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MECHANICAL PROPERTIES OF NANOCRYSTALLINE NOMINALLY
MULTILAYERED HEXAGONAL COBALT ELECTRODEPOSITS
BY
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
ii
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.
iii
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.
iv
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
v
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
vi
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
vii
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]
viii
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 1m
Figure 27: 500nm layered Co deposit (X10.0K); scale bar represents 5m
ix
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 5m [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 10m
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
x
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
1
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.
2
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.
3
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 10m[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
4
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.
5
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.
6
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.
7
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]
8
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
9
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.
10
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]
11
Table 3: Mechanical and structural effects of current type on electrodeposited metals
Material Current
Type
CD
(A/dm2)
Time
On/Off
Surface
roughness
(um)
Hardness
(HVN)
Grain
size
(nm)
Reference
Ni DC 100
ma/cm2
--- Ra=0.37
rz=2.1
350 45 [Sanacian et al.
2014]
Ni PC 100
ma/cm2
2ms-18ms Ra=0.25
rz=1.6
556 31 [Sanacian et al.
2014]
NiWTiO2 DC 1.5 - Somewhat
uniform
467-686 - [Kumar et al.
2013]
NiWTiO2 PC 1.5 40ms:30ms Smooth
and
smaller
spherical
grains
(more
uniform
than
NiWTiO2
plated
under DC)
Higher
than DC
(by about
100)
(Finer
than
NiWTiO2
plated
under
DC)
[Kumar et al.
2013]
NiCo PC 10 50% duty
cycle
- 450 - [Zamani et al
2016]
NiCo DC 5 - Irregular
polyhedral,
~0.4 um
<300 25.99 [Karslioglu et
al. 2015]
NiCo PC 5 10 10 ms Reduced
crystallite
size, ~0.2
um
~300 24.5 [Karslioglu et
al. 2015]
NiCo PRC 5 10/10/10ms Spherical
cluster
equiaxed,
~0.2 um
>400 23.93 [Karslioglu et
al. 2015]
CoP DC 10 - Smooth
fine
globular
- 30 [Kosta et al.
2011]
CoP PC 175 2.5ms/45ms Smooth
fine
globular
- 10 [Kosta et al.
2011]
12
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].
13
(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 ��) [Karimpoor et al. 2003]
14
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.
15
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].
16
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]
17
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
18
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]
19
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]
20
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,
expressed as: 𝜎𝑦 = 𝜎0 + 𝑘𝑑−1/2 or 𝐻 = 𝐻0 + 𝑘𝐻𝑑−1/2, 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].
21
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
22
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].
23
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
24
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]
25
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]
26
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
27
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].
28
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] .
29
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 = 450C + 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
30
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].
31
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}, {1122} and {1121} families of planes, with
32
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
33
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.
34
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.
35
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
36
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 800C - 1000C has been
observed at an elongation of 15-30% [ASM International 2007].
37
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.
38
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
39
(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.
40
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
41
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)
42
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.
44
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)
45
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-
layer thickness Cobalt
Co: 99.3; Fe: 0.254; Mn: 0.107; S: 0.0815;
Si: 0.207
ML100 100nm nominal sub-
layer thickness Cobalt
Co: 99.2; Fe: 0.279; Mn: 0.184; S: 0.0820;
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.
46
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 (1120) 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].
47
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), (1120), (10 13) and (2020) rings are shown in increasing diameter from the
transmitted beam, respectively.
48
(a) (b)
(c) (d)
Figure 23: BF (a) and (c) and DF (b) and (d) TEM images and (e) SADP inset (L=300) of MN Co
49
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.
50
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, Moire 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
51
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) [Peter et al.
52
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].
Figure 26: 100nm layered Co deposit (X50.0K); scale bar represents 1m; deposit thickness is in the
vertical direction, t0
Figure 27: 500nm layered Co deposit (X10.0K); scale bar represents 5m; deposit thickness is in the
vertical direction, t0
𝒕𝟎
𝒕𝟎
𝑴𝑳𝟏𝟎𝟎
53
(a)
(b)
Figure 28: (a) ML20 and (b) ML100 BF TEM images; deposit thickness is in the vertical direction, t0
𝒕𝟎
𝒕𝟎
𝑴𝑳𝟐𝟎
𝑴𝑳𝟏𝟎𝟎
54
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
55
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 ��1)
rings are circled in (b) and (c)
56
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
(ppm)
Sulfur
(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-
57
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]
58
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
59
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.
60
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].
61
Brooks et al. [2008] found that nanocrystalline materials follow the same hardness-
strength relationship as their polycrystalline counterparts: HV = 3UTS 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.
62
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
63
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
64
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
65
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
Considere’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].
66
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.
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
0
200
400
600
800
1000
1200
1400
1600
1800
0 0.1 0.2 0.3 0.4 0.5 0.6 0.7SelectedTrueStress(MPa)atonsetofplasticregion
n-Value(work-hardeningrate)
n-Valuevs.SelectedTrueStress(atonsetofplasticregion)
67
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 5m [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
68
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.
69
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
70
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,
71
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 10m
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
72
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)
(c) (d)
Figure 42: BF (a) and (c) and DF (b) and (d) XTEM images of MN Co
73
(a) (b)
(c) (d)
Figure 43: BF (a) and (c) and DF (b) and (d) XTEM images of ML20 Co
74
(a) (b)
(c) (d)
Figure 44: BF (a) and (c) and DF (b) and (d) XTEM images of ML100 Co
75
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.
76
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, 1m), 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, 1m) and deposits with
nominal sub-layer thicknesses between 100nm and 20nm
77
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)
78
Figure A3: Lift-out of specimen (X2.5K)
Figure A4: Specimen pre-thinning (X700)
79
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)
80
Appendix B: Additional TEM Images of as-deposited and near-fracture surface
specimens
Figure B1: MN BF image
Figure B2: MN BF image
81
Figure B3: MN DF image
Figure B4: MN BF image
82
Figure B5: ML20 BF image
Figure B6: ML20 DF image
83
Figure B7: ML20 BF image
Figure B8: ML20 DF image
84
Figure B9: ML20 BF image
Figure B10: ML20 DF image
85
Figure B11: ML100 BF image
Figure B12: ML100 BF image
86
Figure B13: ML100 BF image
87
Figure B14: Near-fracture surface MN BF image
Figure B15: Near-fracture surface MN BF image
88
Figure B16: Near-fracture surface MN DF image
Figure B17: Near-fracture surface MN BF image
89
Figure B18: Near-fracture surface MN DF image
Figure B19: Near-fracture surface ML20 BF image
90
Figure B20: Near-fracture surface ML20 BF image
Figure B21: Near-fracture surface ML20 DF image
91
Figure B22: Near-fracture surface ML20 BF image
Figure B23: Near-fracture surface ML20 DF image
92
Figure B24: Near-fracture surface ML20 BF image
Figure B25: Near-fracture surface ML20 DF image
93
Figure B26: Near-fracture surface ML100 BF image
Figure B27: Near-fracture surface ML100 BF image
94
Figure B28: Near-fracture surface ML100 BF image
Figure B29: Near-fracture surface ML100 DF image
95
Figure B30: Near-fracture surface ML100 BF image
Figure B31: Near-fracture surface ML100 DF image
96
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