Experimental study of the microstructural evolution of chemical vapour deposited nickel upon annealing

Transcription

1 Experimental study of the microstructural evolution of chemical vapour deposited nickel upon annealing by Chichi Chen A thesis submitted in conformity with the requirements for the degree of Master of Applied Science Materials Science and Engineering University of Toronto Copyright by Chichi Chen 2011

2 Abstract Experimental study of the microstructural evolution of chemical vapour deposited (CVD) nickel upon annealing Chichi Chen Master of Applied Science Materials Science and Engineering University of Toronto 2011 The effect of annealing conditions on the microstructure evolution of CVD nickel was investigated systematically in the present study by differential scanning calorimetry, optical microscopy and transmission electron microscopy (TEM), upon both ex-situ and in-situ annealing. TEM observation revealed the as-deposited CVD nickel possessed a bi-modal grain structure, with large columnar grains embedded in nanocrystalline matrix. Ultrafine and nano growth twins were present as well as multiply twinned grains with five-fold symmetry. Microstructure observation upon annealing showed that grain growth did not occur until annealing at 400ºC. Detwinning was observed at 400ºC and higher temperatures. The ultrafine and nano twins tended to transform into dislocation cell structures and this phenomenon was driven by the excess free energy associated with the high density of grown-in twin boundaries. The five-fold twinned grains were found to be thermally stable up to 600ºC. The hardness was observed to decrease with increasing annealing temperature. ii

3 Acknowledgements I would like to take this opportunity to acknowledge all the people who have showed their support throughout this project. First and foremost, I would like to give many thanks to my supervisor Professor Zhirui Wang. This thesis would not have been possible without his expert professional guidance, patience and encouragement in all aspects of this project. Furthermore, my gratitude is extended to Mike Mei and Sal Boccia for their helpful suggestions and technical assistance with Electron Microscopy, Srebri Petrov and John Graydon for their professional analysis and suggestions during my experiment and all my colleagues Zongshu Li, Jean Hsu, Jun Huang for their helpful discussions and technical support. Finally, the material supply from Weber Manufacturing Technologies Inc. is highly appreciated. iii

4 Table of contents Contents Abstract... ii Acknowledgements... iii Table of contents... iv List of Tables... vi List of Figures... vii List of Appendices... xii 1 Introduction Nanocrystalline materials Definition Preparation of bulk nanocrystalline materials Mechanical properties of nanocrystalline materials Chemical vapor deposition Introduction Nickel Vapour Deposition Structure and grain morphology of CVD nickel Mechanical properties of CVD nickel Advantages of nickel vapor deposition over electroforming Thermal stability of nanocrystalline materials Thermal stability of cold worked materials Thermal stability of nanocrystalline nickel Twinning and detwinning in FCC materials Twinning and its effect on materials properties Stability of nanotwins in FCC materials Experimental Material Thermal stability analysis via DSC Microstructure evolution investigation Optical microscopy iv

5 2.3.2 Grain boundary tracking via semi in-situ optical microscopy Transmission electron microscopy (TEM) Hot stage In-situ TEM Microhardness and Rockwell hardness testing Results Thermal stability investigation using DSC Preliminary analysis using Optical Microscopy Characteristics of as-received CVD nickel Microstructure evolution upon annealing Grain boundary tracking by Semi in-situ OM Detailed analysis using Transmission Electron Microscopy Characteristics of as-deposited CVD nickel Microstructure evolution upon annealing Hot stage In-situ TEM Mechanical properties of CVD nickel upon annealing Discussions Grain growth behaviour of CVD nickel upon annealing Thermal stability of ultrafine and nano twins in CVD nickel Thermal stability of five-fold twinned structures Analysis on the stable structure in hot stage in-situ TEM Thermal stability of CVD nickel in terms of mechanical properties Conclusions Future Work Appendices References v

6 List of Tables Table 1.1 Properties summary of CVD nickel, conventional nickel and electrodeposited nickel..18 Table 1.2 Chemical composition of CVD nickel...18 Table 1.3 Mechanical properties of Ni and Ni-B alloys...20 vi

7 List of Figures Figure 1.1 Atomic structure of a two dimensional nanocrystalline material...2 Figure 1.2 Schematic illustrations of four types of nanocrystalline materials showing structure in different numbers of dimensions..3 Figure 1.3 Schematic representations of severe plastic deformation setup Figure 1.4 A schematic diagram of Hall-Petch relation and inverse Hall-Petch relation...8 Figure 1.5 Schematic representations of the CVD reactor. 10 Figure 1.6 Schematic of three types of structures obtained by CVD.14 Figure 1.7 Optical image showing typical grain structure of pure CVD nickel 15 Figure 1.8 Microstructure of a CVD pure nickel...15 Figure 1.9 (a) Five <011> subgrains meet to form a pentagon (b) zone axis diffraction pattern of the pentagon in (a)..16 Figure 1.10 (a) A regular decahedron consisting of five irregular tetrahedrons (b) Relationship between the fcc and bco structure..17 Figure 1.11 Micro-hardness as a function of boron content for as deposited and Ni and Ni-B alloys..19 Figure 1.12 (A) Electroformed nickel (B) CVD nickel.21 Figure 1.13 Various stages involved in the recovery of a plastically deformed material..23 Figure 1.14 Polygonization of the bent crystal containing edge dislocations 24 Figure 1.15 Mechanism of subgrains rotation and coalescence.25 Figure 1.16 Schematic representation of grain growth through atomic diffusion.26 vii

8 Figure 1.17 DSC scan of a 20 nm grain sized Ni deposits at heating rate of 10K/min. 27 Figure 1.18 Bright field TEM micrographs of electrodeposited nickel in as-received condition.28 Figure 1.19 Optical micrograph of the nanocrystalline nickel after annealing at 773K for 8 hours...30 Figure 1.20 XTEM and cross-sectional HRTEM micrographs of epitaxial nanotwinned Cu films at the tip of a nanoindentor...33 Figure 1.21 Snapshots of in-situ TEM micrographs of Σ3{112} ITBs before indentation...34 Figure 1.23 Schematics of three proposed detwinning processes..35 Figure 2.1 Bulk nickel prepared by nickel carbonyl vapour deposition with 16 mm thickness Figure 2.2 Schematic representation of bulk CVD nickel on the substrate with the growth direction indicated..36 Figure 2.3 Schematic of a heat flux DSC with disk-type measuring system. 37 Figure 3.1 Thermal stability investigation of CVD nickel through Differential Scanning Calorimetry at a heating rate of 20 C/min. 41 Figure 3.2 Optical micrographs showing the microstructure of as-received CVD nickel. 43 Figure 3.3 Grain morphology of CVD nickel at planar view after isothermal annealing for 45 minutes at various temperatures. 44 Figure 3.4 Grain morphology of CVD nickel at transverse view after isothermal annealing for 45 minutes at various temperatures. 45 Figure 3.5 Semi in-situ tracking of grain growth under optical microscope.46 Figure 3.6 Ultrafine and nano grains surrounding the large columnar grains in as-deposited CVD nickel in the transverse direction...48 viii

9 Figure 3.7 TEM micrographs showing the morphology of large columnar grains in as-deposited CVD nickel 49 Figure 3.8 (a) High density of growth twins inside the large columnar grains (b) Distribution of twin thickness showing all the twins are in the ultrafine and nano size range..50 Figure 3.9 TEM images showing five-fold twinned structures embedded in the bi-modal grains of as-deposited CVD nickel...52 Figure 3.10 Microstructure evolution of CVD nickel upon annealing at 200ºC, 400ºC, 600ºC and 800ºC in the planar direction observed under TEM..54 Figure 3.11 Microstructure evolution of CVD nickel upon annealing at 200ºC, 400ºC, 600ºC and 800ºC in the transverse direction observed under TEM 55 Figure 3.12 (a) Twin receding in the form of step movement (b) Dislocation configuration at the migrating step fronts..56 Figure 3.13 Receding of growth twins took place in a five-fold twinned grain 57 Figure 3.15 Partial dislocations formed at both the end of twins and the misfitted sites traversed by dislocation lines.57 Figure 3.16 Tangled dislocations along twin boundaries with some of them extending into the lattice.58 Figure 3.17 Twin boundaries transformed into dislocation array structures.58 Figure 3.18 (a) Twin boundaries turned into dislocation lines and disappeared (b) dislocations emitted from twin boundaries into the twinned crystal..59 Figure 3.19 Dislocation cell structures formed in the cleared grain interior after detwinning; twin boundaries became darker and blurred full of tangled dislocations...60 Figure 3.20 (a) One grain with half of the growth twins transformed into dislocation cells (b) Dislocation cells were formed both at the detwinning front and directly at twin boundaries...60 ix

10 Figure 3.21 (a) Dislocation arrays formed at the end of twins and along twin boundaries (b) detwinning by migration of dislocation arrays and dislocation emission..62 Figure 3.22 Morphology of nanotwins inside the large grains traversed by dislocation walls after annealing at 600 ºC...63 Figure 3.23 (a) Dislocation cell walls formed at the moving steps of a twin lamella (b) Nonuniform size distribution of dislocation cells at 800 ºC.64 Figure 3.24 Schematic demonstration of the formation of dislocation cells at the tip of a receding twin lamella. Traces of dislocations left from detwinning became part of a dislocation cell wall.65 Figure 3.25 Five-fold twinned structures in CVD Ni of as-deposited condition and after annealing at 400ºC and 600ºC 66 Figure 3.26 In-situ TEM images showing the microstructure of the same location upon annealing up to 661ºC. 68 Figure 3.27 Morphology of a five-fold twinned grain upon annealing from room temperature up to 655ºC in hot stage TEM.70 Figure 3.28 Microhardness of CVD nickel on planar and transverse direction upon annealing...71 Figure 3.29 Rockwell hardness of CVD nickel on planar and transverse direction upon annealing 71 Figure 4.1 Model of grain rotation and growth induced by deformation in nanocrystalline Ni 74 Figure 4.2 Proposed possible detwinning mechanisms for CVD nickel upon annealing..76 Figure 4.3 Strain energy density of twin boundaries and dislocations in polycrystalline copper..78 Figure 4.4 HREM image of the central area of the five-fold twinned grain..79 x

11 Figure 4.5 HREM images of the boundary between the matrix and five-fold twinned structure..80 Figure 4.6 Geometry of a grain boundary groove xi

12 List of Appendices Appendix 7.1 DSC scans at heating rate of 10ºC/min and 20ºC/min 87 Appendix 7.2 CVD nickel annealed at 200 ºC, 250 ºC, 300 ºC and 370 ºC for 45 mins...88 Appendix 7.3 TEM images of the black lines traversing the large grains.89 Appendix 7.4 Five-fold twinned structures in the as-deposited CVD Ni.. 90 Appendix 7.5 Detwinning phenomena as observed at 400 ºC, 600 ºC and 800 ºC Appendix 7.6 The hot stage in-situ TEM images showing a large grain full of growth twins surrounded by ultrafine and nano grains upon annealing up to 752 ºC.95 xii

13 1 1 Introduction In the past ten years, the nickel carbonyl vapor deposited nickel received major applications in the commercial fabrication of plastic molding inserts and other plastic forming tools. The technology is capable of growing fully dense deposits on aluminum mandrels with thickness up to 25mm 1. Unlike other deposition techniques, high accuracy replication with uniform thickness can be achieved by CVD process. The nickel shells produced by such chemical vapor deposition technique were found to show superior corrosion resistance as well as excellent strength and toughness due to the fine microstructure 2. With small addition of boron during deposition, extra strength can be added to the material for particular requirements. All the above features lead to the high volume use of CVD nickel to date. However, to further optimize the design and fabrication of the CVD Ni products, reliable thermal and mechanical data of the material are required. To the author s knowledge, up until now, many studies had reported the mechanical and wear properties of CVD nickel, but no one has examined its thermal stability. Thermal stability is especially important for materials with fine structure, which is associated with high internal energy raised from the intercrystalline components. On the other hand, any change in the microstructure upon heating will accordingly affect the mechanical properties. Considering the service temperatures for plastic molding process, for instance, the process temperature range for polypropylene is ºC, 3 the microstructural change of CVD nickel upon annealing becomes particularly important. The objective of this project is therefore set to investigate the thermal stability of CVD nickel through experimental study of its microstructural evolution upon annealing up to 800 ºC. The background knowledge on CVD process and the most up-to-date information on structure and properties of CVD nickel will be introduced in Chapter 1. The specific material used and experimental procedures involved in this study will be explained in Chapter 2. Results from various experiment observations will be presented in Chapter 3 followed by discussion in Chapter 4. The major finding from this project will be summarized in Chapter 5 with possible future work presented in Chapter 6.

14 2 1.1 Nanocrystalline materials Definition A material with lowest free energy is considered to be most stable. However, since last century it was found that when atoms deviated from their ordered structure position, the material would exhibit excellent features. The disordered structure is created by incorporating various defects such as dislocations, grain boundaries, vacancies, etc. Hence, the idea of nanocryatalline material is to generate a disordered structure by introducing all kinds of defects such that 50% or more of the atoms are located at the region of these defects. 4 As an example, Figure 1.1 shows the atomic structure of a two dimensional nanocrystalline material. The atoms in the boundary core region are deviated from their position as in perfect lattice and the degree of misfit depends on the misorientation between crystals. Figure 1.1 Atomic structure of a two dimensional nanocrystalline material, the atoms in the boundary cores are represented by open circles; the hexagonal array of atoms in grains are indicated in black. 4 Nanotechnology and nanomaterials has been the centre of focus in materials research field for a long time; yet there is still no exact definition of nanomaterials. Nevertheless, a working definition for nanomaterials and nanotechnology is generally accepted and is narrated as the design, characterization, production and application of structures, devices and systems by

15 3 controlled manipulation of size and shape at the nanometer scale that produces structures, devices and systems with at least one superior characteristics or property. 5 The term nanometer scale is typically within the range of several nanometers up to one micrometer. Nanostructured materials often have superior properties compared to those conventional materials with larger grain size scale. Nanomaterials are divided into four different classes according to the number of dimension, as shown in Figure The ability to artificially synthesize these nanocrystalline materials of microstructure from zero to three dimensions only became possible over the past twenty years. Figure 1.2 Schematic illustrations of four types of nanocrystalline materials showing structure in different numbers of dimensions. They could be 0-dimensional powders; 1-dementional thin deposited layers; surface coatings with nanostructure in 2 dimensions; and 3-dimentional bulk nanomaterials. 6

16 Preparation of bulk nanocrystalline materials The main synthetic methods in the preparation of bulk nanomaterials will be briefly discussed in this section. The wide popularity of techniques for preparing bulk nanocrystalline materials is because they are more suitable for investigation and observation compared with nanopowders. Nanopowders compaction is one of the most widely used methods to prepare compact nanocrystalline materials. 7 The deposition of films and coating also makes it possible to produce pore free nanocrystalline materials. Films and coatings can be produced by chemical and vapour deposition from gas phase, electrodeposition and sol-gel technology. Nanocrystalline alloys can also be produced by crystallization from amorphous state. Finally, severe plastic deformation is another way to produce pore-free metals and alloys with nano sized grains Compaction of powders As mentioned above, bulk nanocrystalline materials become the interest of fundamental and applied science due to their convenience for observation compared with nanopowders, which is much more complicated. Isolated nanocrystalline particles or nanopowders can be readily prepared by a wide variety of methods. The best known methods such as gas phase evaporation and condensation, deposition from colloidal solutions and plasma chemical synthesis and thermal decomposition are all fully developed and widely used. 7 The method of compacting nanopowders was developed between 1981 and1986 and has become one of the most widely used methods in the preparation of bulk nanocrystalline materials. 8 After the nanocrystalline powders are formed by one of the above methods, they are first pressed in vacuum at a preliminary low pressure around 1 GPa followed by a high pressure up to 10 GPa. At the same pressure, the smaller size of the particles, the less compacting capacity during pressing. As a result, high temperature is always applied to soften the material so that easier plastic deformation is allowed and better contact between the nanopowders is achieved. Meanwhile, diffusion process can take place at high temperature such that porosity can be greatly reduced to improve the final density. 9 However, the temperature should be controlled under a limit such that recrystallization will never occur in the materials. Moreover, sliding and shearing motion between the particles also help to get a better bonding. The resulting materials after pressing and sintering have the density up to 70% to 90% of the theoretical value; whereas

17 5 the density of nanocrystalline metal can reach up to 97%. 8 Compaction of nanopowders introduces very little contaminations by avoiding contact with other environment during the production of nanopowders and the compacting process. Overall, compaction of nanopowders is a promising method in the production of bulk nanocrystalline materials with high density Deposition of films and coatings Several deposition methods are developed and available to produce nanocrystalline coatings and films on a cold or heated substrate. Unlike the synthesis of nanoparticles, nanocrystalline materials form directly on the substrate surface rather than precipitating in the gas phase. Continuous deposition of thicker layer can be considered as a way of producing bulk free standing materials with the substrate stripped away. 9 The deposition process can take place from vapour, plasma and colloidal solutions. 8 Chemical and physical vapour deposition, namely CVD and PVD, are conventional ways to deposit films from gaseous reactants. In vapour deposition, the size of crystallite can be manipulated by varying the substrate temperature and evaporation rate. In the deposition from plasma, the process is mainly carried out using reactive working media and metallic cathodes. The electric discharge is maintained using an inert gas. Controlling the gas pressure and electric discharge parameters can regulate the thickness of the deposited films and the crystallite size in the films. Furthermore, deposition from colloidal solution is used in producing oxide semiconductor films. The deposition on the substrate from prepared solution is followed by drying and annealing. Co-precipitation method can be used to produce nanostructured films containing more than one type of semiconductor nanoparticles Crystallization of amorphous alloys Crystallization from amorphous state of materials is used in the production of nanocrystalline alloys. Amorphous metallic alloys are first produced by melt spinning, i.e., rapid cooling of the melt on a spinning surface. The amorphous alloys are then annealed at certain temperature allowing crystallization to take place. In order to achieve nanocrystalline structure in the materials, annealing is carried out in such manner that nucleation is promoted as much as possible while the grain growth process is kept at a low rate. Any heat treatment and deformation will only affect the grain size in the materials and will not change the phase composition of the alloy. 8 As a whole, the production of nanocrystalline alloys from

18 6 crystallization of quenched amorphous state has been developed rapidly as more and more types of alloys can be produced by this method Severe plastic deformation Severe plastic deformation is known as an attractive method to produce nanocrystalline materials within the grain size of 100nm. The resulting materials undergo multiple severe shear plastic deformation and form highly misoriented structure. Methods to achieve such severe deformation include equal channel angular pressing (ECAP), rolling, quasi-hydrostatic pressure and uniform forging. 8 Two common methods that provide high strain and fine grain structures are High Pressure Torsion and Equal Channel Angular Pressing, as shown in Figure 1.3. Figure 1.3 Schematic representation of severe plastic deformation setup (a) high pressure torsion (b) equal channel angular pressing 10 Severe plastic deformation introduces high density of defects into the final structure of the materials. Accompanying a refinement in grain size, the dislocation density, the concentration of point defects and stacking faults all increase. In addition to specific deformation conditions, the main factor that controls the structure and property of the materials is the mechanism of the deformation process. 8 As long as uniform stress and strain is applied throughout the entire materials, the deformation process is considered most efficient.

19 7 In a word, severe plastic deformation is an efficient way of forming nanostructured metals, alloys and other materials. It is also possible to produce bulk nanostructured materials without damage and residual porosity Mechanical properties of nanocrystalline materials Grain size dependence of strength and hardness Hardness measures the resistance of a material to plastic deformation; therefore it is related to yield strength σ y. Grain size is known to have a strong effect on hardness and strength which has been studied extensively on metals, alloys and ceramics. The Hall-Patch Law describes the relation between grain size D and yield strength/microhardness as shown below. Here σ 0 is the stress that keeps dislocation from moving in the material and k y is a constant. σ y = σ 0 + k y D -1/2 Vickers microhardness H v is related to yield strength by the empirical relation Hv/σ y = 3 at T melt. 11 And its grain size dependence is described in the following equation, where H 0 and k y are constants. H v H 0 + k y D -1/2 The above two equations show that materials exhibit enhanced mechanical properties with a decrease in grain size. It should be noted that Hall-Petch relation is not valid for both very large grains and ultrafine grains. The inverse relationship of size and strength is caused by the deformation mechanisms that dislocations slip and pile up at the grain boundary. In other word, grain boundaries in polycrystalline materials act as a barrier to dislocation motion during plastic deformation. 12 In small sized grains the dislocation pile-up length is shorter than in large grains and therefore allows less capacity for plastic deformation. Hence, the required stress to plastically deform the materials with smaller grain size becomes much higher. 9 On the contrary, when the grain size falls within tens of nano metre range, diffusion controlled creep at grain boundaries becomes an important mechanism even at room temperature that greatly increases deformation rate. It is also worth noting that grain boundary diffusion coefficient increases with decreased grain size. 8 Therefore, the effect of grain size on hardness and strength is ambiguous to define and depends on the ratio of change in yield strength and change in deformation rate.

20 8 Figure 1.4 presents a schematic diagram showing the Hall-Petch strengthening limit up to ~10 nm and the inverse Hall-Petch relation when the grain size falls less than 10 nm. Corresponding deformation mechanisms are also shown. Figure 1.4 A schematic diagram of Hall-Petch relation and inverse Hall-Petch relation Mechanical properties of bulk nanomaterials The elastic modulus in nanocrystalline coppers subjected to severe plastic deformation was reported to decrease compared with conventional polycrystalline copper. 13 This is because the reduction in grain size leads to higher volume fraction of grain boundary constituents, in which the atoms are highly misoriented and close to amorphous state. Since the elastic modulus of amorphous materials is only about half the value of polycrystalline materials, materials with nanometre sized grains show apparent decrease in elastic modulus. According to various review papers on study of the stress-strain behaviour of nanocrystalline Pd, Cu, Mg alloys and many other alloys, it was concluded that the tensile strength of nanocrystalline metals was about1.5 to 8 times higher than that of polycrystalline metals with coarser grains. 8

21 9 1.2 Chemical vapor deposition Introduction Chemical vapour deposition (CVD) process has rapidly developed in the past few decades. High purity materials with high performance can be produced from the process. It is used extensively in the commercial deposition of metals, metal alloys and ceramic coatings. In a typical CVD process, a properly heated substrate is exposed to one or more gaseous chemical precursors, which are then thermally decomposed near the substrate or reduced by an appropriate gas on the substrate surface to produce the desired deposits. CVD process has been used substantially to fabricate thin films. However, when given longer deposition time, it is also capable of producing deposits with thickness exceeding 15mm. These bulk deposits can be readily obtained because the deposition rate of CVD process is much faster than that of commercial electrodeposition. 14 The deposited CVD films usually exhibit good adherence and excellent uniformity in thickness, even on those substrates with irregular shape and sharp corners. With the nature of uniform thickness achieved and bulk deposits capability, 3D components with highly complicated geometry could possibly be fabricated without additional machining or welding. According to a National Technical Information data search of CVD research report, the applications of the CVD technique ranges from optically and electrically sensitive thin films less than 0.01mm thick to bulk SiC production in the ceramic turbine blade research. 15 The composition and concentration of gaseous precursors during deposition process can be varied to control the microstructure the final deposits. Two or more than two precursors can be decomposed simultaneously Nickel Vapour Deposition CVD nickel is used in many diverse applications, such as plastic molding inserts and other plastic forming tools, inserts for printing plates, reflectors and other optical applications, nickel shapes, nickel tooling, etc. When small amount of boron is added, it can also produce high strength structural parts. 16

22 Methodology Most metals have been deposited by CVD process in the production of thin films. Among them, CVD nickel is a promising material owing to its great corrosion resistance. 17 In fact, the low temperature thermal decomposition of nickel tetracarbonyl (Ni(CO) 4 ) to form nickel deposits with very little contaminations from foreign atoms has been of commercial interests for many years. Due to the toxic and corrosive features of most reactants, a closed system is usually required when depositing materials with CVD process. Figure 1.5 Schematic representation of the CVD reactor. R for gas regulator, L for Liquid flow meter, G for gas flow meter 18 The sketch in Figure 1.5 schematically illustrates the CVD reaction apparatus for decomposition of Ni(CO) 4. The liquid Ni(CO) 4 is first vaporized in a heated line with added CO gas. The gas mixture is then transferred into the sealed reaction chamber together with argon gas. The substrate in the chamber is usually resistively heated at temperature between 110 ºC and 200 ºC; the deposition rate ranges from µm/h. 2 The exhaust gas mixture must be pumped away through two getter furnaces to remove any un-reacted Ni(CO) 4 before discharging it into the

23 11 atmosphere. 18 The mandrels or substrates used in the nickel vapor deposition process could be made of aluminum, steel or brass Chemistry of Nickel Vapour Deposition The thermal decomposition of nickel tetracarbonyl to form nickel deposits involves the reaction as described in the following: Ni(CO) 4 Ni + 4CO This reaction was originally used to purify nickel in the early 1900s and it is still actively used in the production and purification of metallic nickel to date. The mean metal-carbon dissociation energy was calculated to be kj/mol 19. The optimum temperature for this reaction to occur is at 180 ºC-200 ºC. When the temperature goes beyond 200 ºC, carbon tends to deposit together with nickel. The operating pressure in the chamber is up to 1 atm. The decomposition product could be solid film or nickel power, depending on the deposition parameters of the process. The reactant, Nickel carbonyl vapor, can be obtained by direct reaction of CO gas with finely divided nickel powder at 80 ºC and 1 atm. Chemically, the reactions occur in two steps as follows: 16 Ni + O 2 NiO 2 NiO 2 +6CO Ni(CO) 4 + 2CO 2 The deposition process from Nickel tetracarbonyl gas was first discovered by Ludwig Mond. Even today, the process of nickel carbonylation and its decomposition at low temperature is still used to produce high purity nickel (99.998% nickel). 20 However, due to the extremely toxic nature of Ni(CO) 4 (threshold limit for human exposure is 10-3 ppm) and the bi-product carbon monoxide gas, many other non-toxic precursors have been investigated and used in the processing of nickel thin film by chemical vapour deposition, such as NiCp 2, Ni(MeCp) 2, Ni(den) 2, etc. 17 The choice of precursors depends on the specific requirement of the process and commercial application of nickel deposits. Among all these possible precursors, nickel carbonyl still remains the best choice when the lowest operating temperature is required.

24 Control of reaction rate Understanding the kinetics of the CVD process is particularly important to know the process as well as to optimize the controlling parameters in order to obtain the deposits with desired structure and properties as required in various applications. 21 For instance, more ductile nickel can be obtained by lower reaction rate. The influence of various deposition conditions on the reaction rate is discussed in this section. The rate limiting step of CVD process may be described as two controlling types, the mass transfer/diffusion and the surface reaction kinetics. At medium gas flow rate and/or high substrate temperature, the deposition rate of CVD deposits is solely determined by diffusion in the gas phase. 21 This type of mass transfer control assumes that there is a gas layer boundary above the substrate surface on which the deposition reaction takes place. A concentration gradient exists between the reaction zone and the substrate surface. While at high gas flow rate and/or low temperature, the deposition rate is determined by kinetically controlled surface process/chemical reaction. It occurs when the reaction at the surface is a slow step. Although many studies of CVD process made the assumption that the heterogeneous reaction took place at the freshly deposited nickel surface, many observations showed that the homogeneous reaction in the gas phase also played an important role in CVD process. The homogeneous thermal decomposition of Ni(CO) 4 was distinguished from heterogeneous reaction, and two different kinetic laws were proposed. 22 The possible parameters that would influence the deposition rate in CVD nickel process such as 17, gas flow rate, temperature, total pressure have been studied and described extensively In the decomposition process of nickel carbonyl, the reaction rate increases remarkably with the increase of gas flow rate at any given temperature and pressure. According to Clements et al., the rate of reaction was controlled by the rate of mass transfer of reactant or product across the gas phase boundary layer above the heated substrate. 19 Thickness of nickel deposits can be achieved with prolonged duration of the process; the typical deposition rate in actual process is approximately mm/h 19.

25 13 As temperature increases, the reaction rate first increases slightly up to 437K, then it either remains constant given a high flow rate or decreases when the flow rate is relatively low. 19 The initial increase of the reaction rate with the increasing temperature is possibly due to the increased thermal convection, which results in the accelerated removal of the CO gas from the reaction site and therefore pushes the reaction forward. Apparently this is a mass-transfer controlled process in the gas phase as reported by Clements et al. The transition of reaction rate at a specific temperature as mentioned above is possibly caused by the interference of the removal of carbon monoxide gas away from the substrate with the arrival of the nickel carbonyl to the substrate, therefore leading to a suppression of the deposition reaction. Thus, when the temperature goes beyond the limit, in the case of nickel, above 437k, due to the enhanced thermal convection, the decomposition of the nickel carbonyl occurs in the gas phase layer above the heated substrate and produces black metallic powders that are evident in the pipes and at the outlet of the reactor. This is also known as gas phase precipitation. Gas phase precipitation is not desired in the production of CVD films because the precipitated particles incorporate into the deposits in the form of soot, leading to surface roughness of the deposited film, the non-uniformity in film thickness and poor adhesion. 16 However, it is the major mechanism in the production of extremely fine powders. Pressure is also a rate limiting factor same as temperature since the diffusivity of gas is inversely proportional to its pressure. Low pressure reduces the possibility of the homogeneous nucleation and promotes heterogeneous nucleation. In the work of Clements, however, pressure appeared to have no significant effect on the rate of reaction, where the investigated pressure range is 4-10mmHg and the temperature range was 423K-548K. The low sensitivity of reaction rate with pressure change indicated a diffusion controlled process under which the rate of reaction was mostly dependent on flow rate that affects the boundary layer thickness Structure and grain morphology of CVD nickel Property of CVD materials is closely related to their structure, which is further controlled by deposition parameters. The exact deposition mechanism on how the film is formed on a surface by CVD process is still controversial. 27 Therefore this section will only describe the resulting structure and grain morphology of CVD nickel deposit.

26 14 In general, the structure of CVD deposits can be classified into three types shown in Figure 1.6. These structures can be controlled by manipulating the deposition conditions such as temperature, total pressure and gas flow rate. When the operating temperature is high, grains tend to grow in columnar shape as no interruption to the growth is induced. The columnar structure appears to be more pronounced when the thickness of the deposits increases. These structures may lead to the anisotropic nature in electrical, chemical and mechanical properties. On the other hand, at low pressure or/and low temperature, kinetically controlled surface process becomes the rate controlling step and the deposited structure tends to be fine grains as shown in Figure 1.6(c). 16 The fine grain structure is usually desired due to their great properties. Sometimes two or all the three types of the structure may co-exist in the structure of CVD deposits, which can be frequently found in thick deposits since achieving uniform structure is difficult. Figure 1.6 Schematic of three types of structures obtained by CVD: (a) columnar grains with domed tops, (b) faceted columnar grains, (c) equiaxed fine grains. 16 The observation of CVD nickel using optical microscopy showed it possessed a dendrite structure with the grains fall in two size ranges 2. The large columnar grains were separated by many much smaller equiaxed grains, as shown in Figure 1.7. This bi-modal grain morphology is similar to the result of the co-existence of columnar and fine grain structures shown in Fig 1.6 (b) and (c). And this type of microstructure is typical when the substrate temperature is much lower than the melting point of the deposited nickel. 14

27 15 Growing surface Figure 1.7 Optical image showing typical grain structure of pure CVD nickel 2 At the planar view parallel to the substrate surface (Figure 1.8), the bi-modal microstructure consisted of equal fraction of large grains with diameters of µm and much smaller grains around 0.1µm in diameter. The latter appeared to have much higher density of defects. 18 According to the microstructure observation, most of the large grains were twinned with low density of dislocation. While the ultrafine grains did not show any texture, the large columnar grains showed preferred orientation or growth direction in the <110> direction. 2 Figure 1.8 Microstructure of a CVD pure nickel. 18

28 16 Figure 1.9 (a) Five <011> subgrains meet to form a pentagon (b) zone axis diffraction pattern of the pentagon in (a) 18 A special form of grain structure worth mentioning was found in the CVD nickel deposit, where five large grains in the <011> directions meet at 72ºangles to form a pentagon, as shown in Figure 1.9 (a). The zone axis diffraction pattern on the pentagon done by Campbell et al. indicated a near five-fold symmetry {Figure 1.9 (b)}. Such structure containing five identical tetrahedrons is termed as Multiply Twinned Particles (MTP s). In face-centre cubic (FCC) crystal structure materials the twin boundary planes is {111}. Therefore, the angle between two sets of (111) planes in MTP s should be 70.5º. However, analysis on the diffraction pattern from the pentagon suggested a rotation of the grain about the zone axis < 011 > by multiples of 72º. An explanation for the mechanism of formation of this kind of structure in gold was proposed by Yang in In the early stage of vapour deposition process, small clusters of gold atoms formed with diameter of 5-20nm and some of them had pentagonal profile. Five identical tetrahedral units that form the regular decahedron had the body-centre orthorhombic (BCO) crystal structure, which was slightly distorted from FCC structure. The resulting structure then had the {111} twin planes in FCC replaced by {101} planes in BCO where the five twin components were separated exactly by 72º. As the clusters grow to larger size, usually a few tens of nanometre, the BCO structure reverted to more stable FCC structure. Therefore, the

29 17 pentagonal structure evolved from the regular decahedral cluster with BCO structure in the tetrahedrons. The geometrical relation between the BCO and FCC structure was elaborated in Figure The resulting misfit at the twin boundaries is accommodated by dislocations and secondary twins. 18 Figure 1.10 (a) A regular decahedron consisting of five irregular tetrahedrons, such as ABCD. AC = AD = BC = BD, AB =1.0515AC. (b) Relationship between the fcc and bco structure. The bco unit cell is shown with a unit tetrahedron A'B'C'D' indicated. The lattice is fcc if a = b and c = a. A'B'C'D' is then a regular tetrahedron. With a small distortion from fcc, the lattice becomes bco. The pertinent relationships for bco are b = a and c = a. A'B'C'D' is then irregular and similar to ABCD in (a) Mechanical properties of CVD nickel Much work has been done on CVD nickel to collect the important data in its mechanical and thermal properties. The data summarized in Table 1.1 was obtained from the tests carried out at Ortech, Ontario Hydro s Research division, Canmet and University of Toronto. 2

30 18 Table 1.1 Properties summary of CVD nickel 2, conventional nickel and electrodeposited nickel. 29,83 (GS: Average grain size) Properties CVD nickel Conventional nickel (GS: 10 µm) Electrodeposited nickel (GS: nm) Yield strength MPa 103 MPa MPa Ultimate tensile MPa 403 MPa MPa strength Modulus of elasticity GPa 207 GPa 204 GPa % Elongation % 50% 4-7% Hardness HV 140 HV 650 HV It can be seen that CVD nickel has superior strength compared with the conventional nickel, but not as good as the electrodeposited nickel due to the bigger grain size. However, it also has quite decent %elongation that is better than electrodeposited nickel with very small grain size. Therefore, CVD nickel is provided with the toughness required for most applications. In addition to the high strength and great toughness, CVD nickel also has excellent corrosion resistance and wear resistance due to its full density and low porosity; low residual stress in the deposits helps to ensure the accurate duplication of the substrate or mandrels. Table 1.2 Chemical composition of CVD nickel. 2 Ni 99.98% C <150 ppm S < 1.0 ppm H < 7.3 ppm CVD nickel is usually deposited with very little contaminations from other elements. Typical chemical composition of CVD nickel is presented in Table 1.2. Boron was usually added during deposition process to achieve better product performance. The addition of small concentration of boron could effectively solute strengthen the pure CVD nickel and meanwhile result in a finer

31 19 grain structure. Experimental data showed that increasing in boron content led to an increase in the micro-hardness as presented in Figure 1.11 below. 18 Figure 1.11 Micro-hardness as a function of boron content for as deposited and Ni and Ni-B alloys. ( as-deposited, 350, 1 h, 600, 1 h, 800, 1 h.) 18 Skibo et al. summarized the yield and tensile strength, % elongation for pure CVD nickel and CVD nickel with various amount of boron added in Table However, the samples were obtained directly from Vaporform Inc., thus the specific information about the deposition parameter was unknown. The yield and tensile strength were found to increase as a function of boron content. On the other hand, the total elongation was affected very little by the boron addition.

32 Table 1.3 Mechanical properties of Ni and Ni-B alloys

33 Advantages of nickel vapor deposition over electroforming CVD nickel process is based on thermal decomposition of reactants rather than an electrolytic deposition. Therefore, as long as the substrate/mandrel provides uniform temperature, the deposits will be uniform in thickness no matter how complicated the geometry of the substrate surface is. However, this is not the case in electrodeposited nickel as substrate geometry has significant effect on the thickness of deposits. At certain places like sharp corners, the electrical fields are much higher than other places causing nickel to build up at these locations. The comparison of nickel deposits formed by electroforming (EF) and CVD process is illustrated in Figure With CVD nickel process, both the internal and external corners grow at the same rate; hence it gives more uniform thickness and properties in the deposited shells. Therefore, much more flexibility is provided by CVD nickel process that could not be achieved in EF nickel sometimes. Figure 1.12 (A) Electroformed nickel; (B) CVD nickel 2 Generally speaking, CVD nickel is also denser and more ductile than electroformed nickel. The full density of CVD nickel contributes to its superior corrosion resistance. The enhanced toughness and strength gives high quality nickel products with prolonged lifetime. The possibility of products failure due to brittle cracking is greatly reduced. Moreover, EF nickel usually contains sulphur making good welding impossible. But for CVD nickel, welding can be performed at any time for any repair or modification purpose since it contains almost no sulphides. High purity deposits production is a promising advantage and makes CVD the only acceptable technique for nuclear applications. Furthermore, the deposition rate of CVD nickel is 0.25mm/hr, which is 20 times faster than EF nickel, making the bulk nickel production feasible. Lastly, CVD process requires shorter delivery time and lower cost to duplicate shells.

34 Thermal stability of nanocrystalline materials Nanocrystalline materials are considered thermally unstable structures due to the excess free energy associated with their large grain boundary area and triple junctions. Study of the energy release and microstructure evolution upon thermal activation is quite significant in the better understanding their thermal stability. This is usually accomplished by anisothermal or isothermal annealing Thermal stability of cold worked materials A metal is said to be cold worked when it is plastically deformed at a temperature relatively low compared to its melting point. 30 Plastic deformation process introduces high density of defects and strain hardens the material while the grain shape is also changed. A fraction of energy is stored in the metal as strain energy associated with the newly created dislocations and point defects by deformation. This energy retained in the materials results in a thermally unstable structure. A cold worked material would release its stored energy and return to its original state in both structure and property under appropriate heat treatment. Studies of the energy release process of cold worked metals showed that three stages are involved during annealing, namely recovery, recrystallization and grain growth Recovery Recovery refers to the change in a deformed material prior to recrystallization that leads to the partial restoration of properties to the state before deformation. 31 During recovery, the stored internal strain energy is lowered by dislocation movement, which is a result of enhanced atomic diffusion at high temperature. There are two main mechanisms involved in the recovery process, dislocations annihilation and dislocations rearrangement into lower energy configuration. 31 Besides, in recovery stage, the mechanical properties are partially recovered to their original values before deformation. 12 A series of events as shown in Figure 1.13 is involved in the evolution of dislocation configuration during recovery. Some of these stages may have occurred during deformation before annealing, known as dynamic recovery. Upon cold working, unequal number of dislocations with opposite signs is generated. Therefore, annihilation process could not remove the excess dislocations. These dislocations will rearrange themselves into new configurations

35 23 with lower energy in the form of dislocation arrays or low angle grain boundaries (LAGBs). The simplest case is polygonization with tilt boundaries produced in a bent single crystal when all the dislocations left have only one burger s vector, shown in Figure In alloys with medium to high stacking fault energy, the dislocation configuration after deformation is usually more complex in the form of 3D dislocation cell structure with the cell walls being dislocation tangles 31. Upon annealing, annihilation of some redundant dislocations occurs while the rest in the tangled cell walls rearranges into low angle grain boundaries. The resulting structure is called subgrains. The stored energy can be further recovered by growth of the subgrains leading to the reduction in total area of LAGBs. The above events from tangled dislocations cells to subgrain growth are depicted as distinct stages in Figure 1.13(c), (d), (e). 31 Figure 1.13 Various stages involved in the recovery of a plastically deformed material 31

36 24 Figure 1.14 Polygonization of the bent crystal containing edge dislocations. (a) as-deformed, (b) after annihilation, (c) tilt boundaries formation 31 Two different mechanisms were proposed for subgrain growth, the subgrain boundary migration and the subgrain rotation and coalescence. Based on the in-situ TEM observation of Fe-Si alloys upon annealing, Hu suggested that two adjacent subgrains might rotate until they obtain similar lattice orientation through diffusion process along grain boundaries. Subsequently, the adjacent grains with similar orientation coalesced into one larger grain. 31 The mechanism is schematically presented in Figure This process of rotation and coalescence is thermodynamically feasible and the driving force arises from the reduction in grain boundary energy by eliminating low angle boundaries.

37 25 Figure 1.15 Mechanism of subgrains rotation and coalescence Recrystallization Upon completion of recovery, the deformed metals still retain relatively high strain energy and recrystallization occurs at this stage. Recrystallization is known as the nucleation and growth of new strain-free grains with low dislocation density in the plastically deformation grains. 12 The driving force for recrystallization arises from the internal energy difference between strained and unstrained grains. The process depends on both time and temperature because the major mechanism involved is short range diffusion. The resulting structure is the same as the one in undeformed condition: the large equiaxed grains. Thereby the superior mechanical properties due to cold work are restored to the original value before plastic deformation Grain growth The strain free grains continue to grow at an appropriate temperature after recrystallization is completed. This process is called grain growth and it is a phenomenon commonly seen in all polycrystalline materials, not necessarily after recrystallization. The driving force for grain growth is the total energy reduction via the progressive elimination of grain boundary area as

38 26 grains grow to larger size. Normal grain growth takes place by grain boundary migration where some grains grow larger and some grains shrink and eventually vanish. The grain boundary motion is accomplished by short range diffusion of atoms from the shrinking grain to the growing grain across the boundary. The process is depicted in Figure The grain growth law for most polycrystalline materials follows the relation below: where d 0 is the initial grain size at time t = 0, K is a kinetic constant and n is the grain growth exponent. 32 Again, because normal grain growth is diffusion dominated process, it depends on both temperature and time. At increasing temperature, diffusion rate is enhanced; therefore grains grow much faster. Figure 1.16 Schematic representation of grain growth through atomic diffusion Thermal stability of nanocrystalline nickel Owing to the ultrafine grain size of nanocrystalline materials, there is high amount of energy stored in the intercrystalline components, i.e., grain boundaries, triple junctions and quadruple

39 27 nodes, and other types of defects. For instance, a nanocrystalline material having an average grain size of 5 nm contained 50% volume of interfacial boundaries. 34 As such, the nanostructured materials are thermodynamically unstable and tend to transform back to the coarse grain structure through enhanced grain growth. This leads to the loss of superior properties of nanocrystalline materials due to grain growth. Therefore, a fundamental understanding of the thermal stability of these materials becomes a critical issue for industrial applications. Over the past few decades, the thermal stability of nanocrystalline nickel had been studied extensively 33-42, most of which were performed on electrodeposited nickel. The experimental results on grain growth of nanocrystalline nickel upon annealing will be briefly reviewed. According to many of the studies on nanocrystalline nickel electrodeposits, it was found that different annealing conditions gave rise to very different post-annealing structures in. Little agreement existed between different studies. Base upon the distinctive growth mechanisms brought up at different annealing temperature, it was proposed that grain growth of nickel electrodeposits went through a multi-stage process. 38 Stage I Stage II Stage III Figure 1.17 DSC scan of a 20 nm grain sized Ni deposits at heating rate of 10K/min. 37 By looking at the DSC scan of a 20nm electrodeposited nickel heated at 10 K/min in Figure 1.17, three exothermic stages were shown 37. The first exothermic process that started at 394K and ended at 561K was identified as the nucleation and abnormal grain growth stage. The major

40 28 peak of heat release (561K-643K) was the normal grain growth stage. The third and weakest reaction was assigned to be the growth towards equilibrium stage extending from 643K to 773K. The nucleation stage, or the formation of first growing grain, was similar to the subgrain rotation and coalescence known in recovery process. In other word, the recovery was realized by accidental removal of low angle grain boundaries via climbing of edge dislocations. 37 This led to the subsequent abnormal grain growth of minority of grains rapidly consuming the surrounding nanograins. This process was quite similar to conventional recrystallization process where strain free grains nucleated and grew at the expense of adjacent highly deformed/strained regions 31. From 561K, normal grain growth started with all the grains growing at similar rates. The activation energy calculated from DSC signal KJ/mol was close to the activation energy of grain boundary self diffusion in nickel. 42 This indicated the migration of high angle grain boundaries during grain growth was similar to diffusion at the temperature corresponding to the DSC signal peak. Bi-modal grain structure was detected at this peak temperature. Figure 1.18 Bright field TEM micrographs of electrodeposited nickel in as-received condition (a) and after annealing at 420 ºC for ~1s (b), 30s (c), 3.6ks (d), 39.6ks (e), 432ks (f). 39

41 29 The microstructure evolution and multi-staged grain growth of nanostructured nickel upon annealing was summarized in Figure The initial average grain size of the nickel deposits was 20 nm. As shown in Figure 1.18b, abnormal grain growth occurred right after ~1s of annealing at 420 ºC. During this stage, a small fraction of grains grew rapidly by consuming the surrounding nanocrystalline matrix. The nanograined matrix remained unchanged in size until they were consumed by the growth front of the abnormal grains. This was identified as the stage I abnormal grain growth and it totally completed after 30s of annealing at 420 ºC, as shown in Figure 1.18c. Upon completion of stage I, the stage II normal grain growth took place but at a greatly reduced growth rate (Figure 1.18d). Beyond this point, stage III abnormal grain growth occurred between 3.6ks and 39.6ks shown in Figure 1.18e. Large planar grains (5~50µm) were found embedded in the matrix of nanograins with mean grain size of 500nm. This was also referred to as the late stage abnormal grain growth. Finally by 432ks of annealing the large abnormal grains consumed all the submicron grains and all the grains continued to grow normally leading to a uniform microstructure (Figure 1.18f). The late stage abnormal grain growth following the normal grain growth attracted much interest. A typical micrograph of the abnormal grains with planar boundaries were shown in Figure Further investigation by STEM with energy dispersive X-ray spectroscopy (EDS) on the planar boundaries in Stage III abnormal grain growth showed significant amount of sulphur-rich second phase at the planar boundaries of the abnormal grains. 39 The wetting effect could probably explain the planar boundary front resulting from the rapid migration of grain boundaries without interfered by the misorientation dependent boundary velocity.

42 30 Figure 1.19 Optical micrograph of the nanocrystalline nickel after annealing at 773K for 8 hours 41

43 Twinning and detwinning in FCC materials Twinning and its effect on materials properties Nickel is likely to form twins when subject to straining or annealing attributing to its low stacking fault energy. Twinned structure usually exhibits favourable properties such as corrosion resistance and improved ductility. 43 Moreover, nanocrystalline materials with high density of twin boundaries possess better electrical conductivity than those without twins. 49 Twins are also considered not susceptive to crack propagation. The twin boundary energy varies depending on the structure of the boundaries. For example, a perfect coherent twin boundary (Σ3) lies on a {111} plane has the lowest energy of ~ 0.01J/m 2, while a random grain boundary possesses 1J/m 2 or more. 43 In recent years, twin and twin boundaries have attracted increasing attention because nanostructured metals often contain large number of twins that are either induced by deformation or grown in during synthesis. 45 Nanotwins were observed to show strengthening effect during plastic deformation of nanocrystalline metals. 46 The mechanical properties of nanocrystalline materials depend on the dislocation interaction with grain boundaries and twin boundaries. The twin boundaries formed within a grain cause change of crystal orientation between the twinned region and the matrix, leading to a discontinuity of slip planes across the twin boundaries. Experiments showed that nano-sized growth twins introduced in nanostructured materials acted as strong barriers to dislocation motion and hence gave an unusual combination of high strength and good ductility Furthermore, the excellent high temperature fatigue behaviour of superalloys is contributed to a large extent by the presence of twins and twin boundaries. 52 As such, twinning plays an important role in the mechanical properties of many materials Stability of nanotwins in FCC materials Twinning is normally a result of permanent strain and considered to be not recoverable. However, the molecular dynamics simulation performed on nanocrystalline Al, Cu, Ni and Co recently showed that under high stress twin boundaries could accommodate plastic strain by boundary migration through activity of partial dislocations on the planes right next to the twin boundaries. In other word, twin boundaries can be moved when they are subject to external stress.

44 32 In a recent study of polycrystalline aluminum, twinning and subsequent detwinning were observed near a crack tip during tensile straining. 57 Twinning was totally reversible in this high purity polycrystalline Al. Two stages for the detwinning were observed: the thinning of twinned region followed by shortening of twin boundaries. The observed twin had a thickness of 3 nm and a length of 25nm with its twin boundary plane parallel to the crack plane. Steps were seen to form on the twin boundaries which were regarded as a bundle of partial dislocations emitted from the crack tip. 57 Movement of these steps, i.e., propagation of the partial dislocations along the twin boundary, led to the thickness reduction of the twin lamella. The author attributed the spontaneous detwinning to the high stacking fault of aluminum and also the relatively low Peierls-Nabarro force against gliding of partial dislocations. In the first stage, thinning of twins by step movement was mainly driven by the line tension of the partial dislocations. After the twin lamella was thin enough, the whole twin receded together in the second stage more quickly with the additional driving force from the twin boundary energy. Twin boundary migration was also observed in electrodeposited copper with nanoscale growth twins during tensile straining. 45 The copper contained large number of nanoscale growth twins and coherent twin boundaries with Σ3 configuration. The twins had thickness of a few to tens of nanometres and twin boundaries were perpendicular to the applied stress. Under in-situ TEM, twin boundary migration was observed via slip of shockley partial dislocations that emitted from the twin boundary/grain boundary intersections during tensile straining. 45 In addition, the authors suggested that twin boundary migration might be the preferred deformation mode in the early stage of plastic straining. Another study performed on sputtered high purity copper showed similar results along with more extensive analysis. 58 Topographical analysis together with atomistic simulations showed the detwinning process was accomplished via collective glide of partial dislocations that form an incoherent twin boundary (ITB). Figure 1.20 showed the cross sectional view of the nanotwinned Cu films before indentation was applied. Two types of twin boundaries were shown in Figure 1.20 (b), Σ3 {111} coherent twin boundary (CTB) and Σ3{112}incoherent twin boundary (ITB). ITBs indicated the termination of twin structure in the lattice.

45 33 Figure 1.20 (a) XTEM micrograph of epitaxial nanotwinned Cu films at the tip of a nanoindentor; (b) Cross-sectional HRTEM of the films along <110> zone axis; (c) Magnified HRTEM micrograph of the Σ3{112} ITB shown in (b). 58 Upon nanoindentation test, the in-situ XTEM micrograph presented in Figure 1.21 showed the migration of two ITBs. It was revealed that the ITBs migrated in the form of a step with the height of several {111} atomic distance at an average speed of 3.6nm/s.

46 34 Figure 1.21 Snapshots of in-situ TEM micrographs of Σ3{112} ITBs before indentation (a), at 33s (b) and at 55s (c); the migration distance of three ITBs as a function of time (d). 58 For both growth and deformation twins in FCC materials, the ITBs, i.e., the end of twins, are comprised of arrays of Shockley partial dislocations. 59 Twinning dislocations in deformation twins usually have the same burgers vector so that the dislocations move in the same direction. But this is not necessary for growth twins in sputtered nanotwinned copper. 58 The twinning dislocations would dissociate into one pure edge partial dislocation and the other two mixed partial dislocations with opposite screw components. With applied shear stress the ITBs would migrate easily; meanwhile, the dissociation of twinning dislocations reduced the Peierls force and further assisted the migration. Wang et al. 58 predicted the detwinning mechanism of nanotwins with both ends intersecting with grain boundaries. Three processes were proposed and schematically shown in Figure (1) A twinning dislocation was nucleated at the grain boundary and glided inside the twinned region without interfering with the twin boundaries, as shown in Figure 1.23 (a). Shear stress was required to move the dislocation.

47 35 (2) A single twinning dislocation was nucleated at the grain boundary and glided along the twin boundaries, shown in Figure 1.23 (b). The twin thickness was changed by one {111} planar distance and shear stress was needed. Another case presented in Figure 1.23 (c) showed a number of twinning dislocations nucleated at the grain boundary and continuously glided along the twin boundaries, leaving behind the same number of twinning dislocations at grain boundaries with opposite sign to the gliding dislocations. Local relaxation at grain boundary was required due to these edge dislocations, either by shear deformation of neighboring grains or climbing in grain boundaries. (3) Multiple dislocations nucleated at grain boundary and glided together, reverting the twinned lattice back to matrix stacking. The two twin boundaries were removed simultaneously. External shear stress was not necessary and the process was driven by the twin-twin interaction. 58 The third mechanism proposed about nucleation and glide of twinning dislocations was considered the most energetically favoured by eliminating the twin boundaries. However, the Peierls barrier was higher than the other two processes. Thus, the process described in Figure 1.23 (d) would most likely occur for thin twins. Figure 1.23 Schematics of three proposed detwinning processes. The red lines represent CTBs, the black dash lines represent (111) planes and the blue area is twinned region. 58

48 36 2 Experimental 2.1 Material The CVD nickel used in this project was provided by Weber Manufacturing Technologies Inc. The nickel deposits were produced by a well developed chemical vapor deposition process at INCO s Coppercliff Carbonyl refining facilities. Nickel tetra-carbonyl gas, Ni(CO) 4, was used as the reaction gas in this process. Yet the exact deposition parameters are not available to the public. Figure 2.1 Bulk nickel prepared by nickel carbonyl vapour deposition with 16 mm thickness The nickel deposit is of rectangular shape, measuring 16 mm in thickness, 40.5 mm in length and 5 mm in width. The actual dimension of the sample is shown in Figure 2.1. Figure 2.2 Schematic representation of bulk CVD nickel on the substrate with the growth direction indicated. The transverse direction is perpendicular to the substrate and the planar direction is parallel with the substrate.

49 37 Since the microstructure of CVD nickel is known to be anisotropic, the investigation on the evolution of its microstructure and mechanical property will be carried out on two directions, namely the planar direction and the transverse direction as schematically shown in Figure Thermal stability analysis via DSC Differential Scanning Calorimetry (DSC) is a thermal analysis technique that has been used to measure the heat flow related with any kind of phase transition or chemical reaction in materials. 60 Qualitative and quantitative information can be obtained about any change caused by the exothermic and endothermic reactions. The DSC apparatus used in this project was disk-type heat flux DSC as shown in Figure 2.3. During a DSC scan, the sample of interest and a reference material are placed in separate chambers sitting symmetrically to the centre. The heat supply from the furnace is passing to the samples through the conducting chambers. The two chambers are heated linearly with time and maintained at the same temperature. If the sample material experienced any physical or chemical transition, it would require either more or less heat input so as to maintain the same temperature with the reference material. In that case, a differential signal Δ T is then generated and the corresponding heat flow difference is calculated by the machine. This difference in heat flow is recorded as a function of temperature to produce the DSC curve. Figure 2.3 Schematic of a heat flux DSC with disk-type measuring system. S: crucible with testing sample, R: crucible with reference sample. 60

50 38 It has been reported that DSC was an effective tool to study the thermal stability and nature of nanocrystalline materials. 61 The principal of using DSC to study grain growth was described by Chen 62. By using DSC, the thermal effect related to reduction in grain boundary component during microstructure evolution was measured. Nucleation and grain growth during annealing of nanocrystalline materials would give different DSC signals. In this study, the thermal stability and grain growth behaviour of the as-deposited CVD nickel was analyzed by Dupont 910 Differential Scanning Calorimeter purged with argon gas. The samples were first prepared from various locations of the bulk nickel and then heated from room temperature up to 600 C at two different heating rates: 10 C/minute and 20 C/minute. Heat flow (W/g) difference during the non-isothermal annealing was obtained as a function of increasing temperature. Any microstructure change would be reflected by the exothermic and endothermic peaks on the resulting DSC curve. 2.3 Microstructure evolution investigation The microstructure evolution CVD nickel was first analyzed by optical microscopy and transmission electron microscopy (TEM). Samples were annealed at 200 C, 400 C, 600 C and 800 C, respectively. Semi in-situ optical microscopy and hot stage in-situ TEM were then employed in order to track the grain boundary migration and twin boundary movement. Samples were cut into small pieces from random locations of the bulk material for further preparation. Detailed experimental methods are described in the following sections Optical microscopy A preliminary analysis on the microstructure of as-received CVD Ni and the microstructural evolution of CVD Ni upon annealing was first carried out using optical microscopy. The samples were cut into dimensions about 5mm X 10mm X 8mm. Both planar and transverse directions can be observed on the same sample. Each sample was first annealed in a box furnace for 45 minutes, and then taken out to cool down to room temperature in the air. After that, each sample was first grinded with abrasive papers from a grit of 120, followed by 320, 400, 600, 800 and Finally the sample was polished on a soft cloth with 1µm diamond suspension solution. A copper sulphate based etchant was used to etch the surface in order to reveal the grain boundaries for observation under optical microscope.

51 2.3.2 Grain boundary tracking via semi in-situ optical microscopy In order to track the grain boundaries migration at the beginning of grain growth, the same sample was used for observation. The sample surface was dented by microhardenss machine to mark the location. The same sample was annealed at increasing temperature from 25 C to 250 C, 370 C and 480 C, each for 45 minutes. After each annealing, the sample was taken out from the furnace and observed under optical microscope for any microstructure change Transmission electron microscopy (TEM) After the preliminary analysis using optical microscopy, Transmission Electron Microscopy (TEM) was carried out using FEI Tecnai F20 TEM. In order to prepare the special TEM sample, thin samples measuring 1cm thickness were cut from the bulk nickel. Since the microstructure at different temperatures will be observed, the samples were first annealed in a box furnace and taken out to cool down. Afterwards, they were first grinded on both sides using abrasive paper from 120 grit to 800 grit and ended up in a thin foil with approximately µm thickness. A TEM puncher was used to cut into several pieces of disc-shaped samples with 3mm in diameter. These disks were then electropolished in a solution of 15% perchloric acid and 85% ethanol using double jet polishing machine. The well polished sample has a tiny hole in the centre with certain thinned area that is ready to be observed under TEM Hot stage In-situ TEM Hot stage in-situ TEM was carried out to trace the microstructure evolution back to initial state. The TEM sample prepared by the method mentioned above was placed in a heated chamber and the microstructure was observed constantly. The temperature of the chamber was manually controlled thereby the heating rate was not consistent throughout the annealing. To avoid from the effect of non-uniform heating rate, the samples were specifically held at high temperatures for a long time. One microstructural feature was monitored on each sample, and the microstructure was recorded every 2-3 minutes. 39

52 Microhardness and Rockwell hardness testing Since the microstructure changed, the corresponding mechanical properties were expected to change accordingly. Therefore, the evolution of mechanical property was measured by micro hardness and Rockwell hardness testing on as-received sample and samples annealed at 200 C, 400 C, 600 C and 800 C. The hardness value from both planar and transverse directions was obtained. In Vickers microhardness testing, the sample was well polished and etched to the extent that was observable under optical microscopy. The load used was 2.942N, HV0.3 and hold for 10 seconds. In the Rockwell hardness, Scale B indenter made of 1/16 steel ball was used. Each hardness value was obtained based on the average of at least 10 tests.

53 Heat Flow W/g 41 3 Results 3.1 Thermal stability investigation using DSC 0.05 DSC st run 2nd run difference Curie temperature: 354 º C Temperature ºC Figure 3.1 Thermal stability investigation of CVD nickel through Differential Scanning Calorimetry at a heating rate of 20 C/min. The blue line shows the first heating cycle from room temperature up to 600ºC and the red line shows the consecutive run of the same sample after it was completely cooled. The green line is the difference in heat flow obtained by subtracting the second run from the first run. Several DSC scans were carried out on a series of samples, all of which showed exothermic heat release. However, no apparent heat release peak was detected from these DSC scans results. The fluctuating noise in Figure 3.1 resulted from the background drift of the DSC machine and was found to be comparable to the actual heat flow rate measured. It can be seen that even at the maximum point of the curve, the flow rate is only about twice the magnitude of the fluctuating noise range. Therefore, based upon all the testing results obtained from different samples, conclusion could be drawn that DSC curve of CVD nickel, unlike the electrodeposited nanocrystalline nickel that has a distinct exothermic peak and large amount of heat release, did not show apparent exothermic peak. Besides, rather small enthalpy can be detected from

54 42 microstructure change of CVD nickel. More DSC scans at heating rate of 10ºC/min and 40ºC/min can be found in Appendix Preliminary analysis using Optical Microscopy Characteristics of as-received CVD nickel The microstructure of as-deposited CVD nickel was quite inhomogeneous as shown in Figure 3.2. Ultrafine and nano grains were found embedded in between the large columnar grains that grow in the direction perpendicular to the substrate (Figure 3.2 (b)). The cross sectional dimension of the large columnar grains is shown in Figure 3.2 (a) in the planar direction parallel to the substrate. As can be seen, the large grains were asymmetrical with two narrow ends, as similar to elongated rhombuses. The width of the columnar grains as shown in the transverse direction was in the range of 2-12μm. In the optical micrographs, the size and shape of the small grains matrix were quite difficult to identify due to the limitation of the microscope resolution. (a)

55 43 (b) Figure 3.2 Optical micrographs showing the microstructure of as-received CVD nickel (a) planar direction (b) transverse direction. Note: the red arrow indicates the growth direction of the columnar grains Microstructure evolution upon annealing The microstructure evolution of CVD nickel upon annealing at different temperatures in the planer view is shown in Figure 3.3. Initially the material had a bi-modal grain size distribution. There was no obvious change in grain morphology observed in the sample annealed at 200ºC for 40 minutes. In the sample annealed at 400ºC, major change was found in the grain morphology as all the ultrafine and nano grains grew into size identifiable under optical microscope. The large grains also grew with two ends blunted. At 600ºC, grain growth continued as both the large and small grains grew towards equiaxed shape. The grain size distribution became much narrower compared to the initial bi-modal grain size distribution in as-deposited samples. Twin boundaries were clearly seen in most of the large grains. At 800ºC, all the grains grew uniformly into equiaxed shape with straight grain boundaries. The twin boundaries were still visible with twin width in the micrometer range. After annealing at 900ºC, the average grain size was estimated to be 8-12μm. The bi-modal grain size distribution did not exist anymore.

56 44 Figure 3.3 Grain morphology of CVD nickel at planar view after isothermal annealing for 45 minutes at various temperatures: 200ºC, 400ºC, 600ºC, 800ºC and 900ºC. The microstructure evolution upon annealing observed in the transverse direction showed consistent results with the planar direction, as shown in Figure 3.4. Grain growth behaviour in CVD nickel was not observed until the sample was annealed at 400ºC. The ultrafine and nano grains grew continuously as can be seen in the samples annealed at 400ºC, 600ºC and 800ºC. The grain boundaries of the large columnar grains were no longer straight as found in the asdeposited samples but rather curved at various locations and merged with the surrounding small grains. Growth twins were clearly seen inside the large columnar grains aligned parallel with the long axis of the grains.

57 45 Figure 3.4 Grain morphology of CVD nickel at transverse view after isothermal annealing for 45 minutes at 200ºC, 400ºC, 600ºC and 800ºC, respectively. In order to narrow down the temperature range in which grain growth starts to occur, further investigation was carried out on samples annealed at 250ºC, 300ºC and 370ºC for 45 minutes. The results can be found in Appendix 7.2. The optical micrographs showed that grain growth did not start at 250ºC and 300ºC whereas it indeed occurred at 370ºC. This results indicated that the onset temperature for grain growth of CVD nickel lies somewhere between 300ºC and 370ºC upon annealing Grain boundary tracking by Semi in-situ OM The same as-deposited sample was first annealed at 250ºC followed by 370ºC annealing; no grain growth was detected under either condition. As compared with the sample annealed directly at 370ºC, apparent delay of grain growth was evident. After 480ºC annealing, grain

58 46 growth occurred and the microstructure at the same location was compared between sample annealed at 370ºC and 480ºC as shown in Figure 3.5. Site #1: Site #2: Figure 3.5 Semi in-situ tracking of grain growth under optical microscope. The comparison was made between the same sample annealed at 370ºC for 45 minutes and at 480ºC for the same amount of time. The letter M indicates the indent marked by microhardness machine. As indicated in the multi-coloured circles, the ultrafine and nano grains in some regions seem to have very similar shape and orientation and these grains tended to merge into one large irregular grain. For example, in the red circles at site #1, all the small grains between the two columnar

59 47 grains grew into one large grain lying in between the two columnar grains. Similar examples exist at many other locations in the sample used for grain boundary tracking. On the other hand, some small grains only grew into slightly larger grains with equiaxed shape. This indicates that abnormal grain growth occurred before uniform grain growth between 370ºC and 480ºC. 3.3 Detailed analysis using Transmission Electron Microscopy Characteristics of as-deposited CVD nickel Transmission Electron Microscope (TEM) was employed for further investigation at higher magnifications. Using TEM, the as-deposited CVD Ni was characterized by three distinct features worth to be noted. They were ultrafine and nano grains, large columnar grains with high density of ultrafine and nano twins and multiply twinned particles (MTP s) with five-fold symmetry. (i) Ultrafine and nano grains As mentioned in Section 3.2.1, the microstructure of as-deposited CVD nickel was bi-modal consisting of large columnar grains embedded in the ultrafine and nano grain matrix. The TEM micrograph of these small grains is present in Figure 3.6.

60 48 Figure 3.6 Ultrafine and nano grains surrounding the large columnar grains in as-deposited CVD nickel in the transverse direction.

61 49 (ii) Large columnar grains with high density of growth twins The morphology of large columnar grains in both planar and transverse directions is shown in Figure 3.7 (a), (b). Growth twins tended to grow in the same direction with the longitudinal axis of the large grains. Among them, there were some special grains with high aspect ratio that were sectioned by multiple black lines perpendicular to the growth twin boundaries, as shown in Figure 3.7 (c), (d). Upon further investigation at higher magnifications as shown in Appendix 7.3, these lines were confirmed to be dislocation wall structure comprised of tangled dislocations. It was also found that the twins from the two sides of the dislocation walls were slightly displaced with respect to each other. In other word, the continuity of growth twins inside these large grains was interrupted by tangled dislocation walls. Planar Direction Transverse Direction (a) (b)

62 50 (c) (d) Figure 3.7 TEM micrographs showing the morphology of large columnar grains in as-deposited CVD nickel, (a) planar direction, (b) transverse direction, (c) and (d) special large grains with dislocation wall structure traversing twin boundaries across the grains. As can be seen in Figure 3.8, high density of growth twins was found inside the large columnar grains. The thickness of the twins, i.e., the spacing of twin boundaries, was obtained based upon measurement of 200+ twins. It was found that all the twins had thickness below 500nm with 83.78% of them being nano twins with thickness below 100nm. According to the thickness measurement, these twins are categorized as the ultrafine and nano twins.

63 Distribution (%) 51 (a) Twin thickness (nm) 83.78% 16.22% Tt < 100nm 100nm < Tt < 500nm (b) Figure 3.8 (a) High density of growth twins inside the large columnar grains. (b) Distribution of twin thickness showing all the twins are in the ultrafine and nano size range.

64 52 (iii) The five-fold twinned structure In addition to the small grain matrix and large columnar grains observed in as-deposited CVD nickel, there was another unusual grain morphology observed, the five-fold twinned structure as shown in Figure 3.9. In the structure, five subgrains joined together to form a pentagon. The size range of these structures was estimated to be 1.2~5.6µm. It is worth noting that five-fold twinned structure with such big size has never been reported in other bulk materials. More five-fold twinned structures in the as-deposited CVD Ni can be found in Appendix 7.4. (a) (b) Figure 3.9 TEM images showing five-fold twinned structures embedded in the bi-modal grains of as-deposited CVD nickel (a) with extra nano twins to accommodate the misfit, (b) without extra twin boundaries.

65 Microstructure evolution upon annealing Grain growth behavior TEM study of microstructure evolution of CVD Ni clearly showed that at 400ºC, all the small grains grew bigger compared to the as-deposited sample. The density of growth twins did not change apparently in the sample annealed at 400ºC. However, upon further annealing at 600ºC, great reduction in twin density was observed along with uniform grain growth; dislocation cell structures were found in some grains. At 800ºC, even less twins with wider thickness were observed. Besides, dislocation cell structures were all over in the material.

66 54 Figure 3.10 Microstructure evolution of CVD nickel upon annealing at 200ºC, 400ºC, 600ºC and 800ºC in the planar direction observed under TEM. The microstructure evolution of CVD Ni upon annealing in the transverse direction was similar to the planar direction in terms of grain growth behaviour. Besides, grain boundaries of the large columnar grains were curved at multiple locations, as evident in Figure ºC, 600ºC and 800ºC samples. This suggested the active interaction of the large columnar grains with the surrounding small grains through grain boundary migration.

67 55 Figure 3.11 Microstructure evolution of CVD nickel upon annealing at 200ºC, 400ºC, 600ºC and 800ºC in the transverse direction observed under TEM.

68 Detwinning process in the large grains When observed at higher magnification, the twin receding was found to occur at 400ºC and higher temperatures. In the following sections, the detailed detwinning process was captured and studied in samples annealed at 400ºC, 600ºC and 800ºC for 45 mins. More TEM images on the detwinning process can be found in Appendix (i) 400ºC Figure 3.12 showed that at 400ºC, some of the growth twins receded in the form of step movement. And dislocation arrays were often seen at the migrating step front of these twins. In addition, detwinning by the gliding of dislocation arrays was also observed in the five-fold twinned structure as shown in Figure Figure 3.12 (a) Twin receding in the form of step movement (b) Dislocation configuration at the migrating step fronts

69 57 Figure 3.13 (a) Receding of growth twins took place in a five-fold twinned grain; (b) higher magnification of (a) showing how dislocation arrays were involved in the detwinning process. Figure 3.15 Partial dislocations formed at both the end of twins and the misfitted sites traversed by dislocation lines.

70 58 Figure 3.15 showed detwinning process in the special large grains with dislocation walls traversing growth twins. It can be seen that the dislocation arrays not only formed at the ends of the twins, but also formed at the sites of these dislocation walls. Figure 3.16 Tangled dislocations along twin boundaries with some of them extending into the lattice Figure 3.17 Twin boundaries transformed into dislocation array structures.

71 59 It can be seen from Figure 3.16 that the twin boundaries were no longer stable at 400ºC while some of them turned into thick lines full of tangled dislocations. Some other twin boundaries transformed into low angle boundaries comprised of dislocation arrays as shown in Figure (a) (b) Figure 3.18 (a) Twin boundaries turned into dislocation lines and disappeared (b) dislocations emitted from twin boundaries into the twinned crystal Some twin boundaries became single dislocation lines and disappeared locally as can be seen in Figure 3.18 (a). On the other hand, dislocations were seen to emit into the twinned lamella from the twin boundaries as shown in Figure 3.18 (b). It is therefore very likely that before twin boundaries disappeared, the twinned region had been converted back to matrix stacking by those dislocations emitted from the twin boundaries.

72 60 (ii) 600ºC From the TEM micrographs (Figure 3.19 and 3.20) of CVD nickel after annealing at 600ºC, it can be seen that after growth twins were reverted back into matrix lattice, dislocation cell structures formed at the twin-free region. The cell walls were comprised of tangled dislocations with the interior volume free of dislocations. The size of these cells appeared to be non-uniform. Besides, the dislocation cells in Figure 3.19b were seen to not only form at the twin-free region but also form inside the twin lamellae across twin boundaries. The twin boundaries at 600ºC became much darker and blurred compared to the samples annealed at 400ºC. High density of dislocations nucleated and tangled on the twin boundaries. Twin boundaries were found to be the precursor of dislocation cell structure. (a) (b) Figure 3.19 Dislocation cell structures formed in the cleared grain interior after detwinning; twin boundaries became darker and blurred full of tangled dislocations.

73 61 Precursor of dislocation cell walls at TBs (a) (b) Figure 3.20 (a) One grain with half of the growth twins transformed into dislocation cells; (b) Dislocation cells were formed both at the detwinning front and directly at twin boundaries. As shown in Figure 3.20, newly formed cells usually had some portions of the walls either located at the moving step of a twin or at halfway of an activated twin boundary. The dislocations at these locations emitted into the lattice, tangled with the adjacent dislocations and formed cell walls. In addition, as shown in Figure 3.20, at the vicinity of the detwinning front indicated by the yellow arrows, high density of dislocations were randomly distributed. This revealed the initial state of dislocation structure right after detwinning and before they regrouped into cellular structure. Thus, the transition from random dislocation structure to organized cell structure might be another possible cell forming mechanism.

74 62 (a) (b) Figure 3.21 (a) Dislocation arrays formed at the end of twins and along twin boundaries (b) detwinning by migration of dislocation arrays and dislocation emission Dislocation arrays were observed to form both at the moving ends of twin lamellae and along twin boundaries as shown in Figure 3.21 (a). Is is worth to be noted that the twins in Figure 3.21 no longer receded in the form of step movement but receded the whole structure all at once. And the thickness of these twins were around 100 nm. Thus the twin thickness might be a critical factor determining the way of twin receding by dislocation migration. In addition to detwinning through massive gliding of dislocation arrays, the twin boundaies were also seen to disappear by dislocation emission, as evidence in Figure 3.21 (b). Figure 3.22 showed the detwinning phenomenon in the large grains traversed by dislocation walls. It was found that detwinning could start from both ends of twins at the intersection of twin boundaries and grain boundaries. Besides, detwinning can also start from the sites where

75 63 dislocation walls intersect with twin boundaries. Dislocation arrays were seen to form at these locations and some twins already receded away from the intersecting sites. Figure 3.22 Morphology of nanotwins inside the large grains traversed by dislocation walls after annealing at 600 ºC.

76 64 (iii) 800ºC The samples annealed at 800 ºC showed similar morphology as the samples annealed at 600ºC inside the large grains. Nevertheless, the twins became much wider with much lower twin density. Dislocation cells were found everywhere both in the cleared grain interior and inside the twin lamellae. Most of the dislocation cells had cleared cell interiors. Non-uniform cell size still existed in the 800ºC sample as shown in Figure 3.23 (b). Figure 3.23 (a) Dislocation cell walls formed at the moving steps of a twin lamella (b) Nonuniform size distribution of dislocation cells at 800 ºC Figure 3.24 clearly showed the traces of dislocations remained from detwinning process turned into part of a dislocation cell wall observed in the transverse direction.

77 65 Figure 3.24 Schematic demonstration of the formation of dislocation cells at the tip of a receding twin lamella. Traces of dislocations left from detwinning became part of a dislocation cell wall Five-fold twinned structures The five-fold twinned structures remained in the material after annealing at 200ºC, 400ºC and 600ºC. However, when annealed at 800ºC such structure no longer existed in the material. The morphology of the five-fold twinned structures in as-deposited, 400ºC and 600ºC annealed samples is shown in Figure Since no apparent microstructure change was observed at 200ºC, the five-fold structure at 200ºC was exempted for comparison.

78 66 25 ºC 400 ºC

79 ºC Figure 3.25 Five-fold twinned structures in CVD Ni of as-deposited condition and after annealing at 400ºC and 600ºC In the as-deposited CVD Ni, five-fold twinned structures were frequently observed. It can be seen that after 400ºC annealing, the grain boundaries of the structure became curved and smoother, indicating that grain boundary migration had occurred as a result of grain growth in the material. The growth twins inside the structure were also subject to detwinning process similar to the twins inside the large columnar grains. The five-fold twin boundaries remained stable at 400ºC. Yet when the temperature further increased to 600ºC, the five-fold twin boundaries turned into thicker lines possibly with many dislocations. The outer boundaries also became rounded. Some of the structures consisted of dislocation cells in their subgrains. Nevertheless, five-fold twinned structures were seen to totally disappear at 800ºC.

80 Hot stage In-situ TEM As shown in Figure 3.26, upon anisothermal annealing in the hot stage TEM, the ultrafine and nano grain pocket in the as-deposited sample remained stable until 375ºC. The grain interiors were seen to clear up at 375ºC and higher temperature. At 600ºC, these grains grew bigger with their boundaries getting smoother and clearly seen. Nevertheless, the grain growth behaviour of the small grains was only confined in the original area between the large grains. Grain boundary migration of the large grains was not observed even after the sample was held at 661ºC for 126 minutes. Furthermore, detwinning process and dislocation motion were not observed elsewhere. The results described above are very different from isothermally annealed bulk samples. In other words, the large grains along with the high density of growth twins remained stable upon annealing up to 661ºC for 126 minutes. (to be continued)

81 69 Figure 3.26 In-situ TEM images showing the microstructure of the same location upon annealing up to 661ºC. The time specified was accumulated from 600ºC and 661ºC, respectively. Another in-situ monitoring of microstructure evolution was done by hot-stage TEM specially focusing on the five-fold twinned structure in Figure No apparent change was observed on the structure. The dislocations initially existed in the sample remained in place regardless of the high temperatures. In-situ observation was also made on another location, which can be found in Appendix 7.6.

82 70 Figure 3.27 Morphology of a five-fold twinned grain upon annealing from room temperature up to 655ºC in hot stage TEM 3.5 Mechanical properties of CVD nickel upon annealing Both Figure 3.28 and Figure 3.29 showed that starting from 400ºC, both the microhardness and Rockwell hardness of CVD Ni decreased with increasing temperature. This result is in good agreement with the observation on microstructure change mentioned in previous sections, since hardness decreases accordingly with increasing grain size.

83 HRB HV Vickers Microhardness planar 174 transverse Annealing temperature (ºC) Figure 3.28 Microhardness of CVD nickel on planar and transverse direction upon annealing Rockwell hardness Annealing temperature (ºC) planar transverse Figure 3.29 Rockwell hardness of CVD nickel on planar and transverse direction upon annealing

84 72 4 Discussions 4.1 Grain growth behaviour of CVD nickel upon annealing The DSC curve in Figure 3.1 showed that the microstructure change of CVD nickel upon heat treatment is a gradual process with relatively small enthalpy. It was introduced in Section that for nanocrystalline Ni major exothermic peak was detected on the DSC curve. This is not surprising because of the following two reasons. First of all, CVD Ni has bimodal grains structure and only about half of its volume is comprised of the nano-grains. Thus the average grain size is much bigger than the 100% nanocrystalline nickel leading to lower intercrystalline energy in the material, which is the main driving force for grain growth. In addition, the coherent twin boundaries inside the large grains are known to have much lower energy than grain boundaries. Therefore, lower energy of the system and less tendency to reduce the intercrystalline volume resulted in the small amount of heat release and the absence of nucleation process that is frequently observed in nanocrystalline Ni. The preliminary microstructure evolution results from both optical microscopy and TEM showed that grain growth started to occur at 400ºC and higher temperature. Further investigation on samples annealed from 250ºC to 370ºC suggested that the onset temperature for grain growth of CVD nickel lies somewhere between 300ºC and 370ºC. The results from semi in-situ OM observation showed that some small grains lying in between the columnar grains merged into one larger grain with irregular shape before 480 ºC. It is therefore very likely that abnormal grain growth occurred before curvature driven normal grain growth proceeded. Abnormal grain growth occurs when a population of grains grows faster than the rest by consuming the matrix of grains leading to inhomogeneous microstructure. 63 Anisotropy of grain boundary energy and mobility is believed to be the main cause of abnormal grain growth. 64 It is therefore proposed that the initial grain growth of CVD Ni was accomplished by grain rotation and coalescence, similar to the subgrain rotation and coalescence model in the conventional recovery and recrystallization as explained in Section In nanocrystalline nickel with 20nm grain size, it has been suggested that the rotation and coalescence type mechanism might be responsible for the initial abnormal grain growth observed. 34 Recent MD simulation also suggested that the grain rotation induced coalescence and

85 73 curvature driven grain migration might occur in those FCC metals with very small starting grain size. 82 It is known that materials fabricated from deposition process usually have texture. For instance, the electrodeposited nanocrystalline Ni from watts baths with saccharin shows a strong [200] texture. 35 Hence, the nano grains only have to slightly rotate to join each other after which large grains can be formed. The low angle boundaries between the small grains are then removed through climbing of edge dislocations. 34 This is in consequence of structural relaxation of the grain boundaries upon annealing as suggested by Rupp and Birringer for nanocrystalline Cu and Pd. 65 A proposed physical model for the deformation induced grain rotation and coalescence is shown in Figure 4.1. The only difference is that when subject to external stress the rotation is accomplished by dislocation sliding at grain boundaries, while during annealing, the diffusion process along grain boundaries leads to rotation. The shape of the grown grain is not spherical but depends on the size and shape of the nano grains involved in the coalescence process. After all the grains have recovered, the remaining grains now have different orientations with high angle grain boundaries. The dominating mechanism will then become grain boundary migration leading to uniform grain growth. In CVD Ni, it was evident that curved boundaries existed in the samples annealed at 600ºC and 800ºC, indicating the occurrence of grain boundary migration.

86 74 Figure 4.1 Model of grain rotation and growth induced by deformation in nanocrystalline Ni (a) before deformation (b) shear of grain 1 and 2 by dislocation sliding at grain boundaries and subsequent rotation of grain 3(c) multiple grain rotation (d) a large grain formed with subgrain boundaries. 66 In a word, initially some of the ultrafine and nano grains with similar orientation combined to form large grains by rotation and coalescence. Afterwards, normal grain growth took place where all the grains grow uniformly with curvature driven grain boundary migration. Abnormal grain growth with faceted grain boundaries in the stage III grain growth observed in electrodeposited nanocrystalline nickel (Section 1.3.3) was not observed in CVD nickel. It is believed that thermodynamically the grain growth in CVD nickel was associated with the reduction of the free energy of grain boundaries and triple junctions. During the chemical vapour deposition process, high energy was introduced into the materials through grain boundaries, twin boundaries and other forms of defects such as dislocations and point defects, all of which made the materials thermally unstable. As such, the system tends to release the excess free energy when it is either thermally or mechanically activated. 4.2 Thermal stability of ultrafine and nano twins in CVD nickel In general, twin boundaries are considered stable structures compared to grain boundaries. In the case of nickel, a coherent twin boundary possesses free energy of 43 mj/m 2 while grain boundary has free energy of 866 mj/m Twinning can be a deformation mechanism by means of forming

87 75 deformation twins under stress. The presence of twins can also play an important role in strengthening the materials similar as the effect of grain boundaries, where the boundaries act as barriers to impede the dislocation motion. Recent researches on growth twins of FCC nanocrystalline materials demonstrated that detwinning can also be the dominating deformation mechanism at the early stage of plastic straining. 45, 58, 68 To the author s knowledge, detwinning due to thermal activation in FCC materials has never been publically reported elsewhere. According to the TEM micrographs of samples annealed at different temperatures, growth twins in the large columnar grains were found to start receding at 400 ºC. MD simulation by Hu et al suggested that migration of twin boundaries under stress first needed nucleation of twinning partial dislocations. 68 In the case of CVD Ni, arrays of twinning partials were observed at the migrating twin ends, or referred to as the incoherent twin boundaries (ITBs) as introduced in Section Hence, it is suggested that twin receding was accomplished by gliding of these twinning partials collectively at ITBs. For thick twins ( > 100nm), ITBs migration occurred in the form of step movement; whereas in the case of thin twins with thickness round or below 100nm, gliding of partial dislocations reoriented the whole twinned crystals back to lattice stacking all at once and removed the two twin boundaries simultaneously. This is in accordance with the detwinning models proposed by Wang et al for nanotwins with both ends at grain boundaries. 58 The occurrence of different types of receding modes according to twin thickness was probably due to the higher friction stress incurred during dislocation migration at wider ITBs in those thick twins. Most of the twins receded from their two ends intersecting with grain boundaries and twinning partials were believed to nucleate at grain boundaries. In addition to that, dislocation arrays also formed at the sites where dislocation walls traverse the twin boundaries across some narrow and large grains as shown in Figure Thus, it seems that the twinning partials involved in the detwinning process preferred to not only nucleate at the grain boundaries, but also form at the sites where twin boundaries were traversed by dislocation walls. MD simulation by Hu et al suggested that nucleation of twinning partials was heterogeneous such as at stress concentration sites, at twin boundaries or at the surface of the crystal. 68 Likewise, twinning partials in CVD Ni seem to prefer nucleation at highly faulted regions, such as grain boundaries and dislocation walls.

88 76 Apart from detwinning through collective glide of partial dislocations, the twin boundaries were also seen to disappear through dislocation emission at the same time. Before twin boundaries could disappear by themselves, the twinned region must have been converted back to the matrix stacking. Therefore, it is very likely that this conversion was done by the movement of dislocations emitted from the twin boundaries. The two possible detwinning mechanisms based upon the observation of the microstructure evolution upon annealing are summarized in the following chart (Figure 4.2). Detwinning mechanisms 1 2 Collective glide of partial dislocations at twin ends Local dislocation movement Thick twins (> 100nm): detwinning in the form of step movement Nano twins (< 100nm): whole twin receded together 1. Twinned region converted back to matrix stacking by dislocation movement 2.Twin boundaries disappeared by dislocation emission Figure 4.2 Proposed possible detwinning mechanisms for CVD nickel upon annealing Upon thermal activation, twin boundaries became very unstable as evident in Figure 3.16, 3.17, 3.18 and Dislocations arrays and densely tangled dislocations were activated along the twin boundaries. Some of the dislocations bulged from the twin boundaries and extended into the lattice, while some others emitted into the twinned lamellae. The tendency of transformation of twin boundaries into dislocation configurations upon annealing suggested an energy reduction process. Our results suggested that migration of ITBs and CTBs under thermal activation first needed nucleation of twinning partial dislocations, which was a result of diffusion and dislocation climbing. Considering the formation of annealing twins was reported to be associated with the

89 77 migrating grain boundary during grain growth 69, it is possible that migrating grain boundaries is also a necessary requirement for the nucleation of twinning partials involved in detwinning process. However, the latter statement needs to be confirmed by further investigation in the future. At 600 ºC and higher temperature, the dislocations left from detwinning process reorganized themselves into cellular structure with the boundaries comprised of tangled dislocations, as shown in Figure 3.19, 3.20 and These boundaries should be low angle boundaries with small misorientation because very little contrast can be seen on any adjacent dislocation cell units. The rearrangement process from randomly distributed dislocations into organized cellular structure with well-defined boundaries is similar to the recovery process observed in heavily deformed metals as described in Section Since high density of randomly distributed dislocation is considered unstable in terms of stress field screening, inhomogeneous dislocation distribution will form spontaneously, leaving the major fraction of the grain free of dislocations. 70 The driving force for this recovery process is therefore to minimize the total elastic energy associated with the high density of dislocations. In addition to self re-organizing from randomly distributed dislocations, dislocation cells were also seen to form part of the cell walls directly at the twin boundaries or migrating fronts of ITBs where dislocations were active. In order to explain the transformation from twin boundaries to dislocation configurations in CVD Ni upon annealing, the energy involved in this process was considered. Work of Luo et al 71 showed that in the electroformed copper where high density of growth twins were present, the density of the twin boundary area was around 1.8~8 x10 4 cm 2 /cm 3. As can be seen in Figure 4.3, the total energy of these twin boundaries is equivalent to the dislocation with density as high as m -2, which is equivalent to 10%~20% plastic strain applied to the polycrystalline copper. On the other hand, the density of the twin boundaries area in CVD nickel was calculated to be 6.8 x 10 4 ~1.8 x 10 5 cm 2 /cm 3, which was even higher than that of electroformed copper. Therefore, although coherent twin boundaries has much lower energy than grain boundaries and they are more stable against migration than grain boundaries, ultrafine and nano twins with such high density are still very unstable structure thermodynamically. They tend to transform into dislocation cell structures with lower free energy through a series of process described above. As

90 78 a whole, the detwinning process in CVD Ni upon annealing was driven by the large amount of energy stored in the high density of twin boundaries. Figure 4.3 Strain energy density of twin boundaries and dislocations in polycrystalline copper Thermal stability of five-fold twinned structures Five-fold twins have been studied for more than 50 years yet the formation mechanism of fivefold twins in nanocrystalline FCC metals is still unclear. 72 Two requirements were proposed to form five-fold twins, the external stress and the orientation variation of the stress. 73 MD simulation showed that the external stress was not necessary when the fivefold twins were formed in nano-grained Cu of 5~30nm. 74 This was further confirmed by Huang et al. 72 through experimental study on nanocrystalline Cu upon annealing. All the above findings suggested that five-fold twins, similar to the conventional twins, could be formed as a result of deformation, annealing and as in our case, during film growth. Such structure should be a stable configuration with relatively low energy. Interestingly, the five-fold twins in FCC materials reported so far were all in the early stage of grain growth. Due to the geometrical arrangement of FCC structure, five-fold twinned grains should have distortion along the central axis. As the five-fold twins grow, internal stress will

91 79 build up due to the distortion. The maximum size of the structure reported without stress was 40nm. 75 Almost all the work so far was devoted to the five-fold twins in nanowires, nanoparticles and thin films; yet very little discussion was carried out in bulk materials. In our case, the fivefold twinned structures do exist in bulk nickel with thickness of 16mm. The estimated size of these five-fold twinned grains ranges from 1.2 to 5.6 µm, which is much bigger than any literature reported previously. Wang et al. have performed detailed analysis on the internal structure of five-fold twinned grains in CVD Ni by High-resolution TEM. 76 They found that the mismatch inside the structure was shared irregularly among the five twin boundaries (Figure 4.4) and along the grain boundaries (Figure 4.5) through stacking faults, nano twins and disclinations. Disclinations are wedge-like defects due to angular distortion formed by irregularly packing of atoms. Stress field and elastic distortion usually exists in the region of disclinations. The disclination approach has been used to effectively explain the unique structure of pentagonal particles. 77 Figure 4.4 HREM image of the central area of the five-fold twinned grain. The twin boundaries are indicated by TB1 to TB5. 76

92 80 2 nm 2 nm Figure 4.5 HREM images of the boundary between the matrix and five-fold twinned structure with (a) nanotwins and stacking faults (b) disclinations and stacking faults 76 According to our experimental results on the evolution of five-fold twinned structure, it was found that the structure was gradually eaten by the surrounding grains through grain boundary migration. At 600 ºC, the five-fold grains became rounded without showing the star-tips. This indicated that under thermal activation the internal stress developed at the grain boundary during growth of nickel deposit was relaxed and disclinations also disappeared. The five-fold twin boundaries also became unstable and full of dislocations along the boundaries. Moreover, such structure totally disappeared in the sample annealed at 800 ºC, suggesting the five-fold twinned structure was very unstable at high temperature and would be totally consumed by surrounding growing grains at 800 ºC. Nevertheless, when compared with the conventional growth twins existed in CVD Ni, such multiply twinned structure is still much more stable against thermal activation because detwinning process for those growth twins was found to occur in the material as early as 400 ºC.

93 Analysis on the stable structure in hot stage in-situ TEM Hot stage TEM was used to in-situly monitor the grain boundary migration of CVD Ni upon annealing. The results from tracing at several locations showed that although the ultrafine and nano grain pocket started to grow at 375 ºC, the large grains were always stable regardless of the high temperature and long time provided. Therefore, absence of grain boundary migration due to insufficient conditions that might have prohibited diffusion process was excluded from the possible causes of this phenomenon. It is known that the preparation of TEM thin foil introduces artificial surfaces to the sample. It is very likely that the thin foil effect was responsible for the stable boundaries since such stabilizing phenomenon was not observed in the conventional TEM with bulk samples. And the free surfaces of TEM thin foil have been reported to actually inhibit the grain boundary migration of nanocrystalline nickel during in-situ TEM annealing experiment. 78 The nano grains in regions directly surrounding the jet-polished hole appeared much smaller than the grains far from the hole on the 3mm disk where uniform grain growth took place. This indicated a free surface pinning effect and this effect is apparently in close relation with the thickness of the foil. Novikov has simulated the growth of films on substrate where all the grains were in columnar shape with only the diameters different. 79 According to his work, the grain growth process was inhibited by a dragging force that was further influenced by the film thickness. The ratio of the initial grain diameter to the film thickness was a key parameter in governing the grain growth character. When the ratio was 3, the grain boundary was completely stagnated. If the ratio was between 3 and 0.3 abnormal grain growth occurred while normal grain growth took place at ratio < 0.3. Mullins has established a well developed theory to explain the stagnated normal grain growth in thin films when the average grain size was comparable to the film thickness. 80 He attributed the stagnation effect to the grain boundary grooving at the free surface of the thin film (Figure 4.6). The grain boundary grooves form at the intersection of grain boundary with the free surface of the films through surface diffusion. By redistributing matters at the grain boundary, equilibrium of the surface tension is established. Consider a grain boundary running through the thickness of the film perpendicular to the surface, grooves will form on both the top and bottom surface of the film. The grain boundry would need to increase its area in order to escape from the

94 current position and proceed further, which is not energetically favored. Mullins model showed that such boundary was actually pinned effectively by the surface. 82 Figure 4.6 Geometry of a grain boundary groove 80 According to the microstructure observation of as-deposited CVD nickel, columnar grains located in the TEM sample near the jet-polished hole must run through the thickness. This is because the thinned region near the hole might be only one or two nanocrystalline grains thick, whereas the height of the columnar grains measures more than 50 µm. The diameter of the columnar grains was estimated to be 2~12µm. Hence, the ratio of grain diameter to the film thickness is much higher than 3, above which complete stagnation of grain boundaries would occur as mentioned by Novikov. With such a high ratio and both ends of the columnar grains pinned at the top and bottom surface, it is therefore not surprising that the large columnar grains remained stable during annealing under in-situ TEM. As observed in the in-situ TEM, not only the grain boundaries of the large columnar grains remained stable, but the twin boundaries were also stable. No evidence of detwinning process was observed even at very high temperature (~752 ºC). If the reason mentioned above stands, then our results further suggested that the free surface pinning effect in thin film was also applicable to twin boundaries in addition to grain boundaries. Another possible reason leading to the absence of detwinning could be the release of residual stress when the TEM thin foil was thinned to around 80 µm. It is generally accepted that films synthesized by deposition techniques usually have residual stress generated during film growth. It is possible that detwinning process and grain boundary migration in the large grains observed under conventional TEM was triggered by the residual stress in the bulk samples. Whereas the

95 83 thin foil sample in the case of in-situ TEM had almost no residual stress left and thus could not trigger the detwinning and grain growth process. 4.5 Thermal stability of CVD nickel in terms of mechanical properties As shown in Figure 3.28 and 3.29, hardness remained high after 200 ºC annealing. This is in consistent with the microstructure results obtained from both OM and TEM, where no evidence of grain growth was observed for the sample annealed at 200 ºC. Upon further annealing at higher temperatures, the hardness of CVD nickel showed a decreasing trend. According to the microstructure investigations by OM and TEM, grain growth started to occur at 400 ºC annealing for 45 mins and the grains continued to grow and finally reached uniform grain size distribution at 800 ºC annealing for 45 mins. The decrease in hardness with increase in grain size followed conventional Hall-Petch relation. Hardness reflects a material s ability to resist plastic deformation and dislocation movement. Therefore, it is an important parameter that is closely related to the yield strength of a material. While the grain boundary acts as a barrier to dislocation movement, twin boundaries are considered to have the same prohibiting effect. 81 In the as-deposited CVD nickel, the large amount of twin boundary area in the columnar grains together with the large grain boundary area of nano grains all contributed to its superior strength. By that means, upon thermal activation, the reduction of twin boundary and grain boundary area led to the softening of CVD nickel.

96 84 5 Conclusions In the present study, the microstructure evolution of CVD Ni upon annealing was investigated systematically. A series of techniques was employed to study its grain growth behaviour including DSC, optical microscopy, TEM, semi in-situ optical microscopy and hot stage in-situ TEM. The evolution of mechanical properties upon annealing was investigated by microhardness and Rockwell hardness testing. Through the above experimental investigations, distinct microstructural features were observed in the as-deposited CVD Ni, including the large columnar grains with high density of growth twins, the ultrafine and nano grains embedded in-between the columnar grains and the large sized five-fold twinned grains; more importantly, the following conclusions have been drawn. 1. The grain growth behaviour of CVD Ni was found to be a gradual heat release process upon annealing. There is no evidence of recrystallization either on the DSC curve or from the microstructure evolution observation. 2. Massive grain growth in CVD Ni occurred upon annealing above 400 ºC. All grains grew towards equiaxed shape at higher temperatures. The initial bi-modal grain size distribution gradually diminished upon annealing. 3. At 400 ºC, apparent decrease in growth twin density occurred accompanying the grain growth in CVD Ni. The growth twins became much wider at 800 ºC. 4. The ultrafine and nano twins were found to start receding at 400 ºC. Two detwinning mechanisms were proposed: (1) Detwinning by glide of partial dislocations collectively at the ends of twins. These partial dislocations were believed to nucleate at the intersecting region of twins and grain boundaries. (2) Detwinning by dislocation movement within the twinned region. These dislocations were believed, however, to have emitted from the twin boundaries. 5. After the twins receded, the remaining dislocations formed dislocation cell structures at 600 ºC and higher temperatures. The cell structures may have formed by re-organization of randomly distributed dislocations at the detwinning front; they may also form directly at the twin boundaries where dislocations were active.

97 85 6. The driving force for the detwinning process is believed to be the reduction of total energy associated with the high density of ultrafine and nano twins. 7. The five-fold twinned structures were found to be unstable upon annealing up to 600 ºC and were gradually eaten by the surrounding grains and finally disappeared at 800 ºC. 8. Hot stage in-situ TEM detected moderate grain growth of the ultrafine and nano grains, but no grain boundary migration and detwinning process were present in the large grains. Possible reasons leading to the stabilizing behaviour could be (1) free surface pinning effect of the twin boundaries and grain boundaries on the TEM thin foil and (2) release of residual stresses in the as-deposited material due to TEM sample thinning effect, which reduced the driving force for microstructure change in the material. 9. The hardness of the CVD Ni showed a decreasing trend with increasing annealing temperature on both planar and transverse directions. These results are in good agreement with the microstructural evolution of CVD Ni, and comply with the conventional Hall-Patch relation.

98 86 6 Future Work Up to present, more investigations need to be carried out as listed in the following. 1. Measure the residual stresses in the as-received CVD nickels using X-ray diffraction technique. 2. More conventional TEM observation could be carried out on samples annealed from 300ºC to 400ºC, between which grain growth has occurred, in order to verify if a moving grain boundary during grain growth is required for detwinning partial nucleation. 3. The characteristics of the dislocations involved in the detwinning process and dislocation activated along the twin boundaries could be studied by HRTEM. 4. Study the deformation mechanism of CVD Ni by conducting compression tests and analyzing via TEM.

99 Heat Flow W/g Heat Flow (W/g) 87 7 Appendices Appendix 7.1 DSC scans at heating rate of 10ºC/min and 20ºC/min DSC scan at 10ºC/min to 600ºC Temperature (ºC) First run Second run Difference DSC scan at 20ºC/min to 600ºC First run Second run difference Temperature ºC

100 8 88 Appendix 7.2 CVD nickel annealed at 200 ºC, 250 ºC, 300 ºC and 370 ºC for 45 mins Grain growth was observed at 370ºC, suggesting the onset temperature for grain growth of CVD Ni lied between 300 ºC and 370ºC Annealed at 200 ºC Annealed at 250 ºC Annealed at 300 ºC Annealed at 370 ºC

101 Appendix 7.3 TEM images of the black lines traversing the large grains In the large grains traversed by black lines, the lines were found to be comprised of tangled dislocations. The growth twins at the two sides of the dislocation walls were slightly shifted from each other. 89

102 Appendix 7.4 Five-fold twinned structures in the as-deposited CVD Ni 90

103 91 Appendix 7.5 Detwinning phenomena as observed at 400 ºC, 600 ºC and 800 ºC. 400 ºC Detwinning process occurred in the large grains at 400 ºC. Dislocation arrays formed at both the ends of twins and the sites where twins were sectioned by dislocation lines

104 92 Dislocation arrays formed at the sites where twins were sectioned by dislocation walls. Parallel dislocation lines densely aligned in twinned zone

105 ºC Growth twins and dislocation cell structure at 600ºC

106 ºC Growth twins and dislocation cell structure at 800ºC

107 95 Appendix 7.6 The hot stage in-situ TEM images showing a large grain full of growth twins surrounded by ultrafine and nano grains upon annealing up to 752 ºC. The time specified was accumulated from each temperature respectively. Note that at 350 ºC the ultrafine and nano grain started to grow and revealed clearer grain boundaries.