Mechanical Properties of Nanocrystalline Nominally Multilayered Hexagonal Cobalt Electrodeposits

MECHANICAL PROPERTIES OF NANOCRYSTALLINE NOMINALLY MULTILAYERED HEXAGONAL COBALT ELECTRODEPOSITS

KRISTA MARIJA VIOLA

A THESIS SUBMITTED IN CONFORMITY WITH THE REQUIREMENTS FOR THE DEGREE OF MASTER OF APPLIED SCIENCE GRADUATE DEPARTMENT OF MATERIALS SCIENCE AND ENGINEERING UNIVERSITY OF TORONTO

 2016 KRISTA MARIJA VIOLA MECHANICAL PROPERTIES OF NANOCRYSTALLINE NOMINALLY MULTILAYERED HEXAGONAL COBALT ELECTRODEPOSITS

Krista Marija Viola

Master of Applied Science Graduate Department of Materials Science and Engineering University of Toronto

ABSTRACT The microstructure and mechanical properties of electrodeposited nanocrystalline cobalt were investigated and compared to cobalt electrodeposits produced under waveforms that would result in a nominal multilayered material by alternating electrodeposition conditions in the same electrolytic solution.

All sample types were of the hexagonal crystal structure and a preferred orientation was prominent with the introduction of nominal multilayers, in which the basal plane preferentially was oriented parallel to the surface of the deposit. Transmission electron microscopy was used to compare the starting microstructure and post-failure microstructure of cobalt electrodeposits. Tensile tests were performed at a strain rate of 5 x 10-4 s-1 and microhardness tests were performed under a 100g load. The average hardness, yield, ultimate tensile and fracture strengths increased when the electrodeposited cobalt followed a nominal multilayered pulse train. Tensile elongation for cobalt electrodeposits with 100 nm nominal layer thicknesses are more than twice that observed for monolithic cobalt.

ii ACKNOWLEDGEMENTS

I wish to thank my supervisor, Professor D. Perovic for his excellent supervision and support throughout this research program, and my committee: Professor G.D. Hibbard and Professor U. Erb.

I wish to acknowledge Dr. J. McCrea and Integran Technologies Inc for nanocrystalline cobalt production. I thank Mr. S. Boccia, Mr. D. Grozea, Mr. M. Daly, Mr. H. Kuntz, Dr. A. Lausic, Mr. J. Tam, Mr. A. Delhaise (Department of Materials Science and Engineering, University of Toronto), Ms. J. Howe, Mr. P. Woo, Mr. C. Soong (Hitachi High- Technologies Canada Inc) for their assistance and contributions to this research.

Finally, I wish to thank my husband, Perry Haldenby, for his constant guidance and patience throughout the course of this research.

iii TABLE OF CONTENTS ABSTRACT ...... ii ACKNOWLEDGEMENTS ...... iii LIST OF TABLES ...... vi LIST OF FIGURES ...... vii LIST OF APPENDICES ...... x 1 INTRODUCTION ...... 1 2 LITERATURE REVIEW ...... 2 2.1 Nanostructured Materials ...... 2 2.1.1 Synthesis ...... 2 2.1.1.2 Electrolyte Constituents ...... 6 2.1.1.3 Current Parameters ...... 8 2.1.2 Crystallographic Structure ...... 12 2.1.2.1 Deformation Mechanisms ...... 14 2.1.1.2 Multilayered Materials ...... 17 2.1.3 Mechanical Properties ...... 20 2.1.3.1 Strength and Hardness ...... 20 2.1.3.2 Young’s Modulus ...... 21 2.1.3.3 Ductility ...... 22 2.1.3.4 Wear resistance ...... 25 2.1.3.5 Corrosion resistance ...... 27 2.2 Nanostructured Electrodeposited Cobalt ...... 29 2.2.1 Crystallographic Structure ...... 29 2.2.2 Deformation Mechanisms ...... 31 2.2.3 Multilayered Materials ...... 34 2.2.4 Mechanical Properties ...... 35 2.2.5 Applications ...... 38 2.2.5.1 Wear Resistance and Tribological Behaviour ...... 38 2.2.5.2 Corrosion Resistance ...... 39 3 EXPERIMENTAL...... 40 3.1 Electrodeposition ...... 40 3.2 Characterization ...... 41 3.2.1 X-ray Diffraction...... 41 3.2.2 Bulk Alloy Composition ...... 42 3.2.3 Scanning Electron Microscopy ...... 42 3.3.3 Transmission Electron Microscopy ...... 42

iv 3.3 Properties ...... 44 3.3.1 Microhardness ...... 44 3.3.2 Tensile Testing...... 44 4 RESULTS AND DISCUSSION ...... 45 4.1 Sample Identification ...... 45 4.2 Crystallographic Structure ...... 45 4.2.1 Monolithic Cobalt ...... 47 4.2.2 Multilayered Cobalt ...... 51 4.2.3 Solute Concentration ...... 56 4.3 Properties ...... 60 4.3.1 Microhardness ...... 60 4.3.2 Tensile Testing...... 62 5 CONCLUSIONS ...... 75 6 RECOMMENDATIONS ...... 76 APPENDICES ...... 77 Appendix A: FIB microsampling procedure ...... 77 Appendix B: Additional TEM Images of as-deposited and near-fracture surface specimens...... 80 7 REFERENCES ...... 96

v LIST OF TABLES

Table 1: Mechanical properties of Polycrystalline vs. Nanocrystalline Ni [A. Robertson et al. 1999]

Table 2: Consolidated vs. Electrodeposited nanocrystals with respect to polycrystalline counterparts [Erb et al. 1997]

Table 3: Mechanical and structural effects of current type on electrodeposited metals

Table 4: Tensile Properties of polycrystalline and nanocrystalline Co at different strain rates [Karimpoor et al. 2003] Table 5: Sample identification and bulk purity analysis via XRF

Table 6: C and S concentration as determined via ASTM E1019-11 (ppm)

Table 7: Solute concentration for sulfur [S] and carbon [C], grain size (d) and grain size range (r) and standard deviation (s) of nanocrystalline electrodeposited Co [Hibbard et al. 2006]

Table 8: Vickers microhardness results, taken under a 100g load and dwell time of 15 seconds

Table 9: Mechanical properties obtained from engineering stress-strain curves

vi LIST OF FIGURES

Figure 1: Electrocrystalization behaviour; density of metal ions at the cathode surface is function of distance away from surface [Koch 2007]

Figure 2: Grain size with respect to saccharin concentration in the electrolyte solution of nickel electrodeposits [El-Sherik and Erb, 1995]

Figure 3: XRD patterns of nickel electrodeposits with changing saccharin concentrations in the electrolyte solution [El-Sherik and Erb, 1995]

Figure 4: Pulse current plating waveform with time on and off (Ton and Toff) peak CD values (Jp) and average current density (Jm). [El-Sherik et al. 1995]

Figure 5: XRD and DP patterns for polycrystalline cold and hot rolled annealed electrowon (a) and nanocrystalline (b) cobalt [Karimpoor et al. 2003]

Figure 6: Volume fraction of crystalline and intercrystalline components with respect to grain size where grain boundary thickness is assumed as 1nm [Wang et al. 1997]

Figure 7: Schematic diagram of deformation evolution in nanocrystalline nickel including dislocation motion, void formation and unconstrained ligaments [Kumar et al. 2003]

Figure 8: Schematic diagram in plane (a) and perspective (b) view of edge crack mechanisms as propograting through brittle and tough layers [Srolovitz et al. 1995]

Figure 9: Fracture resistance as a function of crack length [Srolovitz et al. 1995]

Figure 10: Intermediate fine layers as observed following tensile testing [Fiebig et al. 2016]

Figure 11: Hardness with reduction in grain size of nanocrystalline nickel electrodeposits [El-Sherik et al. 1992]

Figure 12: Electrodeposited Ni-P grain size as a function Young’s modulus [Zhou et al. 2009]

vii Figure 13: Fracture orientation of nanocrystalline cobalt at a strain rate of 5 X 10 -4 s-1 [Karimpoor et al. 2006]

Figure 14: Fracture surface of nanocrystalline cobalt following tensile testing [Karimpoor et al. 2006]

Figure 15: Taber wear index with respect to average grain size for Ni electrodeposits [Jeong et al. 2001]

Figure 16: Vickers hardness as a function of taber wear index for nanocrystalline Ni [Jeong et al. 2003]

Figure 17: XRD patterns of nanocrystalline Co from cathode facing surface (a), mid- section (b), and free surface (c) [Karimpoor et al. 2007] Figure 18: Stress-strain curves for polycrystalline and nanocrystalline Co at strain rates of 1 X 10-4 s-1 (1), 5 X 10-4 s-1 (2) and 2.5 X 10-3 s-1 (3) [Karimpoor et al. 2003] Figure 19: Schematic diagrams of (1) MN, (2) ML20 and (3) ML100 electrodeposited Co Figure 20: X-ray diffraction patterns for reference FCC and HCP Co using Cu-K radiation

Figure 21: Tensile coupon measurements (mm) Figure 22: X-ray diffraction patterns of MN, ML20, and ML100 specimens using Cu-K radiation

Figure 23: BF (a) and (c) and DF (b) and (d) TEM images and (e) SADP inset (L=300) of MN Co

Figure 24: Grain size distribution for MN Co with log-normal distribution

Figure 25: BF (a) and DF (b) images of grain near perforation in MN Co (DF image producted by selected diffraction around the (100) plane Figure 26: 100nm layered Co deposit (X50.0K); scale bar represents 1m

Figure 27: 500nm layered Co deposit (X10.0K); scale bar represents 5m

viii Figure 28: ML20 (a) and ML100 (b) BF TEM images

Figure 29: Grain size distribution for ML20 Co with log-normal distributiom Figure 30: Grain size distribution for ML100 Co with log-normal distributiom Figure 31: DP of MN(a) ML20 (b) and ML100 (c). Textured regions along the (0002) and (10 ̅11) rings are circled in (b) and (c) Figure 32: XRD patterns for electrodeposited nanocrystalline Co as discussed in Table 7 above [Hibbard et al. 2006]

Figure 33: Tensile test results of monolithic cobalt (MN) at a strain rate of 5 X 10-4 s-1

Figure 34: Tensile test results of 20nm multilayered cobalt (ML20) at a strain rate of 5 X 10-4 s-1 Figure 35: Tensile test results of 100nm multilayered cobalt (ML100) at a strain rate of 5 X 10-4 s-1 Figure 36: Average tensile test results of monolithic and multilayered cobalt at a strain rate of 5 X 10-4 s- Figure 37: SEM imaging periodic features on fracture surface of coarse grained and nanocrystalline grained multilayered NiCo; scale bar represents 5m [Daly et al. 2015]

Figure 38: XRD peak intensities for ML100, ML20 and MN deposits

Figure 39: Tensile elongation vs index of (200) peak for Ni-Fe and Ni-W [Matsui et al. 2013]

Figure 40: Fracture location from DIC imaging (a) and facture surfaces (b) imaged via secondary electrons (SE) at 15kV for MN, ML20 and ML100 Co. Scale bar in (b) represents 1mm

Figure 41: SE images of fracture surfaces for MN (a), ML20 (b), and ML100 (c) at 10kV showing dimpled fracture surfaces. Scale bar represents 10m

Figure 42: BF (a) and (c) and DF (b) and (d) xTEM images of MN Co Figure 43: BF (a) and (c) and DF (b) and (d) xTEM images of ML20 Co

ix Figure 44: BF (a) and (c) and DF (b) and (d) xTEM images of ML100 Co

LIST OF APPENDICES

Appendix A: FIB Microsampling Procedure

Appendix B: Additional TEM Images of as-deposited and near-fracture surface specimens

x 1.0 INTRODUCTION Electrodeposition of nanocrystalline cobalt is a desirable production method to form fully dense nanostructured materials with high strength, corrosion and wear resistance as a hard chrome replacement or biomaterial [Hibbard et al. 2001; Karimpoor 2001; McCrea 2010; Spriano 2005]. Recently, multilayered electrodeposits have been investigated for their mechanical properties following a rule-of-mixtures principle when polycrystalline and nanocrystalline electrodeposits are in alternating order throughout the thickness of the material [Daly et al. 2015]. The development of these materials as produced within a single electrolytic solution is beneficial to achieve nanocrystalline materials with advantageous mechanical properties over their polycrystalline counterparts without a significant reduction in ductility that has been previously observed in monolithic nanostructures [Aus et al. 1992; Brooks et al. 2011; Wang et al. 1995]. The focus of this research is the nanostructure characterization and mechanical properties investigation of multilayered electrodeposited cobalt materials with two sub- layer thicknesses: 20nm and 100nm, in comparison to monolithic cobalt. In each case, the starting microstructure and post-failure microstructures are observed and their relation to mechanical properties is investigated. In the experimental sections the characterization techniques: tensile tests, microhardness tests, bulk chemical analysis and fractography methods are discussed. Results, conclusions and recommendations are presented and additional electron microscopy images are included in appendices.

1 2 LITERATURE REVIEW

2.1 Nanostructured Materials

2.1.1 Synthesis Nanocrystalline materials offer exceedingly improved materials properties in comparison to their polycrystalline counterparts. Nanomaterials may be formed through a multitude of methods: nanocrystalline materials are produced through techniques including deformation of pre-formed materials (eg. ball milling) and direct formation (eg. chemical vapour deposition, consolidation, inert gas condensation, sol-gel processes, and electrodeposition) [Gleiter et al. 1989]. Electrodeposition may be used to produce non-equilibrium nanocrystalline metals or alloys with desirable properties tailored by electrolytic solution constituents and conditions (eg. current density, frequency, duty cycle). Methods applied in regards to the production of electrodeposited nanocrystalline materials are discussed elsewhere (US Patent # 5,352,266 and 5,433,797).

Considering nanocrystalline materials with a grain shape of a 14-sided tetrakaidecahedron with faces as grain boundaries and edges as triple junctions, it has been calculated that the volume fractions of intercrystalline components within a material increases and perfect crystal regions within grains decrease as grain size is reduced. The properties of nanocrystals under 10nm in diameter are influenced majorly by the large presence of triple junctions and grain boundaries [Palumbo et al. 1990].

Table 1 outlines the mechanical properties of nanocrystalline nickel in comparison to its polycrystalline counterpart [Robertson et al. 1999]. For the purpose of this discussion, polycrystalline or coarse-grained materials will hereafter refer to materials with average grain sizes greater than 1 um to a few mm in diameter. Ultra-fine grained (UFG) materials will hereafter refer to materials with average grain sizes in the range of 100nm to 1um in diameter. Nanocrystalline materials will hereafter refer to materials with an average grain size at or below 100nm in diameter.

2 As shown in Table 1, nanocrystalline Ni with an average grain size of 10nm shows significant improvements in mechanical properties, such as increases in yield and ultimate tensile strength values but with a significant decrease in tensile elongation from 50% for polycrystalline Ni and 1% for nanocrystalline Ni. Electrodeposited nanocrystalline materials that have similar mechanical property improvements but that maintain high tensile elongations or ductility are desired [Brooks et al. 2011].

Table 1: Mechanical properties of Polycrystalline vs. Nanocrystalline Ni [A. Robertson et al. 1999]

Property Ni 10m[12] Ni 100nm Ni 10nm

Yield Strength, MPa (25 C) 103 690 >900

Ultimate Tensile Strength, MPa (25 C) 403 1100 >2000

Tensile Elongation, % (25 C) 50 >15 1

Elongation in Bending, % (25 C) - >40 -

Modulus of Elasticity, GPa (25 C) 207 214 204

Vickers Hardness, kg/mm2 140 300 650

Work Hardening Coefficient 0.4 0.15 0.0

Fatigue Strength,, MPa (108 cycles/air/25 C) 241 275 -

Wear Rate (dry air pin on disc), m3/m 1330 - 7.9

Coefficient of Friction (dry air pin on disc) 0.9 - 0.5

2.1.1.1 Electrodeposition Non-equilibrium structured materials with reduced grain sizes, a large volume fraction of grain boundaries and triple junctions, and negligible porosity may be produced from electrodeposition [Erb et al. 1997]. The electrodeposition parameters and conditions used are capable of controlling properties of the deposits produced (i.e. grain size, surface

3 roughness, preferred crystallographic orientation, tensile ductility, etc.) and will be explored further in this discussion. The parameter variables include electrolyte constituents, electrolyte temperature, pH, current density, duty cycle, frequency, among others.

The electrodeposition of metals is well-known to be dually influenced by competing crystal nucleation and growth. For the purpose of creating nanocrystalline deposits, high nucleation and low grain growth rates are desired. These processes are affected by the rate of charge transfer and diffusion of adsorbed ions (adions) at the electrode surface and may be retarded by plating parameters as mentioned above [Choo et al. 1995]. Crystal nucleation is favoured by high overpotential and low surface diffusion rates, whereas crystal or grain growth is promoted by low potential and high surface diffusion rates. Figure 1 shows a schematic diagram of competing nucleation and growth processes at the cathode and the metal ion density as a function of the distance away from the cathode surface. The overpotential has been found to be a function of current density adjustments and may also be reduced by certain additives [Koch 2007; Ma et al. 2015].

The consolidation of nanocrystalline powders has been shown to create materials with undesired porosity and impurities that influence material’s properties – specifically, through reducing the elastic constant, saturization magnetization, and Curie temperature, and through increasing specific heat and thermal expansion. Electrodeposited materials have been noted to not experience such changes at all, or only to some minor degree (i.e. only 5% change vs. 40-50% as with consolidated materials) [Erb et al. 1997]. Table 2 outlines the changes in properties for consolidated and electrodeposited nanomaterials in comparison to their polycrystalline counterparts. It is evident that electrodeposited materials can achieve the desired properties similar to those measured with polycrystalline counterparts owing to the reduced porosity and impurities present.

Figure 1: Electrocrystalization behaviour; density of metal ions at the cathode surface is function of distance away from surface [Koch 2007]

Table 2: Consolidated vs. Electrodeposited nanocrystals with respect to polycrystalline counterparts [U. Erb et al. 1997]

Property Consolidated Materials Electrodeposited Materials

Young’s Modulus Reduced by 80% (15) Unchanged (18)

Thermal Expansion Increased by 80% (16) Unchanged (19)

Specific Heat Increased by 50% (16) Increased by < 5% (19)

Saturation Magnetization Reduced by 40% (16) Decreased by < 5% (20)

Curie Temperature Reduced (17) Unchanged (21)

Recent studies have shown that electrodeposited fully-dense nanocrystalline material properties are not simply affected by grain size alone, but properties such as tensile ductility, plastic deformation and early Bauschinger effect are impacted by microstructural homogeneity and grain orientation [Matsui et al. 2013; Rajagopalan et al. 2011]. Accordingly, further exploration of electrodeposited nanocrystalline material properties as influenced by deposit microstructure is required.

2.1.1.2 Electrolyte Constituents Electrolyte composition for electrodeposition of nanocrystalline metals may be tailored to favour high nucleation rates and low grain growth. The two competing parameters for electrodeposition of nanocrystalline materials are grain growth and nucleation; the rate determining steps are surface diffusion of adions on the crystal surface and charge transfer at the electrode surface [El-Sherik and Erb, 1995; Karimpoor, 2001]. The addition of organic additives has been shown to increase the overpotential and reduce diffusion rates, which promote nucleation, of electrodeposited metals and additionally the hardness of the deposited metal [Ma et al. 2015; Yang et al. 2010].

Watts’ type electrolytic solutions (eg. for nickel electrodeposition constituents include nickel sulfate, nickel chloride and boric acid) are commonly used in combination with organic additives to increase the cathodic overpotential or tailored pH and current density to achieve textured electrodeposits [Alimadadi et al. 2014; Alimadadi et al. 2016; Li et al. 2011].

Organic additives, namely saccharin, have been studied for their effect on electrodeposited metals with increasing presence in electrolytic solutions. Figure 2 shows nickel electrodeposit grain size as a function of saccharin electrolyte concentration. The authors attributed grain refinement to saccharin lowering the overpotential for nickel ion reduction and blocking crystalline growth and reduced surface diffusion [El-Sherik and Erb, 1995]. Figure 3 shows the change in preferred orientation, or texture, and peak breadth in the X-ray Diffraction (XRD) patterns as a result of increasing saccharin concentration in the electrolyte for nickel deposits. The preferred orientation is (200) with no saccharin present in the electrolyte bath and reduces in intensity with increase in saccharin concentration.

6 The texture of electrodeposited metals has been shown to be affected by the substrate crystal structure to a certain extent [Knock et al. 2000] and organic additives [Yang et al., 2010]. Changes in preferred deposit orientation may promote desirable qualities, such as tensile ductility [Matsui et al. 2013], therefore a change in saccharin content may be an employable method to tailor electrodeposit properties.

Figure 2: Grain size with respect to saccharin concentration in the electrolyte solution of nickel electrodeposits [El-Sherik and Erb, 1995]

Figure 3: XRD patterns of nickel electrodeposits with changing saccharin concentrations in the electrolyte solution [El-Sherik and Erb, 1995]

2.1.1.3 Current Parameters The density of atomic packing in electrodeposited materials has been found [Pangarov 1962] to be a factor of the current density (CD) and temperature. Low CD and high temperatures induce densely packed crystallographic planes parallel to the substrate surface, compared to high CD and low temperatures inducing packed planes in the perpendicular direction to the substrate [Pangarov 1962].

Crystallographic texture of electrodeposits has been attributed to overpotential, direct current (DC) vs pulsed (PC) currents, pH, and surface adsorbates or inhibiting species parameters [Bhardwaj et al. 2011; Pangarov 1962]. Figure 4 shows typical PC parameters, cathodic square wave pulses with time on and off (Ton and Toff) and average and peak CD values (Jm and Jp, respectively). PC parameters are employed to reduce grain

8 size, increase hardness and change preferential texture in electrodeposited metals as compared to direct current plated counterparts [Sanacian et al. 2014].

Pulse current plating conditions may also be used to adjust alloy compositions, allow higher current densities than DC plating to give high overpotential and low surface diffusion rates that promote nucleation (vs. low overpotential and high surface diffusion rates favoured by grain growth) [Choo et al. 1995; Allahkaram et al. 2011]. Table 3 outlines the properties achieved by change in current densities and parameters for electrodeposition of various metals and alloys. Pulse current electrodeposition has been investigated in comparison to direct current plating conditions and found that it produces an overall refined grain size, reduced porosity and improved hardness, with varying surface morphologies, texture and roughness [Allahkaram et al. 2011; El-Sherik et al. 1995; El-Sherik et al. 1996; Kumar et al. 2013; Saitou et al. 2001]. It has also been found, to have higher current efficiencies than DC plating [Allahkaram et al. 2011].

Pulse reverse (PR) current conditions have been investigated [Liu et al. 2007] for electrodeposition of copper and it was found that within pulse polarization plating conditions, the positive pulse current promotes crystal nucleation at the cathode surface, whereas the negative pulse promotes grain growth to achieve decreased grain size in deposits with varying morphologies (e.g. pyramidal, granular crystals). The negative current (dependent on duration and amplitude) was also found to partially dissolve plated metal at the cathode surface, which increases the adion concentration at the cathode surface and recrystallization may occur.

Figure 4: Pulse current plating waveform with time on and off (Ton and Toff) peak CD values (Jp) and average current density (Jm). [El-Sherik et al. 1995]

10 Table 3: Mechanical and structural effects of current type on electrodeposited metals

Material Current CD Time Surface Hardness Grain Reference Type (A/dm2) On/Off roughness (HVN) size (um) (nm) Ni DC 100 --- Ra=0.37 350 45 [Sanacian et al. ma/cm2 rz=2.1 2014] Ni PC 100 2ms-18ms Ra=0.25 556 31 [Sanacian et al. ma/cm2 rz=1.6 2014]

NiWTiO2 DC 1.5 - Somewhat 467-686 - [Kumar et al. uniform 2013]

NiWTiO2 PC 1.5 40ms:30ms Smooth Higher (Finer [Kumar et al. and than DC than 2013] smaller (by about NiWTiO2 spherical 100) plated grains under (more DC) uniform than NiWTiO2 plated under DC) NiCo PC 10 50% duty - 450 - [Zamani et al cycle 2016] NiCo DC 5 - Irregular <300 25.99 [Karslioglu et polyhedral, al. 2015] ~0.4 um

NiCo PC 5 10 10 ms Reduced ~300 24.5 [Karslioglu et crystallite al. 2015] size, ~0.2 um NiCo PRC 5 10/10/10ms Spherical >400 23.93 [Karslioglu et cluster al. 2015] equiaxed, ~0.2 um CoP DC 10 - Smooth - 30 [Kosta et al. fine 2011] globular CoP PC 175 2.5ms/45ms Smooth - 10 [Kosta et al. fine 2011] globular

11 2.1.2 Crystallographic Structure Electrodeposited nanocrystalline metals are often characterized by means of Transmission Electron Microscopy (TEM), electron diffraction patterns (DP) and X-ray diffraction (XRD). TEM imaging is commonly employed to resolve the nanostructured material when the composition, phase, or orientation across the bulk material is unvaried. For example, while Scanning Electron Microscopy (SEM) may image the morphology and composition of specimens, the microstructure of monolithic electrodeposits is commonly imaged via TEM analysis to compare to their respective polycrystalline analogs [El-Sherik et al. 1995; El-Sherik and Erb 1995; Karimpoor 2002].

XRD of nanocrystalline materials demonstrate peak-broadened patterns in comparison to patterns acquired from polycrystalline materials, as a result of the refined grain size. Figure 5 shows a comparison of XRD patterns obtained from polycrystalline and nanocrystalline Co [Karimpoor et al. 2003]. As shown in Figure 5, peak-broadening is observed as well as a change in crystal structure from mixed FCC-HCP (25% FCC) polycrystalline to pure HCP nanocrystalline Co with a strong (0002) texture. Polycrystalline Co was produced by hot and cold rolling, and annealing (up to 1000 C in Ar in an electric tube furnace) electrowon Co [Karimpoor et al. 2003].

12 (a)

(b)

Figure 5: XRD and DP patterns for polycrystalline cold and hot rolled annealed electrowon (a) and nanocrystalline electrodeposited (b) cobalt (Co-K radiation, =1.7902 퐀̇ ) [Karimpoor et al. 2003]

13 2.1.2.1 Deformation Mechanisms Plastic deformation of crystalline materials can occur by atomic diffusion via vacancy or interstitial defects along grain boundaries (Coble creep) or through grains (Nabarro- Herring creep) or via dislocation slip, with a steady creep state linear to the applied stress [Wang et al. 1997]. However, as grain size is reduced the material response to stress has shed light on alternative deformation mechanisms that were proposed in response to observed softening effects, strain rate sensitivity, work hardening, and superplasticity in nanocrystalline materials [Dalla Torre et al. 2005; Chan et al. 2014].

Simulation of tensile tested nanocrystalline metallic FCC material has shown that grain boundary (GB) sliding, stress-assisted free volume migration, and dislocation mechanisms were observed. Dislocations and partial dislocations were emitted from GB’s, often forming stacking faults (SF), and were reabsorbed at the GB and generate a twin boundary or full dislocation, depending on the materials stacking fault energy (SFE) and critical grain size for the emission of a trailing dislocation [Van Swygenhoven et al. 2006]. Dislocation pile-up mechanisms have been calculated to break down for FCC metals at grain sizes near 10nm and calculations show that considering nanocrystalline materials as composite models, both crystalline and intercrystalline (GB’s, triple junctions, quadruple nodes) contribute to the nanostructured materials strength [Wang et al. 1995; Wang et al. 1997].

The volume fraction of each of these components with respect to grain size within the nanocrystalline regime is shown in Figure 6. This figure illustrates that within certain grain size limits, the majority of the material components and effective properties (including strength) is described by the strength of the intercrystalline components. Based upon tensile testing of nanocrystalline Ni, it has been found that the strength of the intercrystalline components decreases in the following order: grain boundary, triple junction and quadruple node [Wang et al. 1997]. Therefore, as the grain size reduces and the volume fraction of triple junctions and quadruples nodes increases there is a softening effect as commonly observed in nanocrystalline materials.

14 Diffusional creep of atoms within intercrystalline components, along with the main grain boundary sliding deformation mechanism, is noted as significant within nanocrystalline grains under 20nm [Wang et al. 1997]. The softening effect as described above is only applicable for grain boundary sliding as the dominant deformation mechanism; when the major deformation mechanism is dislocation movement further grain refinement will strengthen the material as per the Hall-Petch relationship [Hahn et al. 1997].

Figure 6: Volume fraction of crystalline and intercrystalline components with respect to grain size where grain boundary thickness is assumed as 1nm [Wang et al. 1997]

In situ and ex situ deformation and high resolution TEM imaging of nanocrystalline Ni with an average grain size of 30 nm demonstrated that ductile fracture occurred through dimpled rupture, where voids at triple junctions were proposed as dimple nucleation sites larger than the average grain size observed [Zhu et al. 2005]. Dimpled fracture surfaces of nanocrystalline electrodeposited Ni have also been observed, particularly involving dimple measurements surpassing the average observed grain size [Zhu et al. 2005].

15 Twinning was also postulated as a possible deformation mechanism, observed as solid bands of alternating contrast within the imaged grains, [Kumar et al. 2003] but have been negated as likely events in simulated deformation of the similar material [Zhu et al. 2005]. However, it has been observed in electrodeposited nickel with an average grain size of 100nm following cyclic loading [Cheng et al. 2009] and in other FCC nanocrystalline FCC materials, such as Pd and Cu [Ebrahimi et al. 2006; Sriram et al. 2008]. The evolution of deformation in nanocrystalline materials as developed from these results is shown in Figure 7 [Kumar et al. 2003]. The figure illustrates how void formation at grain boundaries and triple junctions, dislocation movement and plastic deformation of individual grains form the dimpled fracture surface observed.

Figure 7: Schematic diagram of deformation evolution in nanocrystalline nickel including dislocation motion, void formation and unconstrained ligaments [Kumar et al. 2003]

16 2.1.1.2 Multilayered Materials Electrodeposition of multilayered (ML) materials may be used to tailor the mechanical properties and microstructure of deposits. The applied design of a layered waveform combination to combine coarse grained, or polycrystalline, electrodeposits sandwiched between nanocrystalline grained electrodeposits has been shown to govern the strength and ductility of the layered specimens, often following a ‘rule of mixtures’ principle [Chan et al. 2012; Daly et al. 2015; Fiebig et al. 2016; Kurmanaeva et al. 2014; Kurmanaeva et al. 2016; Srolovitz et al. 1995].

Generally, the coarse grained layers allow for increased plastic deformation and offer high toughness by bridging through-thickness cracks and works in conjunction with the high strength offered by nanocrystalline layers to provide a composite with tailored mechanical properties. Figure 8 shows a schematic diagram of crack propagation through a composite structure composed of strong/brittle and tough layers [Srolovitz et al. 1995]. The tough, more ductile layers bridge fractured strong/brittle layers and subsequently deform by necking until final failure; as the crack length (a) approaches the bridged material length (L), the composite toughness is predicted to reach steady state value, as shown in Figure 9, and is a function of interfacial de-bonding [Srolovitz et al. 1995].

Room temperature tensile testing of electrodeposited ML NiFe alloys with alternating coarse grained and nanocrystalline layers (1:1 ratio of 5um thick layers with average grain sizes of 500nm and 16nm, respectively) showed that post-failure, a refined grain size layer was formed in-between the coarse and nanocrystalline grained layers [Fiebig et al. 2016]. Grains were elongated in the tensile direction and contained deformation twins [Fiebig et al. 2016]. The formed layer is shown in Figure 10 and is explained by the authors by dislocation pile-up causing internal stress to induce nanocrystalline grain rotation and coalescence along layer interfaces.

The same study [Fiebig et al. 2016] also attributed the ML deformation mechanisms to dislocations within the coarse grained layer that pile-up along the coarse grained- nanocrystalline grained layer interface and create a back stress that hinders dislocation

17 movement along the same slip plane and therefore work hardens the coarse grained layer [Fiebig et al. 2016]. The deformation mechanisms and resulting strength of the layered electrodeposits are also influenced by layer interface structure and coherency, i.e. a change in phase between layers is expected to hinder dislocation mobility between layers and the strength is again influenced by dislocation pile-up at layer interfaces, which manifested itself as a slight decrease in ductility for ML NiFe [Kurmanaeva et al. 2014; Fiebig et al. 2016].

The mechanical properties of ML specimens have been studied with decreasing layer thicknesses (1:1 thickness ratio for coarse and nanocrystalline grained layers) from 5um to 30nm for NiFe alloys. It was found that the ML specimens had increasing hardness with decreasing layer thickness up to 100nm layer thickness, below which hardness remained relatively constant [Kurmanaeva et al. 2016].

Figure 8: Schematic diagram in plane (a) and perspective (b) view of edge crack mechanisms as propograting through brittle and tough layers [Srolovitz et al. 1995]

Figure 9: Fracture resistance as a function of crack length [Srolovitz et al. 1995]

Figure 10: Intermediate fine layers as observed following tensile testing of electrodeposited ML NiFe alloys with alternating coarse grained and nanocrystalline layers [Fiebig et al. 2016]

2.1.3 Mechanical Properties

2.1.3.1 Strength and Hardness As discussed above, grain size of electrodeposited materials may be refined through several parameters. Material yield strength or hardness with respect to grain size is defined by the Hall-Petch relationship [Hall 1951; Petch 1953]. Hall-Petch behavior is commonly used to convey the effect of strengthening (yield strength or hardness) through grain size reduction, −1/2 −1/2 expressed as: 휎푦 = 휎0 + 푘푑 or 퐻 = 퐻0 + 푘퐻푑 , where σy and H the material’s

0.2% yield strength and hardness, respectively, k is a material constant, σ0 and H0 represent the stress required to move a dislocation through the lattice and d is the average grain diameter [Wang et al. 1995]. The Hall-Petch behavior is illustrated in Figure 11. A plateauing effect of the measured hardness value of nNi electrodeposits was observed once the grain size became significantly reduced. It is known that there is a limit to the increase in material strength through grain size reduction, i.e. below a critical grain diameter, the material deviates from typical Hall-Petch behaviour and there is a softening effect, explained by increasing interface volume fraction and grain boundary processes that surpass typical deformation mechanisms for polycrystalline materials [El-Sherik et al. 1992; Van Swygenhoven et al. 2006].

Figure 11: Hardness with reduction in grain size of nanocrystalline nickel electrodeposits [El- Sherik et al. 1992]

Electrodeposited metals such as Ni, Cu, Co, Ni-Fe and Ni-Co have seen increases in strength and hardness with decreasing grain size in agreement with the Hall-Petch relationship [Cheung et al. 1994; Daly et al. 2015; Karimpoor et al. 2003; Sriram et al. 2008; Wang et al. 1997].

2.1.3.2 Young’s Modulus Young's modulus has been previously reported [Zhou et al. 2009] to decrease with nanocrystalline grains (< 17nm) owing to interfacial contributions i.e. with respect to excess free volume in the interface regions or increasing volume fraction of intercrystalline components. In some cases, a reduction in Young’s modulus for nanocrystalline materials may be due to a change in crystal structure (i.e. mixed FCC-HCP polycrystalline Co compared to pure HCP nanocrystalline electrodeposited Co) [Karimpoor et al. 2003]. Electrodeposited nanocrystalline materials have not shown the decrease in Young’s modulus as observed with consolidated nanocrystalline materials with high residual porosity [Robertson et al. 1999]. Figure 12 shows the relationship between grain size and Young’s modulus for nanocrystalline Ni-P. The Young’s modulus value was found to

21 decrease with decreasing grain size until reaching approximately the value for amorphous Ni-P. Similar studies [Erb et al. 1997; Karimpoor et al. 2003; Robertson et a. 1999] have concluded that Young’s modulus for nanocrystalline materials decreases to only some minor degree in comparison to their polycrystalline counterparts, or remains unchanged. The small decrease in value has been attributed to change in crystallographic structure, texture, or the increasing volume fraction of intercrystalline components for average grain sizes < 10 nm.

Figure 12: Electrodeposited Ni-P grain size as a function Young’s modulus [Zhou et al. 2009]

2.1.3.3 Ductility Generally, the intrinsic ductility of electrodeposited nanocrystalline Ni has been investigated [Brooks et al. 2011] and found that the tensile ductility of electrodeposited metals was highly dependent on the presence of defects within the deposit but is independent of deposit microstructure within a grain size range of 10-80nm, as the uniform plastic strain did not vary significantly from specimen to specimen tested. The authors concluded that strain-oriented phenomena control grain-boundary mediated damage with respect to nanocrystalline metals and is best defined by a critical plastic strain independent of the material strength [Brooks et al. 2011].

22 It was also found that the gauge volume of the tensile coupons tested had no significant effect on the measured tensile properties for electrodeposited nanocrystalline metals [Wei et al. 2007]. Nanocrystalline Ni specimens were also investigated [Chan et al. 2012] for stress-induced heat generation and it was observed that no significant heating arose and that it is unlikely to cause grain boundary migration during tensile testing.

Brooks et al. [2011] completed a study of nanocrystalline nickel electrodeposited in a Watts’-type bath in tensile testing and found that the intrinsic ductility (maximum uniform plastic strain) was independent of nickel microstructure over an average grain size range of 10nm – 80nm. The conclusions drawn were that deformation mechanisms involving grain boundaries are strain-oriented and are defined by a critical plastic strain. This was also found to be independent of the material strength.

Nanocrystalline materials offer different mechanical properties than their polycrystalline counterparts, including increased tensile and compressive strength, hardness, wear resistance, and corrosion resistance [Erb et al. 1997; Karimpoor et al. 2002; Karimpoor et al. 2003; Wang et al. 2006]. However, nanocrystalline materials have a corresponding decrease in tensile ductility or elongation, which is reduced with respect to grain size.

Electrodeposited nanocrystalline Co (average grain size 12nm) with an HCP structure was investigated [Karimpoor et al. 2002] and compared to equiaxed polycrystalline Co with a 17% FCC – 83% HCP structure (average grain size 5.5um). It was found that with increased hardness for polycrystalline to nanocrystalline Co (232 VHN to 525 VHN), yield (311 MPa to 1002 MPa) and tensile strengths (811 MPa to 1865 MPa) and similar values for Young’s modulus (207 GPa to 200 GPa) the average elongation to failure for nanocrystalline was only decreased 10% to 7% for polycrystalline Co at a strain rate of 5 x 10-4 s-1. This is much higher than the average elongation to failure for similarly prepared nanocrystalline Ni (<1%) with a similar average grain size, although polycrystalline Ni has a higher ductility than polycrystalline Co. Polycrystalline Co was

23 produced by hot and cold rolling, and annealing (up to 1000 C in Ar in an electric tube furnace) electrowon Co.

Fracture surfaces of similarly produced nanocrystalline Co [Karimpoor et al. 2006] exhibited a flat plateau shape with ledges and a fine-dimpled fracture surface (in comparison to polycrystalline Co), as shown in Figures 13 and 14, respectively, indicative of some plastic deformation and microvoid coalescence, respectively.

Figure 13: Fracture orientation of nanocrystalline cobalt at a strain rate of 5 X 10 -4 s-1 [Karimpoor et al. 2006]

Figure 14: Fracture surface of nanocrystalline cobalt following tensile testing [Karimpoor et al. 2006]

24 2.1.3.4 Wear resistance Both sliding and abrasive wear resistance were found to improve for nanocrystalline metals in comparison to their polycrystalline counterparts [Suryanarayana et al. 2000] . In particular, nanocrystalline Ni deposits were found [Jeong et al. 2001] to show improved abrasive wear with respect to decreasing grain size, as shown in Figure 15. Previous studies [El-Sherik et al. 1997] found nanocrystalline Ni adhesive wear resistance and friction coefficient to improve by over 100 times and up to 50%, respectively, compared to their polycrystalline counterparts.

Figure 15: Taber wear index with respect to average grain size for Ni electrodeposits [Jeong et al. 2001]

25 Solid-solution and precipitation hardened electrodeposited nanocrystalline Ni-P linearly improved the abrasive wear resistance with increasing hardness, as shown in Figure 16, to a much greater extent and by purely reducing grain size alone [Jeong et al. 2003].

Figure 16: Vickers hardness as a function of taber wear index for nanocrystalline Ni and Ni-P electrodeposits [Jeong et al. 2003]

The addition of Co to nanocrystalline Ni electrodeposits was found to decrease the coefficient of friction from 0.45-0.5 to 0.25 as the Co concentration in the deposit increased to 70% [Ma et al. 2013]. The authors concluded that these results were due to the layer of HCP-Co wear particles acting as a solid lubricant, or tribofilm, in the pin-on-disc tests [Ma et al. 2013].

Electrodeposited nanocrystalline cobalt-phosphorus alloys have been offered as a replacement for hard chrome coatings in effort to eliminate the use of hexavalent chromium in electroplating processes [McCrea 2010]. The Co-P alloys offer comparable or improved mechanical, corrosion and wear properties to hard chrome coatings, such as similar hardness (up to 680 VHN), increased ductility (5-7%), reduced wear loss volume (6-7 x

26 10-6 mm3/Nm), reduced coefficient of friction (0.4 – 0.5) and pin-on-disk wear, and a 4- fold improvement in corrosion resistance [McCrea 2010].

While the hardness of nanocrystalline Co, with its average grain size remained unchanged, was shown to increase with added phosphorus due to solid solution hardening mechanisms, the wear resistance of such materials did not increase linearly, but rather was reportedly affected by cobalt oxide wear particles that were re-deposited on the sliding wear track surface [Alanazi et al. 2015].

2.1.3.5 Corrosion resistance The corrosion resistance of some nanocrystalline materials has been shown to be superior to their polycrystalline counterparts [Kim et al. 2002; Li-yuan et al. 2010; Srivastava 2006; Wang et al. 2006; Youssef et al. 2004]. In particular, nanocrystalline Zn coatings for galvanization of steel have shown improved passivation kinetics and passive layer stability compared to typical electrogalvanized steel in potentiodynamic polarization tests in NaOH. The Zn coating, although with etch pits present, also showed an overall lower corrosion rate than the electrogalvanized steel that had a more uniform corrosion morphology [Youssef et al. 2004]. Additionally, improved corrosion behaviour has been observed for nanocrystalline Ni and mixed HCP-FCC NiCo in a number of studies [Kim et al. 2002; Li- yuan et al. 2010; Srivastava 2006; Wang et al. 2006]. A reduction in grain size from 8um to 12nm for electrodeposited Co showed little change in corrosion resistance in Na2SO4 solutions following potentiodynamic polarization tests and surface morphologies were similar and showed uniform degradation. However, an aggregation of sulfur solutes was predicted on the corroded nanocrystalline Co surfaces and annealed nanocrystalline Co, although with identical passivity, demonstrated preferential attack along grain boundaries owing to the S accumulation [Kim et al. 2003]. No improvements in passivation were seen for nanocrystalline copper in NaOH as the grain size was reduced from 3um to 45nm and similar surface morphologies were observed for all tested materials [Yu et al. 2007].

27 Pulse-current electrodeposited nanocrystalline Ni-P layered coatings of 4.3nm average grain size were observed to have severe interlaminar cracking and pitting in NaCl solutions, where preferred Ni dissolution occurred leaving passive P-rich layers, accelerated by temperature increase. Deposit layers of 50nm thickness and with expected alternating P levels was concluded to provide a transverse pathway for the NaCl solution and thus accelerated the degradation of material [Lee et al. 2010]. Both nanocrystalline Co and Ni have been tested in alkaline and acidic solutions and it has been found that while enhanced passivity was observed in alkaline conditions, high corrosion rates and pitting corrosion morphologies were observed in acidic HCl [Li- yuan et al. 2010; Wang et al. 2006]. Similar findings were also observed for nanocrystalline Co-P electrodeposits, which were found to be less passive than amorphous Co-P and showed less uniform degradation morphologies. Active-passive behaviour that was seen in

NaOH solutions for both materials was not observed in H2SO4 conditions, where no passive behaviour was found [Sheikholeslam et al. 2010] .

28 2.2 Nanostructured Electrodeposited Cobalt

2.2.1 Crystallographic Structure Nanocrystalline Co with an average grain size diameter of 7nm and prepared by gas condensation has shown mixed 30% ordered-70% disordered atoms, owing to intercrystalline and crystalline atom contributions [Babanov et al. 1995]. Karimpoor and Erb [2003] characterized the crystallographic structures of electrodeposited nanocrystalline Co and polycrystalline electrowon Co (produced by hot and cold rolling, and annealing up to 1000 C in Ar in an electric tube furnace) by means of X-ray diffraction, scanning electron microscopy (SEM), and bright and dark field transmission electron microscopy (TEM) images and diffraction patterns. They found mixed FCC-HCO and pure HCP structures for polycrystalline (average grain size 4.8um) and nanocrystalline (average grain size 12nm) cobalt samples, respectively [Karimpoor et al. 2003]. Some investigations [Aus et al. 1998; Karimpoor et al. 2002,] of electrodeposited nanocrystalline Co via TEM imaging observed a fully dense hcp material with strong <0002> texture. However, other investigations [Fellah et al. 2010; Wu et al. 2005] have also observed martensitic FCC to HCP phase transformations and mixed HCP – FCC structures in nanocrystalline and ultrafine-grained cobalt, produced by flame-spray-derived cobalt nanopowders [Fellah et al. 2010] and electrodeposition [Wu et al. 2005]. The austenitic phase tranformation (HCP to FCC) is noted as a function of heating rate:

As = 450C + 0.28b, where b is heating rate in C/min [Ray et al. 1991]. Zhang et al. [2006] noted that despite XRD peak narrowing following cold-rolling deformation of cobalt, the grain size was not coarsened pre- to post-deformation nor were any SF's or dislocations observed in the deformed cobalt following TEM imaging. They explain this phenomenon through vacancy activity rather than dislocation or SF and twinning deformation mechanisms. Mainly, internal stress reportedly caused vacancies and vacancy clusters nucleate to mediate deformation caused by atom displacement along GB's and within grains at later deformation stages. Interstitial defects also increase the number of atomic planes that contributes to XRD peak broadening (similar to broadening by a high

29 density of SF's and dislocations). Zhang attributed the XRD peak narrowing to a large density of vacancy movement following strain unloading. Hibbard et al [2001] found that nanocrystalline cobalt had a higher activation energy (1.1 J/m2 specific excess interfacial enthalpy) for grain growth than that for nickel, attributed to boundary diffusion as the rate-limiting step for grain growth. Alloying of nanocrystalline cobalt with C and Cu was found [Bachmaier et al. 2015] to improve thermal stability of nanocrystalline cobalt. This is in contrast to typical nanocrystalline metals exhibiting low thermal stability owing to enthalpy stored in the higher GB area (compared to polycrystalline metals) if grain boundary migration is not impeded (in this study by means of alloying). Studies [Hyie et al. 2012] of Co alloyed with Ni and Fe found that alloying with both elements (FCC) increased the corrosion resistance and microhardness compared to pure cobalt (HCP) or that alloyed with one constituent (Fe), resulting in decreased average grain size (~72nm pure Co compared to 40nm CoFe and 35nm CoNiFe). Preferred orientation of 2.5mm thick nanocrystalline Co was found to change from (011 ̅1) to the (0002) texture as the thickness of the deposit increased, suggesting that with deposit growth the basal plane is preferentially oriented parallel to the deposit surface [Karimpoor et al. 2007]. This evolution is shown in Figure 17 [Karimpoor et al. 2007].

Figure 17: XRD patterns of nanocrystalline Co from cathode facing surface (a), mid-section (b), and free surface (c) [Karimpoort et al. 2007]

2.2.2 Deformation Mechanisms Cobalt has a low stacking fault energy (SFE) of 27 ± 4 mL/m2 [Fellah et al. 2010; Korner et al. 1983; Wu et al. 2004]. This has been observed as a lamellar structure in nanocrystalline cobalt material [Fellah et al. 2010; Karimpoor et al. 2003]. The lamellar structure has also been attributed to the presence of twins [Karimpoor et al. 2003; Hibbard et al. 2002] and HCP-FCC platelets [Farhangi et al. 1989]. Preferentially mechanical twinning is known to occur in polycrystalline cobalt and twins are also predominant in HCP nanocrystalline metals [Karimpoor et al. 2003].

Wu et al [2005] noted that twinning occurs early for HCP metals in addition to dislocation slip deformation mechanisms to satisfy the von Mises criterion. They attributed the large presence of stacking faults in HCP cobalt to being caused by the glide transformation of partial dislocations on closed packed planes during the FCC gamma to HCP epsilon phase of Co. They claim that there are three basal plane stacking faults possible that formed during the above-mentioned phase transformation. Twinning was observed in single crystals along the {1012}, {112̅2} and {112̅1} families of planes, with

31 the main mode of low level strain accommodation along the {10 ̅11} planes in HCP grains. FCC grains were dominated by dislocation slip deformation mechanisms.

Zheng et al. [2005] simulated deformation mechanisms in randomly oriented nanocrystalline cobalt (average grain size 10.4nm) composed of SF as well as full and partial dislocation activities rather than twinning mechanisms when deformed at a strain rate of (~1 x 108 s-1). Shockley partial dislocations (1/3 <1100>) were observed in the basal plane; no critical grain size was found where full dislocation slip transitions to partial dislocation slip as per nanocrystalline FCC metals like Ni and Al. Zheng et al. [2005] also noted that a lamellar structure is attributed to SF ribbons with FCC phases in HCP grains (deformation-induced phase transformation at high strain levels), which may restrict dislocation slip to further induce strain hardening and increase ductility of nanocrystalline HCP metals.

Wu et al. [2005] attributed the lamellar or 'platelet' structure of Co to the martensitic phase transformation from FCC to HCP structure with some platelets attributable to twins and intermediate regions of twins and epsilon martensite (not faulted austenite since HCP phase only). They noted that an increase in strain forced the alpha- to-epsilon transformation. However, the group claimed that the critical resolved shear stress for twinning increases more significantly than that for dislocation slip with increasing strain for reduced grain sizes. This would signify that the main deformation mechanism for nanocrystalline grained cobalt may be dislocation slip and not twinning as previously reported.

Fellah et al. [2010] noted that an increase in nanoscale twins resulted in an improvement of mechanical properties of UFG metals. For example, the interfaces introduced by a Co-Cu lamellar structure studied were assumed to act as coherent twin boundaries that enhanced mechanical properties. The group investigated a highly faulted plated microstructure with a large number of SF's and dislocation contrasts and voids owing to the powder metallurgy formation process. They showed that the lamellar boundaries were FCC-FCC twin boundaries and FCC-HCP phase interfaces. Fellah et al also noted

32 that a reduction in final porosity reduced the presence of the lamellar structure that had high faulting tendency. They attributed a strengthening effect to the boundaries present in the lamellar structure and likened them to grain or coherent twin boundaries. In particular, a noteworthy conclusion was the strength of the microstructure was controlled by the thickness of lamellae rather than the size of grain in which there were found. Similar to Wu et al., Fellah et al. noted that an increase in strain resulted in more HCP than FCC phase to be present but that the main deformation mechanism was through twinning. The FCC to HCP transition was explained by Shockley partial emission and gliding or the HCP lamellae growing in an FCC-structured grain.

Morrow et al. [2014] studied polycrystalline HCP magnesium and found twinning to be the main deformation mechanism. High resolution TEM analysis showed twin boundaries at the basal plane aligns with prismatic plane to create a facet and that the faceted boundary allows for twinning dislocation climb along with more typical twinning dislocation glide.

Karimpoor et al. [2003] attributed a highly-faulted microstructure to the presence of stacking faults and twins introduced by cobalt’s low stacking fault energy (SFE). Karimpoor et al. found that in regards to tensile deformation of nanocrystalline metals the strain rate influences the ultimate tensile strength and the flow stress, which both increased with decreasing strain rates. They attributed the increase in strain rate to increase the ultimate tensile strength to dislocation slip for polycrystalline cobalt and for twins present in nanocrystalline cobalt to decrease the flow stress and tensile strength with increasing strain rate. They claim that twinning required a higher activation stress than that required for dislocation movement, which then proceeds with smaller stress increments. This is comparable to conclusions made by Chan et al. [2014], who found a strain rate dependency/sensitivity for nanocrystalline Ni and Ni-Fe electrodeposits that was not present with coarse grained Ni in terms of yield and ultimate tensile strengths.

33 Karimpoor et al. [2003] concluded that increases in stress levels is owed to both heterogeneous and homogeneous (in grain interiors involved overlap of SF's of dissociated dislocations) nucleation of twins which can occur at grain sizes less than 50nm. This is comparable to FCC metals such as copper with low SFE's, which still have a higher tendancy to deform by dislocation slip rather than twinning [Christian and Mahajan 1995]. The main deformation mechanism of nanocrystalline metals is not dislocation dependent and has been well documented as grain boundary-controlled (GB sliding/rotation) [Chan et al. 2015; Li 1962; Luthy et al. 1979; Shi and Zikry 2009; Van Swygenhoven and Derlet 2001] for nanocrystalline metals with average grain sizes less than 10nm. So nanocrystalline metals with average grain sizes near 10nm may incorporate both dislocation and grain boundary controlled deformation mechanisms. Rajagopalan et al. [2011] found that an increase in homogeneity of nanocrystalline aluminum grains results in higher yield strength values for uniaxial tensile testing.

2.2.3 Multilayered Materials Multilayered Co-X systems have been studied for their change in performance criteria associated with layer properties. For example, Co-Pt multilayered systems have been investigated [Lacey et al. 1990; Poulopoulos et al. 1995,] for their structural and magnetic and magnetoresistive properties, and it has been found that the magnetic properties (eg. perpendicular anisotropy) of the material are dependent on both individual Co layer thicknesses and Co concentration within the alloyed layers. Gomez et al. [2002] found that a Co-Cu multilayered system (layer thickness of 180-200nm) showed distinct layer separation under SEM imaging and that magnetoresistance of the structure increased with decreasing Co layer thickness, down to 1nm Co layers where the continuity of the layer was not observed. TEM analysis [El Fanity et al. 1998] of cross-sectioned multilayered electrodeposited polycrystalline Co-Cu films have shown columnar grain growth between defined layers of 6nm and 4nm FCC Co and FCC Cu, respectively, and that substrate roughness had a direct result on layer deformation and film surface profile.

34 Hong et al. [2006] studied the relationship between the addition of organic substances (sulfopropyl disulfide sodium salt or dimethyldithiocarbamic acid) and Co-Cu multilayers plated in electrochemical solutions. They found that an addition of approximately 0.5mmol/L resulted in more defined layer interfaces (Cu-10nm and Co- 42nm layer thicknesses) and a shift from HCP to FCC-structured Co.It is therefore evident that the electrodeposition parameters ultimately influence deposit performance. Co-Ru bilayers have been investigated [Michel et al. 1996] and found that a hexagonal lattice misfit existed at the interface, which was attributed to the layer interface structure as an important causal factor in the material's change in magnetic anisotropy, particularly with Co layer thicknesses at 1.5nm. The interface structure was also a proposed influence on the material's magnetoresistance.

Nanocrystalline and polycrystalline or coarse-grained electrodeposited NiCo alloys have been multilayered in a 1:1 thickness scheme and found by Daly et al. [2015] to combine the ductility of the coarse grained layer with the improved strength of the nanocrystalline layer in a sandwich-type structure following a rule of mixtures relationship. Fracture surfaces of uniaxial tensile tested coupons exhibited periodic features of coarse dimpled protrusions amongst fine dimpled intermediaries, both products of microvoid coalescence, where the coarse grained layers were shown to offer an increase in tensile strain or elongation through improved necking stability.

2.2.4 Mechanical Properties Electrodeposited cobalt has shown to be a favorable method of nanocrystalline cobalt production in terms of its ease of research-to-production manufacturing and ability to produce near-net-shape products. Karimpoor et al. [2003] investigated the performance and deformation mechanisms of nanocrystalline cobalt. It was reported that cobalt was expected to have a lower ductility than nickel owing to a reduced number of slip systems for its hexagonal closed packed (HCP) crystal structure in comparison to nickel’s face- centered cubic (FCC) structure. However, higher ratios of nanocrystalline to

35 polycrystalline tensile elongation was achieved for Co than for similarly produced Ni electrodeposits. The deformation mechanisms for HCP nanocrystalline Co metals reportedly included dislocation slip, diffusional creep, grain boundary sliding and twinning. Karimpoor et al. [2003] compared polycrystalline Co at 25% FCC and 75% HCP structures and average grain size of 4.8 ± 0.2um to nanocrystalline HCP-only Co material. The phase stability in electrodeposited polycrystalline Co was attributed to the electrodeposition parameters, where organic surfactants, presence of FCC-structure metal ions, and co- deposition of hydrogen at the cathode were linked to an increase in FCC favoured deposition [Dille et al. 1997; Morral et al. 1974] despite post-processing treatments. Karimpoor et al. [2003] found that a reduction in average grain size from 4.8um to 12nm resulted in an increase in yield and ultimate tensile strength and a slight reduction in Young's modulus (from 212-223 GPa for polycrystalline Co to 205-209 GPa for nanocrystalline Co). The reduction in Young's modulus was partly attributed to the difference in increased volume fraction of intercrystalline components or to change in crystallographic structure, though the yield and ultimate tensile strength were not discernable as dependent on grain size or on crystallographic structure. Nanocrystalline Co tensile tested at three different quasi-static strain rates exhibited different tensile elongation values, yield strength, ultimate tensile strength, and work hardening exponent. The lowest strain rate resulted in higher flow stress and tensile strength contrary to what is expected for dislocation-controlled deformation mechanisms, which suggests that mechanical twinning was the major deformation mechanism present [Karimpoor et al. 2003]. The stress-strain curves and values for these properties are reproduced in Figure 18 and Table 4. Minor discrepancies were also observed for polycrystalline Co, except for larger variations in tensile elongation. At 99.5% purity, polycrystalline cobalt that has been hot worked and annealed at 800C - 1000C has been observed at an elongation of 15-30% [ASM International 2007].

Figure 18: Stress-strain curves for polycrystalline and nanocrystalline Co at strain rates of 1 X 10-4 s- 1 (1), 5 X 10-4 s-1 (2) and 2.5 X 10-3 s-1 (3) [Karimpoor et al. 2003]

Table 4: Tensile Properties of polycrystalline and nanocrystalline Co at different strain rates [Karimpoor et al. 2003]

Fracture surfaces for nanocrystalline Co had a mixed/slanted plateau with ledges oriented at 37-53, as shown in Figure 13. This is in comparison to polycrystalline Co fracture surfaces, which were oriented perpendicular to the fracture surface [Karimpoor et al. 2006]. Both specimen types exhibited dimpled fracture surfaces indicative of ductile fracture. Nanocrystalline Co produced finer dimples or microvoids. The room temperature Charpy impact energy of nanocrystalline (18nm average grain size) cobalt was investigated [Karimpoor et al. 2007] and was found to be four times lower than that of annealed (1um average grain size) polycrystalline cobalt with a microhardness about twice as high. It is noted that the modulus of toughness values derived from the researchers’ previous study [Karimpoor et al. 2003] showed similar grain sized cobalt to have high elongation values (9%) and high tensile strength up to 2200 MPa.

2.2.5 Applications

2.2.5.1 Wear Resistance and Tribological Behaviour Co has been investigated as a potential hard chromium replacement for its desirable wear and corrosion resistant properties [Hibbard et al. 2001]. Investigation of the tribological behavior of electrodeposited nanocrystalline and polycrystalline Co and Co-based alloys is imperative to predict its performance in high-wear applications. A comparison [Ma et al. 2015; Wang et al. 2006] of electrodeposited, pulsed current nanocrystalline Co and Ni wear properties showed that with comparable grain sizes (16nm Ni grains and 18nm Co grains, averaged), Co exhibited less visible coating wear damage, reduced friction coefficient, and improved wear resistance by an order of magnitude. Wear rates were also shown to improve with reduction in grain size from ~4.25 x 10-5 mm3/Nm at 2.5um to ~3.5 x 10-5 mm3/Nm at 18nm. The authors credited cobalt’s wear resistance to its hexagonal structure and associated resistance to adhesive wear. Cobalt-based alloys are selected for many human contact applications over nickel materials owing to their reduced metal sensitivity [Brandao et al. 2012]. Alloys that are selected for biomedical implants may be subjected to metal-metal interfaces in high wear locations. These interactions have been addressed as potential causes of hypersensitivity and elevated metal particles in the blood and urine of patients [Spriano et al. 2005]. As the effects of these particles have yet to be fully realized, Co or Co-based alloy coatings with high surface wear resistance properties are desirable [Holecek et al. 2009; IARC 1990; Pourzal et al. 2011]. Weston et al. [2009] have investigated electrodeposited Co and their alloys as considerations for hard chromium replacements in the automotive and aerospace industries. Weston et al. showed that nanocrystalline Co-W coating material had a reduced wear rate than Cr equivalents against 440C martensitic steel counterbodies by an order of magnitude for high loads (61N vs. comparable wear rates at 30N). However, monolithic pure Co coatings were observed to have the highest wear rate out of the materials tested

38 (i.e. decreasing wear rates were found in the order of Co, Cr, CoW). Multilayered nanocrystalline Co structures have not yet been compared in such studies. In a comparative study of nanocrystalline, HCP Co produced by four electrodeposition methods Su et al. [2013] observed decreasing wear rates in the order of pulse reverse current, direct current, pulse current, bipolar pulse plating from ~8.5 x 10-5 mm3/Nm down to ~2 X 10-5 mm3/Nm, respectively (against GCr15 steel counterbodies and an applied load of 5.0 N). This order coincided with the decreasing surface roughness of each film. Overall, the authors found the tribological behavior of the Co films to be dependent on their respective hardness, surface roughness, phase structure and morphology. The addition of alloying elements to form composite structures such as Co-GO (graphene oxide) [Lie et al. 2015] was found to reduce average grain size from 50± 5 to 20 ± 2nm and increase microhardness from 340 ± 10 kgf/mm2 to 430 ± 15 kgf/mm2 with improved wear and corrosion resistance.

2.2.5.2 Corrosion Resistance As discussed earlier, there are mixed conclusions on corrosion improvements of nanocrystalline and polycrystalline metals and alloys. Generally, electrodeposited nancrystalline metals showed overall uniform morphologies following degradation and were found to be dependent on the corrosion conditions, i.e. in alkaline or acidic solutions.

Kim et al. [2003] studied the corrosion behaviour of polycrystalline and nanocrystalline grained cobalt by potentiostatic polarization studies in sodium sulfate and found that both materials exhibited no passivity with no preferential grain boundary dissolution, save for preferential GB dissolution observed in annealed nanocrystalline Co which was attributed to the accumulation of sulfur impurities along GB's.

39 3 EXPERIMENTAL

3.1 Electrodeposition Electrodeposited Co foils of ~150 m and ~500 m average thickness were received from Integran Technologies Inc. and produced by methodologies described elsewhere (US Patent # 5,352,266 and 5,433,797). Foils were electrodeposited in a single electrolytic solution containing cobalt salts including cobalt sulfate and cobalt chloride, and with sulfur-bearing organic additives at temperatures of ~ 60 C and pH 2 - 4. Foils were mechanically stripped from substrates. Electrodeposits were formed under pulse waveforms which nominally would translate to three deposit structures, as shown in Figure 19: (1) Monolithic Co (MN) (2) Multilayered Co with nominal 20 nm sub-layer thickness (ML20) (3) Multilayered Co with nominal 100 nm sub-layer thickness (ML100)

(1)

(2)

(3) Figure 19: Schematic diagrams of (1) MN, (2) ML20 and (3) ML100 electrodeposited Co

40 The exact pulse waveforms are not disclosed as they are proprietary waveforms developed by Integran Technologies Inc. Multilayered Co was deposited using two different electrodeposition plating conditions to produce electrodeposits with comparable bulk thicknesses and with nominal sub-layer thicknesses at 20 nm and 100 nm (ML20 and ML100, respectively). The individual sub-layers were electrodeposited under varying current conditions in the same electrolytic bath. Sub-layeres were deposited at a 1:1 thickness ratio. For example, a deposit at 100 m bulk thickness and 100 nm sub-layer thickness would nominally consist of 1,000 sub-layers; 500 sub-layers of each plating condition.

3.2 Characterization

3.2.1 X-ray Diffraction The crystallographic structures of bulk electrodeposited specimens were analyzed via X- ray diffractometry (XRD) using a Rigaku MiniFlex 600 with /2 geometry. FCC and HCP Co reference peaks used are shown below in Figure 20. The specimens were analyzed in ‘as-plated’ conditions and measurements were taken from the cathode-facing surface.

Figure 20: Co X-ray diffraction patterns for reference FCC and HCP Co (Cu K ,  = 1.5418 nm)

41 3.2.2 Bulk Alloy Composition The purity of bulk Co foils were investigated using X-ray Fluorescence (XRF) in a Bruker S2 Ranger. The specimens were analyzed in ‘as-plated’ conditions. Carbon and sulfur concentrations of each received foil were determined as per ASTM E1019-11. The specimens were analyzed in ‘as-plated’ conditions.

3.2.3 Scanning Electron Microscopy Deposit cross-sections and fracture surfaces of tensile tested coupons were imaged using Scanning Electron Microscopy (SEM) with Hitachi S-4800, SU3500 and SU8230 instruments.

3.3.3 Transmission Electron Microscopy The grain size and microstructure of Co foils were investigated using bright field/dark field imaging and selected area electron diffraction in a Hitachi HF3300 Environmental-CFE- TEM operating at 300 kV. Electrodeposited foils were prepared by dual jet electropolishing (MN) and microsampling (ML20 and ML100) using a Hitachi NB5000 FIB instrument. MN foils were mechanically thinned and disc-punched (3mm diameter) then ground and polished using 400, 600, 800 and 1200 grit SiC papers. Dual-jet electropolishing was conducted with a Struers TenuPol-5 in an 80% methanol-20% perchloric acid solution at

30V. The solution temperature was lowered to approximately -40ºC with liquid N2. Cross-sectional and below-fracture surface cross-sectional TEM (XTEM) samples were prepared from tensile tested coupons using FIB microsampling approximately 100- 200 µm away from the fracture ledge on substrate/cathode-side surface. Microsampling was completed by milling a trough around the desired area and followed by sample ‘lift- out’ and thinning, as shown in Appendix A. A protective W layer was deposited onto the surface of the specimen prior to milling.

42 3.3 Properties

3.3.1 Microhardness Vickers microhardness testing was completed along foil cross-sections with a load of 100g and dwell time of 10s. Cross-sectioned samples were first ground at 500, 800 and 1000 SiC grit followed by 5um, 2um and 1um diamond polishing. Hardness measurements were taken at a minimum of 5 points for all samples across sample thickness.

3.3.2 Tensile Testing MN, ML20 and ML100 dog-bone tensile coupons were waterjet cut from bulk ~500um thick foils received from Integran Technologies Inc. The test coupon geometry is as shown in Figure 21. Tensile testing was completed using an Instron machine with a maximum load of 500 kN at a strain rate of 5 x 10-4 s-1. The tensile coupons were first polished and then speckle spray-painted for Digital Image Correlation (DIC) image captured by a camera system.

Figure 21: Tensile coupon measurements (mm)

44 4 RESULTS AND DISCUSSION

4.1 Sample Identification Identification of the monolithic and multilayered Co specimens is outlined in Table 5. The purity of each of the Co foils was confirmed at >99% using XRF.

Table 5: Sample identification and bulk purity analysis via XRF

Sample Description Bulk Purity (Elemental, %) MN Monolithic Cobalt Co: 99.4; Fe: 0.233; Mn: 0.0919; S: 0.0913; Si: 0.208 ML20 20nm nominal sub- Co: 99.3; Fe: 0.254; Mn: 0.107; S: 0.0815; layer thickness Cobalt Si: 0.207 ML100 100nm nominal sub- Co: 99.2; Fe: 0.279; Mn: 0.184; S: 0.0820; layer thickness Cobalt Si: 0.213

4.2 Crystallographic Structure X-ray diffraction pattern comparisons of polycrystalline and nanocrystalline Co have shown peak broadening, indicative of grain size refinement [Karimpoor et al. 2003]. Nanocrystalline Co X-ray Diffraction (XRD) patterns have been shown to display significant peak broadening with reduced grain sizes. Peak overlap is possible from HCP and FCC structures at certain peaks, as identified in Figure 20. Typical polycrystalline Co electrodeposits show strong (0002) texturing with basal plane oriented parallel to the surface of deposit and electrodeposited nanocrystalline Co has been observed to have a hexagonal crystal structure [Aus et al. 1998; Karimpoor et al. 2002; Karimpoor et al. 2003]. X-ray diffraction patterns of MN, ML20, and ML100 specimens are shown in Figure 22.

Figure 22: X-ray diffraction patterns of MN, ML20, and ML100 specimens using Cu-K radiation

Reference XRD patterns for Co-225 (FCC) and Co-194 (HCP) are reproduced in Figure 20. As shown in Figure 20, the FCC (111) and HCP (0002) peaks for Co overlap as well as the FCC (220) and HCP (112̅0) peaks. However, for mixed FCC-HCP electrodeposited Co, a distinct FCC (200) Co peak is commonly observed, which is not present in the patterns for the investigated monolithic and multilayered specimens [Karimpoor et al. 2001].

46 As seen in Figure 22, the patterns for MN, ML20 and ML100 all show characteristic HCP peaks. However, there is an observed change in preferred orientation to the (0002) peak for layered specimens, with increasing (0002) peak intensity as the nominal layer thickness increases from 20nm to 100nm. The (0002) peak in multilayered structures has a strong basal plane preferred orientation aligned parallel to the deposit surface, in comparison to the randomly textured HCP XRD pattern shown in Figure 20. The monolithic specimen has a weaker basal plane texture than the multilayered specimens.

Change in preferred orientation for electrodeposited nanocrystalline Co has been observed with increasing deposit thickness in tested specimens up to 2.5mm thick samples [Karimpoor et al. 2007]: (0002) preferred orientation was observed in deposits that were at least 1.5mm thick. The onset of the texture change has yet to be linked to a specific deposit thickness; i.e. the deposit thickness at which the preferred orientation to (0002) is unknown.

4.2.1 Monolithic Cobalt Monolithic cobalt foils were prepared for TEM imaging by dual jet electropolishing as described in Section. Specimens were kept under vacuum storage and were UV vacuum cleaned immediately prior to imaging. Bright field (BF) and dark field (DF) TEM images of MN specimens are shown in Figure 23, below. The diffraction pattern (DP) is shown as the inset on the DF image. The inner three HCP rings (10 ̅10), (0002), (10 ̅11) and the (10 ̅12), (112̅0), (10 ̅13) and (202̅0) rings are shown in increasing diameter from the transmitted beam, respectively.

(a) (b)

Figure 23: BF (a) and (c) and DF (b) and (d) TEM images and (e) SADP inset (L=300) of MN Co

48 Grain size measurements were taken from >200 distinct grains in bright field and dark field images in the diffracting condition of monolithic Co. The log-normal grain size distribution is shown in Figure 24. The average grain size was measured at 14 ± 7 nm.

Figure 24: Grain size distribution for MN Co with log-normal distributiom

The TEM images show a large density of alternating contrast fringes, indicated with white arrows in Figure 23. These artifacts may be a result of Moiré fringes, faulted structures or twins, or interference fringes from slanted surfaces of grain boundaries. Near- perforation BF and DF images were taken in an attempt to capture a single grain to rule out Moiré contrast as a potential cause of the majority of fringes observed in these images, shown in Figure 25. The DF image was produced using a tilted beam for selected diffraction around the (10 ̅10) HCP plane, however due to the minimum aperture size available, the DF image likely includes diffracted signals from nearby (0002) and (10 ̅11) rings.

49 Moiré fringes occur when grains of similar orientation are overlapped to produce small differences in periodicities. Images captured near perforation of the MN sample also showed a high density of fringes, observable in both BF and DF images. As previously discussed, electrodeposited HCP cobalt has a predicted high density of deformation twins and has low stacking fault energy in comparison to other metals that favour deformation twinning. Based on literature and Co’s low SFE, it is reasonable that fringes observed in both BF and DF images may be a result of a faulted HCP Co structure, Moiré fringes, or surface effects of grains, or combinations thereof. No excessive columnar growth was observed in MN Co microstructure images.

Figure 25: BF and DF images of grain near perforation in MN Co (DF image produced by selected diffraction around the (100) plane

50 4.2.2 Multilayered Cobalt SEM images of nominal multilayered Co specimens did not resolve layer interfaces or differences at sub-layer thicknesses of 100nm or less, as shown in Figure 26. Sub-layers are oriented horizontally in Fig. 26. An additional Co electrodeposit produced by the same methods as ML20 and ML100 was received with nominal 500nm sub-layer thicknesses. Cross-sectional SEM imaging of this deposit was able to resolve sub-layers, as shown in Figure 27. There are detection limits to sub-layer resolution in SEM owing to reduced contrast from preferred orientation or compositional variations between layers. Therefore, SEM imaging was determined as not feasible for imaging of sub-layers for ML20 and ML100 deposits.

Multilayered Co of alternating layers produced Integran Technology Inc. proprietary waveform conditions were prepared for XTEM imaging via FIB microsampling, as previously discussed in Section 3. Bright field images of multilayered Co with nominal layer thicknesses of 20nm (ML20) and 100nm (ML100) are shown in Figure 28 (a) and (b), respectively. In the cross-sectional images obtained from these methods, no indication of differences in layer thickness or, more simply, layer presence was observed. According to sampling methods, the image orientation in Figure 28 would have layers deposited in the direction of increasing nominal thickness, t0, of 20nm and 100nm, respectively. Both ML20 and ML100 Co have a high density of fringes similar to that observed in MN Co and no excessive columnar growth was observed.

Chan [2011] observed a discrepancy in nominal layer thickness and actual observed thickness for iron electodeposits produced in an iron-sulphate electrolyte to thicknesses of ~80-100 m using pulse waveforms developed by Integran Technologies Inc. SEM and TEM imaging of multilayered iron electrodeposits were unable to establish 100 m nominally thick individual sub-layers that were previously observed at nominal thicknesses of 10 m to 250 m [Chan 2011]. The minimum observable layer thickness was attributed to two potential causes: (1) reduced uniformity, particularly in regions of high current density (eg. dendritic, tree-like electrodeposit growth around deposit edges) [Péter et al.

51 2001] and (2) the requirement of a minimum layer thickness to form a continuous layers owing to the Volmer-Weber growth mechanism, which describes nucleation and growth in the electrodeposition process [Watanabe 2004].

푴푳ퟏퟎퟎ

풕ퟎ

Figure 26: 100nm layered Co deposit (X50.0K); scale bar represents 1m; deposit thickness is in the vertical direction, t0

풕ퟎ

Figure 27: 500nm layered Co deposit (X10.0K); scale bar represents 5m; deposit thickness is in the vertical direction, t0

푴푳ퟐퟎ

풕ퟎ

(a)

푴푳ퟏퟎퟎ

풕ퟎ

(b)

Figure 28: (a) ML20 and (b) ML100 BF TEM images; deposit thickness is in the vertical direction, t0

53 Log-normal grain size distributions for ML20 and ML100 are shown in Figures 29 and 30, respectively. Grain size measurements were taken from >200 distinct grains in bright field and dark field images in the diffracting condition for both specimens. The average grain sizes are 11 ± 9 nm and 10 ± 5 nm for ML20 and ML100, respectively.

Figure 29: Grain size distribution for ML20 Co with log-normal distributiom

Figure 30: Grain size distribution for ML100 Co with log-normal distributiom

54 Diffraction patterns were taken for ML20 and ML100 Co across the width of the microsamples. Textured patterns were observed for both multilayered Co samples and from regions of the monolithic Co. Examples of commonly seen textured diffraction patterns are shown in Figure 31. Texture was observed for all three specimens along segments of the (0002) and (10 ̅11) rings.

Figure 31: DP of MN(a) ML20 (b) and ML100 (c). Textured regions along the (0002) and (10 ̅ퟏ1) rings are circled in (b) and (c)

4.2.3 Solute Concentration Use of sulfur-bearing organic additives in electrolytic solutions for Co electrodeposition has produced deposits with sulfur and carbon solutes, which were investigated for their effect on grain size and crystallographic texture [Hibbard et al. 2006]. The concentrations of carbon and sulfur in electrodeposited foils from this investigation are shown in Table 6.

Table 6: C and S concentration as determined via ASTM E1019-11.

Sample Carbon Sulfur (ppm) (ppm) MN 51.0 110 ML20 37.8 269 ML100 49.6 280

Both multilayered Co deposits exhibited higher sulfur content than the monolithic Co. Carbon concentrations were comparable for all three investigated materials. Co- deposited C and S is expected with sulfur-bearing organic additives in electrolytic solutions. Hibbard et al [2006] studied the effect of starting grain size and solute (sulfur and carbon) concentration on the thermal stability of nanocrystalline electrodeposited cobalt. The authors expected grain boundary mobility to be lower for evenly distributed sulfur solute atoms rather than carbon, as well as an increase in activation energy with an increase in bulk sulfur concentration in the deposits.

The solute concentrations from the Hibbard et al. [2006] study are shown in Table 7. The carbon solute concentrations in deposits studied by Hibbard et al. [2006] are an order of magnitude greater than those observed in the current study of MN, ML20 and ML100 Co samples. The XRD crystallographic patterns did not significantly differ across ten cobalt samples with varying sulfur and carbon concentrations, as observed in Figure 32. Sulfur concentrations were also much greater for the majority of deposits; sulfur concentrations at or less than 300ppm were only observed for 3 deposits in this study (Co-

56 4, Co-5, Co-6), all of which demonstrated a strong (0002) basal plane texture, shown in Figure 32. Sulfur concentrations measured were all above 200ppm for the materials investigated by Hibbard et al. [2006].

Hibbard et al. [2006] also found that the sulfur concentration in nanocrystalline Ni- Co electrodeposits was a significant factor in the deposit’s thermal stability owing to solute drag of sulfur impurities at migrating growth fronts.

Table 7: Solute concentration for sulfur [S] and carbon [C], grain size (d) and grain size range (r) and standard deviation (s) of nanocrystalline electrodeposited Co [Hibbard et al. 2006]

Figure 32: XRD patterns for electrodeposited nanocrystalline Co as discussed in Table 7 above [Hibbard et al. 2006]

57 Matsui et al. [2013] studied the effect of electrolytic additives in a sulfamate electrolyte on the tensile properties of nanocrystalline Ni-W deposits and found mixed conclusions regarding the influence of saccharin sodium on mechanical properties and crystallographic texture. In some cases, specimens had a relatively low tungsten level and a reduced grain size that was concluded to further hinder twin boundary formation. The presence of tungsten reduces the stacking fault energy (SFE) required for twin boundary formation.

A reduced grain size with fewer twin boundaries were observed to have the same hardness values as larger grains with more twin boundaries, so it was concluded that the presence of twins was a hardening feature with a similar effect to that of grain refinement (referred to by both the Hall-Petch effect and the Basinski mechanisms [Basinski et al. 1997] of hardening [Kalidindi et al. 2003]). However, the authors found no connection to the presence of twins or grain size to the tensile ductility of the deposits.

The sample deposited in the saccharin sodium bath had a strong FCC (200) texture in comparison to strong (111) texture in other deposits, which led the authors to conclude that the texture of the deposit was a key component in its resulting tensile ductility, along with comparatively higher residual stress in the (111) textured deposits [Matsui et al. 2014]. The authors concluded that deposit orientation and crystal growth modes must be examined to determine production methods of nanocrystalline electrodeposits with high tensile ductility.

These conclusions are supported [Schuler et al. 2013] by observations of grain refiners like saccharin to affect the crystallographic orientation of Ni deposits, along with other influencing factors like pulse parameters for pulsed current deposition. The effect of saccharin is explained, as it acts as a blocker once absorbed on the (111) Ni plane to hinder surface diffusion. Ni absorption then only occurs on the (100) planes. Newly generated sites on the (111) planes are continuously blocked and thus nucleation is promoted, so that the (111) surfaces grow and the (100) surfaces vanish, which results in a transfer from a (200) texture to (111) texture with increasing saccharin content in the electrolyte (as well

58 as acting as a grain refiner as the crystal shape moves away from the equilibrium shape, increasing internal compressive stresses) [Schuler et al. 2013]. An increase in current density increases the twin density and higher saccharin content decreases the twin density and contradicts the diffusion-based creep deformation mechanism [Schuler et al. 2013].

The solute concentrations observed are an indication of grain refiners used in the electrodeposition of monolithic and multilayered Co electrodeposits. Schuler et al. [2013] observed that an increase in electrolyte saccharin concentration from 0 g/L to 0.4 g/L resulted in a change in preferred orientation from (200) to (111) for FCC Ni. The results of this investigation indicate that there is a strong change in crystallographic texture with the introduction of a multilayers. Schuler et al. [2013] did not study the bulk sulfur or carbon concentration in deposits, which would have further explored the relationship between saccharin concentration in the electrolyte and bulk alloy composition of the Ni deposits.

However, Hibbard et al. [2006] found that the bulk sulfur and carbon concentrations had no effect on the crystallographic texture of nanocrystalline Co electrodeposits. The results of the current study align more closely with conclusions made by Hibbard et al. [2006]: although the addition of sulfur-bearing organic additives were used in the electrolytic solution for monolithic and multilayered Co electrodeposition, the bulk sulfur and carbon compositions do not differ significantly with the introduction of nominal multilayers and are not obviously linked to the change in preferred orientation. However, a bulk alloy composition analysis should be completed on a monolithic deposit produced by the second layering conditions to determine if there are significant differences in the two layers that comprise the Co multilayers.

59 4.3 Properties

4.3.1 Microhardness The hardness measurements for the monolithic and multilayered specimens are shown in Table 8. At least five measurements were taken across the thickness of the foils. and indent size transverses many sub-layers within the multilayered specimens.

Table 8: Vickers microhardness results, taken under a 100g load and dwell time of 15 seconds

Sample Hardness (VHN)

MN 432 ± 5

ML20 471 ± 9

ML100 462 ± 2

As shown in Table 8, the monolithic cobalt foils have the lowest measured hardness values. The 20nm and 100nm multilayered foils, ML20 and ML100, respectively, do not differ significantly in hardness. This may be due to the size of the indent that transverses many layers in each measurement, offering a bulk hardness reading. Previous studies of electrodeposited nanocrystalline materials found hardness values of >10% greater than the maximum hardness values observed in this investigation [Karimpoor et al. 2001]. The exact reason for this discrepancy is currently unknown but could be related to differences in grain size distribution, impurity content, and crystallographic texture. Additionally, the exact effect that nominal multilayering has on deposit hardness is not well-defined from these measurements alone; more multilayered specimens of incremental nominal layer thicknesses should be investigated to determine its effect. Nanoindentation may shed further light on this matter by measuring hardness of individual plated layers, assuming actual layer thickness to be greater than the indent size. Multilayered nanocrystalline NiFe electrodeposits showed increasing hardness with decreasing layer thickness but plateaus at layer thicknesses less than 100nm [Kurmanaeva et al. 2016].

60 Brooks et al. [2008] found that nanocrystalline materials follow the same hardness- strength relationship as their polycrystalline counterparts: HV = 3UTS with non-brittle nanometals, i.e. those which are able to reach a high enough ductility to avoid fracture before the UTS was reached. In a separate study, the authors determined that intrinsic ductility or the uniform plastic strain of electroformed nanocrystalline Ni deposits is independent of deposit microstructure within the grain size range of 10-80nm and that the interfacial damage nucleation and growth is best represented by a critical plastic strain or maximum intrinsic ductility [Brooks et al. 2011]. The increase in hardness from monolithic to multilayered Co specimens may be attributed to the slight reduction in average grain size from 14 ± 7 nm (MN) to 11 ± 9 nm and 10 ± 5 nm (ML20 and ML100, respectively), as explained previously by the Hall-Petch relationship regarding nanocrystalline Co [Karimpoor et al. 2003]. Although there is a slight difference in hardness from ML20 to ML100 Co, its influence by the inverse Hall- Petch relationship as explained in Section 1.1.3.1 is inconclusive. Microhardness tests are required from nominally multilayered Co specimens within a range of sub-layer thicknesses to determine if a relationship similar to that observed by Kurmanaeva et al. [2016] exists.

61 4.3.2 Tensile Testing Tensile testing was conducted on three samples of each specimen type: MN, ML20, and ML100 at a strain rate of 5 x 10-4 s-1. The results of three tested samples for each specimen are shown in Figures 33-35. Their stress-strain curves with the highest tensile strength are shown in Figure 36. Young’s Modulus, 0.2% offset yield strength (0.2%), ultimate tensile strength (UTS), fracture strength (fracture), tensile elongation, and strain-hardening exponent (n) properties obtained from this data are shown in Table 9.

Figure 33: Tensile test results of monolithic cobalt (MN) at a strain rate of 5 X 10-4 s-1

Figure 34: Tensile test results of 20nm multilayered cobalt (ML20) at a strain rate of 5 X 10-4 s-1

Figure 35: Tensile test results of 100nm multilayered cobalt (ML100) at a strain rate of 5 X 10-4 s-1

Figure 36: Average tensile test results of monolithic and multilayered cobalt at a strain rate of 5 X 10-4 s-

Table 9: Average mechanical properties obtained from engineering stress-strain curves

Sample E (GPa) 0.2% (MPa) UTS (MPa) fracture (MPa) Elongation (%) n

MN 136 ± 20 691 ± 36 1302 ± 44 1294 ± 48 3.58 ± 0.6 0.45 ± 0.03

ML20 166 ± 28 723 ± 68 1476 ± 35 1476 ± 35 4.51 ± 0.5 0.33 ± 0.01

ML100 164 ±18 703 ± 55 1498 ± 16 1448 ± 34 7.83 ± 0.6 0.42 ± 0.02

Young’s modulus (E) varies from the monolithic Co to both multilayered Co specimens. All measurements are lower than those previously reported for measured nanocrystalline Co in tensile testing, which were about 200 GPa [Karimpoor et al. 2003]. This reduction by about 20% in values may be due to variations in orientation, which is a known common effect on Young’s modulus values along with grain size reduction [Karimpoor et al. 2003; Zhou et al. 2003]. Grain refinement has been found to slightly reduce Young’s modulus for nanocrystalline electrodeposits in comparison to their polycrystalline counterparts. As previously discussed, this is in part owing to the large

64 increase in intercrystalline component volume fraction, or to a change in crystal structure as shown in comparing FCC-HCP Co to pure HCP Co. However, neither of these changes exist in the monolithic-to-multilayer nanocrystalline Co transition. Both multilayered Co specimens show nanocrystalline grains of with a similar grain size distribution as observed with the monolithic Co.

According to Table 9 the average 0.2% offset yield strengths, ultimate tensile strengths and fracture strengths are increased when the electrodeposited cobalt follows a nominal multilayered structure. Ultimate tensile strengths were calculated following Considère’s Criterion. The lowest strength values are observed for the monolithic cobalt electrodeposits. This agrees with the hardness measurements obtained. Again, the yield strengths, ultimate tensile strengths and fracture strengths for ML20 and ML100 deposits do not differ significantly. Most significant in the mechanical properties data is the large range of percentage tensile elongation for the three specimen types. The largest tensile elongation was seen with 100nm multilayered specimens, ML100, which reached an average of ~ 8%. This is over 40% greater ductility than that seen from the 20nm multilayered structure, ML20, and more than double that observed with monolithic Co.

Strain-hardening (or work-hardening) exponents, (n), were calculated from ASTM Standard E646-16 ‘Standard Test Method for Tensile Strain-Hardening Exponents (n – Values) of Metallic Sheet Materials’. The values obtained for strain-hardening exponents of both monolithic and multilayered specimens were two times greater than those previously observed for monolithic nanocrystalline Co, which were found to be 0.20 ± 0.01 at the same strain rate [Karimpoor 2002]. Generally, the work hardening rate has been observed to decrease for nanocrystalline materials compared to their polycrystalline counterparts with decreasing grain size, although this effect was not fully observed with nanocrystalline Co compared to polycrystalline Co [Karimpoor 2002; Karimpoor et al. 2003].

65 This was not an expected finding in comparison to previous results for nanocrystalline FCC Ni: the work hardening rate for Ni decreases with decreasing grain size owing to decreased dislocation activity [Wang et al. 1997], so Karimpoor et al. [2002 2003] attributed their finding to a possible different deformation mechanism, i.e. twinning. The high activation stress required for twinning followed by lower stress requirements to proceed is manifested in the low work hardening rates observed [Karimpoor et al. 2003]; however, the higher strain hardening rates observed in the current investigation than those observed by Karimpoor et al. [2002; 2003] contradict their findings. It should also be noted the strain hardening rates were not calculated for polycrystalline Co references, therefore these results are strictly confined to preliminary conclusions.

Strain-hardening exponents are calculated as per ASTM E646-16 from the logarithmic form of the true stress vs. true strain curves within the plastic region. The definition of the ‘plastic region’ is not clearly defined and the engineering strain range is only specified for low-carbon steels as a reference. A representation of how n-Values change with the selected true stress value corresponding to the onset of the plastic region is shown in Figure 37. Strain hardening rates or n-Values were calculated with ~250 MPa as the true stress value corresponding to the onset of the plastic region in this investigation.

n-Value vs. Selected True Stress (at onset of plastic region) 1800

e 1600

t s 1400

1200

o 1000

)

a 800

(

s 600

400

T 200

e 0

S 0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 n-Value (work-hardening rate)

Figure 37: n-Values from ML100-1 tensile test data when the true stress value at the ‘onset of plastic region’ varies from 0 to the true stress at fracture

66 Previously studied multilayered NiCo specimens showed near 10% ductility when sub-layers were composed of coarse grained and nanocrystalline grained electrodeposits and the coarse grained layers offered improved neck stability therefore high elongation values [Daly et al. 2015]. The evidence of this was observed in SEM imaging of the fracture surfaces, where periodic features were shown to align with sub-layer thicknesses and coarse grained layers had greater protrusions indicated larger plastic deformation, as shown in Figure 37. No such features were observed on either of the multilayered Co fracture surfaces, nor were sub-layers clearly distinguished in BF and DF XTEM imaging.

Figure 37: SEM imaging periodic features on fracture surface of coarse grained and nanocrystalline grained multilayered NiCo; scale bar represents 5m [Daly et al. 2015]

The strain rate sensitivity of electrodeposited nanocrystalline metals and has been investigated in literature. Decreases in strain rates have been found to increase ultimate tensile strengths, flow stresses and tensile elongation, and to decreased yield strengths [Karimpoor 2001; Karimpoor et al. 2002; Wang et al. 1997]. For example, Karimpoor et al. [2002] investigated nanocrystalline Co mechanical properties under tensile strain rates of 1 x 10-4 s-1, 5 x 10-4 s-1 and 2.5 x 10-3 s-1, and found that under a strain rate of 1 x 10-4 s-1 Co showed the highest UTS and a tensile elongation comparable to polycrystalline Co in

67 the same study. Yield strengths and flow stresses were not shown in this particular study [Karimpoor et al. 2002] to be affected by strain rates.

In addition to the differences in mechanical properties of monolithic and multilayered Co electrodeposits, there is a distinct shift in preferential orientation or texture within the bulk deposits from the (10 ̅11) to (0002) for MN to ML20/ML100 specimens, respectively. This shift is shown in the XRD patterns in Figure 22. There is also an observed increase in intensity of the (0002) peak and decrease in intensity of the (10 ̅11) peak for ML100 compared to ML20. This is not clearly visible in Figure 22. The XRD patterns for ML20 and ML100 are overlaid to more clearly show the intensities in Figure 38 below. As shown in this Figure, MN has a higher (10 ̅11) intensity than it does for the (0002) peak.

As discussed previously, additives to electrolytic solutions have been shown to influence the crystal growth mode and subsequent mechanical properties of deposits. The ductility of electrodeposited nanocrystalline Ni-W was found to be significantly affected by the texture and orientation of the microstructure, as shown in Figure 39 [Matsui et al. 2013]. The orientation is dependent on the crystal growth mode during deposition. For example, the preferred orientation has been found [Amblard et al. 1979] for Ni as the (111), (100) and (110) textures, depending on inhibited or uninhibited crystal growth modes.

The presence of nickel hydroxides and hydrogen acted as inhibitors on inhibitor crystal growth mode, or the (111) preferential texture [Amblard et al. 1979]. However, as previously discussed, the carbon and sulfur concentrations in the investigated monolithic and multilayered deposits were not shown to vary significantly or demonstrate a direct influence on crystallographic texture nor mechanical properties. The main observed difference in deposit characteristics is the change in texture upon the introduction of the nominal layers.

Figure 38: XRD peak intensities for ML100, ML20 and MN deposits

Figure 39: Tensile elongation vs index of (200) peak for Ni-Fe and Ni-W [Matsui et al. 2013]

Crystal orientation has been confirmed to influence tensile ductility in HCP materials [Matsui et al. 2013; Sakai et al. 2006; Wang et al. 2016]. Deformation twins have been shown to strengthen polycrystalline HCP materials by both the Hall-Petch effect by twin boundaries inhibiting dislocation movement and the Basinski mechanism [Basinski et

69 al. 1997] by transforming glissile dislocations into sessile dislocations, where the Burger’s vector is immobilized as it does not lie in the primary slip plane [Kalidindi et al. 2003]. However, a softening effect was also observed with high twin densities owing to a reorientation of the material to a favourable texture for slip to occur. A compressive strain of -0.2 caused a reorientation in 40% of HCP titanium to facilitate slip along the basal plane [Kalidindi et al. 2003].

The primary slip system in Co is (0002) <2 ̅1 ̅10>. Both ML20 and ML100 have a preferred basal plane orientation parallel to the surface of the deposit. In tensile testing, this basal plane is perpendicular to the tensile direction. According to the previous studies mentioned, preferential orientation to the slip system allows for a softening effect, which may or may not negate microstructural strengthening observed with the Hall-Petch effect and Basinski mechanism as introduced by twins or other grain refinements.

This may explain the relative strengthening effects observed with multilayered variations of electrodeposited Co. Theoretically, although grain size averages across the multilayered cross-sections remain constant, the layer interfaces offer an additional boundary impeding dislocation slip. However, this layer interface is not obvious in the TEM bright or dark field images and therefore cannot be confirmed as a significant strengthening mechanism.

In addition to the layer interfaces, the potential large density of dislocations or stacking faults may further strengthen the material by the Hall-Petch effect and Basinski mechanism. The preferred basal plane texture may also introduce a higher ductility than in multilayered Co than monolithic Co. Although a shift in preferred orientation from (10 ̅11) to (0002) may improve ductility in these materials, the difference in (0002) peak intensity between ML20 and ML100 material is not extreme and does not clearly explain the 40% increase in tensile elongation that was observed.

Digital image correlation with the help of a camera system was utilized to image the strain observed in tensile coupons during this investigation. Representative images for MN,

70 ML20 and ML100 materials are shown in Figure 40. All specimens had a mixed slanted, saw-toothed fracture surface that is common in electrodeposited nanocrystalline metals [Brooks et al. 2011; Daly et al. 2015; Karimpoor 2001] and fracture surfaces all had fine dimples, indicative of microvoid coalescence and plastic deformation, as shown in Figure 41.

(a) (b)

Figure 40: Fracture location from DIC imaging (a) and facture surfaces (b) imaged via secondary electrons (SE) at 15kV for MN, ML20 and ML100 Co. Scale bar in (b) represents 1mm

Figure 41: SE images of fracture surfaces for MN, ML20, and ML100 at 10kV showing dimpled fracture surfaces. Scale bar represents 10m

Microsamples were lifted out from near fracture surface regions for MN, ML20 and ML100 Co specimens. Bright field and dark field XTEM images of each are shown in Figures 42–44. Additional images are found in Appendix B. A high density of fringes

71 was observed in all specimens, which is a possible indication of a highly faulted HCP Co structure [Hibbard 2002; Karimpoor 2001]. Again, layers were not distinguishable in ML20 and ML100 images.

(a) (b)

Figure 43: BF (a) and (c) and DF (b) and (d) XTEM images of ML20 Co

(a) (b)

74 5 CONCLUSIONS The structure and deformation behavior of nanocrystalline cobalt formed by electrodeposition was investigated. Monolithic and multilayered Co structures with nominal (1) 20nm and (2) 100nm sub-layer thicknesses of alternating electrodeposition conditions were studied. 1. It was found that while all three specimen types were of the hexagonal crystal structure, a change in preferred orientation occurred with the introduction of nominal multilayered structures. Monolithic cobalt had a preferred (10 ̅11) texture and multilayered cobalt had a preferred (0002) or basal plane texture. The (0002) peak had higher intensity with increasing nominal sub-layer thickness. 2. Bulk chemical analysis showed that both multilayered Co deposits exhibited higher sulfur content than the monolithic Co, but all concentrations were relatively low in comparison to previous studies of similar materials [Hibbard et al. 2006] and were not shown to directly influence crystallographic texture. 3. Tensile tests were performed at a strain rate of 5 X 10-4 s-1 and microhardness tests were performed under a 100g load. The average hardness, yield strength, ultimate tensile strength and fracture strength are increased when the electrodeposited cobalt follows a nominal multilayered structure. 4. Multilayered Co of 100nm nominal sub-layer thickness showed tensile elongation of ~8%, which was near a 75% and >100% increase from multilayered Co with 20nm nominal sub-layer thickness and monolithic Co, respectively. 5. As deposit microstructure for monolithic and multilayered cobalt did not show significant differences, connection between tensile properties and crystallographic orientation of the material is proposed: higher tensile elongation values were seen for deposits with preferred orientation to the slip system, (0002) <2 ̅1 ̅10>. A similar relationship that has been previously noted for nanocrystalline Ni-W electrodeposits [Matsui et al. 2013]. Further work is required to determine if this effect carries across multilayered cobalt with varying sub-layer thicknesses and should be compared to polycrystalline Co counterparts.

75 6 RECOMMENDATIONS A complete understanding of the mechanical properties and deformation mechanisms of multilayered nanocrystalline cobalt electrodeposits would benefit from work on the following matters:

1. Perform tensile tests on Co deposits with nominal sub-layer thicknesses above 100nm (i.e. 200nm, 500nm, 1m), Co deposits with nominal sub-layer thicknesses between 100nm and 20nm and polycrystalline Co deposits for reference 2. Perform tensile tests at varying strain rates (eg. 2.5 x 10-3 s-1 , 1 x 10-4 s-1) 3. In-situ TEM tensile or compression tests 4. Study the effect of temperature on the mechanical properties of multilayered cobalt in comparison to monolithic cobalt 5. Analysis of the preferred crystallographic orientation on deposits with nominal sub- layer thicknesses above 100nm (i.e. 200nm, 500nm, 1m) and deposits with nominal sub-layer thicknesses between 100nm and 20nm

76 APPENDICES Appendix A: FIB microsampling procedure

Figure A1: trough milling for micro sampling (X3.5K)

Figure A2: Top-down view of micro sampling specimen (X2.2K)

Figure A3: Lift-out of specimen (X2.5K)

Figure A4: Specimen pre-thinning (X700)

Figure A5: Final view of specimen thinning (X7.0K)

Figure A6: Location of microsampling from tensile tested coupons. The fracture surface is indicated at the white arrow. Vertical and diagonal lines observed in this image are a result of sample polishing (X400)

Appendix B: Additional TEM Images of as-deposited and near-fracture surface specimens

Figure B1: MN BF image

Figure B2: MN BF image

Figure B3: MN DF image

Figure B4: MN BF image

Figure B5: ML20 BF image

Figure B6: ML20 DF image

Figure B7: ML20 BF image

Figure B8: ML20 DF image

Figure B9: ML20 BF image

Figure B10: ML20 DF image

Figure B11: ML100 BF image

Figure B12: ML100 BF image

Figure B13: ML100 BF image

Figure B14: Near-fracture surface MN BF image

Figure B15: Near-fracture surface MN BF image

Figure B16: Near-fracture surface MN DF image

Figure B17: Near-fracture surface MN BF image

Figure B18: Near-fracture surface MN DF image

Figure B19: Near-fracture surface ML20 BF image

Figure B20: Near-fracture surface ML20 BF image

Figure B21: Near-fracture surface ML20 DF image

Figure B22: Near-fracture surface ML20 BF image

Figure B23: Near-fracture surface ML20 DF image

Figure B24: Near-fracture surface ML20 BF image

Figure B25: Near-fracture surface ML20 DF image

Figure B26: Near-fracture surface ML100 BF image

Figure B27: Near-fracture surface ML100 BF image

Figure B28: Near-fracture surface ML100 BF image

Figure B29: Near-fracture surface ML100 DF image

Figure B30: Near-fracture surface ML100 BF image

Figure B31: Near-fracture surface ML100 DF image

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