr at low temperature [37]. Although the O site is unfavorable for H atoms, it also plays an important role for the H atom diffusion. Both jumps between adjacent T and O sites in the basal plane and between O sites along the c-axis were observed. Besides, the H atoms can share the bonds of Zr atom and lead to a local softening of the bond between these Zr atoms and the surrounding bulk, which may change the cohesion property of Zr.

Hydrogen damage mechanism

Hydrogen enhanced decohesion mechanism is a simple mechanism. When hydrogen atoms are available inside the material, the internal stress gradient drives the diffusion of hydrogen atoms towards crack tips and results in the decrease of interatomic strength or cohesive strength [47]. Simultaneously, the surface energy is reduced by the decreased cohesive strength and so that fracture stress is also reduced and cleavage like fracture occurs below its permissible value. The hydrogen accumulation near crack tips can also decrease the resistance for dislocation motion [81]. The local yield stress drop accelerates the dislocation movement which corresponds to the slip bands at the crack tips on the fracture surface [82]. The increase of dislocation mobility also leads to the local plastic deformation inside embrittled material. By fractography examination, there appears more local plastic deformation with the reduction of macroscopic ductility, but not direct link is observed between hydrogen enhanced local plasticity and actual embrittlement.
Micro-void coalescence (MVC) is inherently a ductile fracture mechanism [83,84]. The MVC crack propagation happens in various stages such as void nucleation, void growth, void coalescence and crack extension. The MVC crack growth happens in a zig-zag pattern by joining of voids in crack propagation direction. MVC dimple produced by the effect of hydrogen possesses poor ductility and final fracture occurs due to shear stress. This is called hydrogen assisted micro-void coalescence. Besides, the previous research on Zr alloy [85] with various hydrogen contents proposed that the increase in void density with strain is enhanced by the presence of hydrogen, and the ductility decreases with increasing hydrogen content. The increased number of void nucleation sites is due to the formation of hydride precipitates, which increase the void nucleation kinetics [86].
In some materials with low solubility of hydrogen, such as titanium and zirconium, formation of hydride usually takes places once the local hydrogen solubility limit is exceeded. The hydride is a brittle phase, which actually causes the embrittlement of the material. Hydrogen agglomeration at crack tips with high stress intensities also results in the hydrogen-embrittlement behaviors due to delayed hydride cracking, which is a subcritical crack growth mechanism and crack propagates in a discontinuous mode [80,87–89]. A crack growth cycle includes stress-directed diffusion of hydrogen towards the crack-tip, hydride formation, which subsequently induces cleavage in front of a growing crack. Regardless of whether the hydride formation is driven by high internal hydrogen concentration or by stress-gradient driven high hydrogen flux at a crack tip, it leads to detrimental effects on mechanical properties, such as loss of ductility and sustained load cracking. Thus, it is significant to develop a more comprehensive understanding of room temperature hydride formation, and of its influence on the damage evolution mechanisms.

Hydride phase transformation

Hydrogen atoms in solid solution state can diffuse very fast in Ti and Zr alloys towards the areas of local stress concentrations, such as crystal defects, precipitates, grain and phase boundaries. Thus, these areas are the preferential nucleation position for hydride precipitations in order to reduce the activation energy. When hydrogen concentration surpasses the terminal solubility limit, hydride nucleation happens and then the hydrogen atoms will be bonded in the interstitial sites of the crystal lattice. The Ti-H and Zr-H binary phase diagram is shown in Fig. 1.10 [90]. At 300 ºC, the hydrogen solubility is around 0.15 wt. % in α-Ti and 0.07 wt. % in α-Zr, which rapidly decreases as the temperature drops. Thus, hydrides are extremely easy to precipitate at room temperature (especially for Zr) due to the low hydrogen solubility.

Hydride phases and orientation relationships

Various hydride phases were found in titanium at room temperature depending on hydrogen concentration, metastable γ-hydride (TiH, face-centered tetragonal (FCT), c/a > 1), stable δ-hydride (TiHx, 1.5 < x < 1.99, face-centered cubic (FCC)) and ε-hydride (TiH2, FCT, c/a < 1)) [91– 95]. The δ-phase hydride is the most commonly observed hydride phase formed by multi-step process. Bair et al. [96] proposed that the shape evolution of the δ-hydride is highly dependent on the intermediate ζ and γ phase, which was explored by a multiphase field model. The γ, δ and ε phase hydrides also happen in Zr alloys [97,98], while the fourth hydride phase ξ-Zr2H, a metastable hexagonal phase, was first observed in [99] with a trigonal symmetry structure fully coherent with HCP Zr matrix.

Hydride enhanced hardening process

With the hardness increase by the generation of hydride precipitation, the material failure occurs easier due to the decrease of plastic and ductile deformation. The origin of such failure is attributed to the brittle nature of the hydride, micro-crack or void nucleation by the interaction between hydride and dislocation at the hydride interface [112]. The internal stress produced by strain incompatibility between hydride and matrix material is the reason for the hydride enhanced hardening process. The hardening is dependent on the crystal orientation and the relationship between slip plane and hydride habit plane; it has been shown that the hardening is more significant when the slip plane is close to the habit plane of hydride [113]. Sometimes, the effect of hardening is weakened because hydrides can also experience plastic yielding under certain working conditions, which was confirmed by Guillot et al. [114] in α-titanium using TEM technique. After plastic strain, the cross slip and dislocation loops debris were observed inside hydride. Thus, hydrides cannot always be regarded as a hard phase in a soft matrix. By the method of ab initio calculation, Udagawa et al. [112] suggested that {111} plane slip in FCC hydride shows less resistance to both detachment and slip than that in {001} and {110} planes.
Chen et al. [115] deeply studied the deformation behaviors of γ-hydride platelet in CP-Ti with 77 ppm hydrogen during cyclic process. Both OR1 and OR2 γ-hydrides were observed under scanning electron microscopy (SEM) and TEM techniques and OR1 hydrides are easier sheared by slip band, which is due to the interface structure and the relative orientation of hydride with matrix. The OR1 hydrides were deformed by the crossed slip through the coherent interface (Fig. 1.12). The equation for slip crossing process can be written as: 1 [112̅0] ̅ ↔ 1 [11̅0] + br (1-4) 3 2 (110)γ (1010)α.


Crystallographic calculations

In the present work, there are two phases: hydride (cubic: FCC/FCT) and α-Ti (hexagonal: HCP), thus only FCC, FCT and HCP phase were involved in the crystallographic calculations. Before the discussion of coordinate transformation, the different types of coordinate systems used in HKL Channel 5 acquisition system should be introduced. To obtain the relative crystal orientation, the orientation image map uses sample surface to define the orthogonal Cartesian sample frame (X1-X2-X3), as shown in Fig. 2.6a. On the sample surface, each crystal has corresponding frame, that is Cartesian crystal coordinate frame (x1-x2-x3). A crystal orientation is defined as the rotation between crystal and sample coordinate frame. The crystal coordinate frames corresponding to the Bravais lattice coordinate systems of different crystal structures are presented in Fig. 2.6b and c. In cubic structure, the Bravais coordinate system is aδ-bδ-cδ (a=b=c and α=β=γ=90º) for FCC and aγ-bγ-cγ (a=b≠c and α=β=γ=90º) for FCT, while the relationship between crystal and Bravais system is x1//a, x2//b and x3//c. The Bravais coordinate system of HCP structure is a1-a2-a3-cα (a1=a2=a3≠cδ, α=β=90º and γ=120º), the relationship with crystal system is x2//a2, x3//cα and x1//(a2×cα).

Table of contents :

Chapter 1 Literature review
1.1 Hexagonal closed packed structure
1.1.1 Deformation modes Slip Twinning Kinking
1.1.2 Solute effect on plastic deformation
1.2 Hydrogen damage
1.2.1 Practical cases
1.2.2 Hydrogen damage mechanism
1.3 Hydride phase transformation
1.3.1 Hydride phases and orientation relationships
1.3.2 Misfit accommodation
1.3.3 Hydride enhanced hardening process
1.4 Chapter summary
Chapter 2 Experimental procedures and crystallographic calculations
2.1 Experimental procedures
2.1.1 Material
2.1.2 Hydrogen charging procedure
2.1.3 Polishing preparation
2.1.4 Nanoindentation test
2.1.5 Tensile tests
2.1.6 Microstructural characterization
2.2 Crystallographic calculations
2.2.1 Coordinate transformation Bravais ↔ crystal coordinate system Rotation of crystal coordinate system Crystal ↔ sample coordinate system
2.2.2 Misorientation (Disorientation)
2.2.3 Trace analysis
2.2.4 Deformation theory Basic concepts Twinning Hydriding
2.3 Chapter summary
Chapter 3 Crystallographic orientation dependence of hydride precipitation in commercial pure titanium
3.1 Introduction
3.2 Experimental process
3.3 Character of hydride layer
3.3.1 Microstructure characterization before and after hydrogen charging
3.3.2 Classification of initial α-Ti grains
3.3.3 Crystal orientation dependence of hydride precipitation
3.4 Orientation relationship preference of α-Ti / δ-hydride transition
3.5 Strain anisotropy of α-Ti / δ-hydride phase transformation
3.6 Anisotropy of hydride nucleation
3.6.1 Microstructure of hydride platelets
3.6.2 Variant selection of hydride platelets Grains with one or two preferential hydride variants (Group I) Grains with more than three hydride variants (Group II)
3.7 Chapter summary
Chapter 4 Multi-dimensional hydride characterization and accommodation behavior in pure titanium
4.1 Introduction
4.2 Experimental process
4.3 Microstructure evolution of diffusion surface
4.4 Hydride microstructure inside hydride layer
4.4.1 Near-matrix hydride nucleation (Region I) Grains with internal orientation change {101̅2} and {112̅2} twin induced by OR2 hydride
4.4.2 Near-surface hydride precipitation (Region II)
4.5 Cross section of hydride layer
4.6 Intergranular hydride
4.7 Discussion
4.7.1 Hydride transformation mechanism
4.7.2 Variant selection of {101̅2}, {112̅2} and {101̅1} twin
4.7.3 Formation mechanism of adjoining hydride pair
4.8 Chapter summary
Chapter 5 Hydride induced hardening in commercial pure titanium
5.1 Introduction
5.2 Experimental process
5.3 Nano-indentation test
5.3.1 Mechanical property of α-Ti at different hydrogen charging times
5.3.2 Anisotropic hardness of α-Ti before and after hydrogen charging Microstructure of grid indentation array Anisotropic hardness of α-Ti Anisotropic hardness of δ-hydride Comparison between α-Ti and δ-hydride
5.4 Tensile test
5.4.1 Tensile property influenced by hydrogen charging
5.4.2 Interaction between hydride and plastic deformation modes Hydride-dislocation interaction Hydride-twin interaction
5.5 Chapter summary
Chapter 6 Conclusions and prospects
6.1 Conclusions
6.2 Prospects


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