Mat., 2,377, (1993). S. J. Savage, and F.H Froes, J. Metals, April, 29 (1 984). Y.
Suwa, R. Roy, and S. Komarneuis, Mat. Sci. Eng., 83, 15 1, (1 986). D.W. Hohan ...
Magnetic Properties of Nanocrystalline Transition Metals
Martin J. Aus
A thesis submitted tu the Department of Materials and Metallurgical Engineering in
confomity with the requirementsfor the degree of Doctor of Philosophy
Queen's University
Kingston, Ontano, Canada Copyright Q Martin J. Aus, 1999.
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In Loving Memory of My Mother and Fatber...
Abstract In the past decade, considerable attention has been devoted to the nanoprocessing of magnetic materials to enhance specifk magnetic properties. For nanocrystalline
materials in which the grain size approaches the dimensions of the domain wall thickness
of conventional rnaterials, considerable changes in magnetic behaviour are expected.
in the present work, various electrodeposited ferromagnetic nanocrystalline pure metals and alloys were characterized by uimg a vibrating sarnple magnetometer. The systems investigated include pure Ni and Co as well as alloys of Ni-P,Ni-Fe and Co-Fe. These studies explored the eEect of grain sue on coercivity, indicating that the crystallographictexture is more significant than grain size. In addition, these studies reported, for the first t h e , that saturation magnetization of pore-fkee electroplated bulk nanocrystalline transition metals and their alloys is relatively linle affected by grain size.
In contrast, previously reported results for ultra-fine particles and nanomaterials produced fiom compacted powders showed a strong decrease in saturation magnetization with decreasing grain size. The dierence in results for pore-fiee electrodeposits and ultra-
fine particles/compacted powders has been attnbuted to antiferromagnetic surface oxide layers, which is a direct result of large intemal porosity in the latter group of materials. Further magnetic studies were completed on nanocrystalline electrodeposits produced by magnetoelectrohydrol).sis. The effects of applied magnetic field strength and substrate orientation on saturation magnetition and coercivity of NoFe and Co were explored. The results have s h o w that both nanoprocessing and electroplating in a
magnetic field can improve soft magnetic properties by lowering the coercivity. Thermomagnetic studies examineci saturation magnetkation as a function of The Curie temperature, Curie temperature and coercivity changes d u ~ annealing. g temperatures of electrodeposited nanocrystals does not differ fiom that observed in polyc~stdlinematerials. The coercivity of nanoc~stallineNi, Ni- 1&/OP and Ni2ûwt??Fedecreased upon annealhg at temperatures below the grain growth temperature
as a result of internai stress relieving. An overall picture of the rnagnetic properties of nanocrystalline metals is now emerging and theory is under development to explain recently observed domain structures and other magnetic phenornena observed in nanocrystaliine materials.
Acknowledgements
First and foremost, I'd like to thank U.Erb for giving me the oppomullty and encouragement to complete a doctoral thesis in the field of nanocrystalline materials, in
which he is both a pioneer and world-class scientist. Professor Erb has taught me a great deal about academia, attention to detail, and work ethics. For this I am grateful. Secondly, I would like to thank two professors from Queen's University, narnely;
D.L. Athedon, who had faith in my abilities and W.M.Conke who did not. The polar opposite attitudes that these gentlemen displayed not oniy motivated me, but ensured my success. During the body of this work 1 was fortunate enough to work with the briliiant
minds of K.T. Aust, G. Palumbo, B. Szpunar and J. Szpunar. I would also like to thank those researchers who came before me, namely; AM.El-Sherik, A. Alfantazi, S. Wang, and also T. and M-L.Turi for remernbering my mother. I would also iike to wish J. McCrea, R Zugic and G. Hbbard the best of luck in their future work in this field. There have been many research associates who have added not only their ideas but their fiiendship during this project and they are as foilows: C. Cheung, D. Clark, E. Griswold, R Langer, G. Panagiotopoulos, and D. Ueno. Last but not least, 1 would like to thank my farnily for their support during this lengthy research period. Leila, Denise, Stephanie and Delons, Thank You.
TABLE OF CONTENTS
A b s t c t .m.....m...................m.........................................~................... i Acknowledgements ..........................................................................
..
II
...
Table o f Contents ............................m.............................................. III
Lit of Figures ............................m..........................m....................... ix
List of Tables ...............................................................................
..................................................................
Glossary of Symbois List of Units
xiii xiv
...............................m..........................................m... xvi
Selected Values of Physical Constants
CHAPTER ONE
CEIAPTER TWO
.....................m.......................... xvii O.
........................m.................m.....m..................m...... 1
.........................................................................
8
THEORY OF MAGNETISM ...........................................................8
............................................................. 8 2.2 Atomic Magnetism ....................................................... 1 1
2.1 Historical Perspective
iii
2.3 Classifications of Magnetic Materials .......................................
13
2.3.1 Diamagnetism ............................................................ 13
2.3.2 Paramagentism ............................................................ 15 2.3.3 Ferromagnetism
........................................................... 18
....................................................... 25 2.3.5 Ferrimagnetism ............................................................ 27 2.3.6 Superparamagnetismand Helimagnetism ............................. 28 2.3-4 Antiferromagnetism
2.4 Energy Contributions to Magnetic Domain Formation ....................... 29 2.4.1 Exchange Energy ......................................................... 30
2.4.2 Magnetostatic Energy .................................................... 31 2.4.3 Anisotropy Energy ....................................................... 32 2.4.4 Magnetoelastic Energy ................................................... 36
2.4.5 Domain Wall Energy ................................................... 37
2.5 Magnetic Properties Derived fiom a Hysteresis Loop ......................... 38 2.6 Hard and Soft Magnetic Materials ........................................... 43
2.7 References .......................................................................... 45
CHAPTER 3
...............................................................................
OVERVIEW OF NANOCRYSTALLINE MATERL4LS ...............
46
46
3.1 h i n Boundary Engineering ................................................... 46 3.1.1 Introduction to Grain Boundaries ...................................... 46 3.1.2 Special Grain Boundaries 3.1.3
........................................... 47
Property Differences Between Law Z and General Boundaries ... 48
3.1.4 Grain Boundary Design ................................................. 48
3-2 Nanocrystalline Materials ........................................................ 49 3.2.1 Introduction to Nanocrystalline Materials
............................. 49
3.3 Structure of Nanocrystalline Materials ......................................... 50 3-4 Synthesis of Nanomaterials ...................................................... 54
3.4.1 Introduction to Nanoprocessing ........................................ 54 3.4.2 Nanoprocessing Techniques Requiring Consolidation ............... 55
.........................................55 3.4.2.2MechanicalAlloying ............................................ 57
3 A2.1 Inert Gas Condensation
3 .4.3 Nanoprocesshg Techniques Yielding Nanomaterials Directly ..... 58
3.4.3.1 horphous Precursors .......................................... 58
3.4.3.2Electrodeposition ................................................ 59
3.4.4 Cornparison of Nanoprocessing Techniques ......................... 61 3.5 Prope~iesof Nanocrystalline M ~ t z i a l s ....................................... 62 3 .5 .1 Physical Properties ...................................................... 62
........................................................... 62 3.5.1.2Hardness ......................................................... 63 3 S.2 Chernid Properties ..................................................... 66 3.5.2.1Corrosion ........................................................ 66 3.5.1.1 Density
3.5.3 Electrical Properties ..................................................... 68 3.5 -4 Magnetic Properties ..................................................... 72 3.5.4.1 CoerciMty ......................................................
72
3.5.4.2 Saturation Magnetization ....................................... 74
3.5.4.3Magnetic Microstructure ....................................... 75 3.5.4.4 Magnetic Transitions ............................................. 76
3.6 References .......................................................................... 77
CHAPTER FOUR
.......................................................................
84
DEVELOPMENT OF APPARATUS ........................................
84
4.1 Measurement of Magnetic Properties using a Hybrid VSM ...............
84
4.2 Development of Hybnd Vibrating Sample Magnetometer .................. 86
.................................................... 89 4.4 Advantages and Disadvantages of Technique................................ 90
4.3 Calibration of Hybrid VSM
4.5 References
......................................................................... 92
MAGNETIC PROPERTIES OF PURE NANOCRYSTALLINE TRANSITION METALS .......................... 93 5.1 Magnetic Properties of Pure Nanocrystalline Ni and Co ..................... 93 5.1.1 Introduction ...................... ;. ...,,. ................................ 93
.............................................................. 5.1.3 Results and Discussion .................................................. 5.1.4 Conciusions .............................................................. 5.1.5 References ............................................................ 5.1.2 ExperYnental
93
99 106
108
5.2 Magnetic Propenies of Nanocrystalline Nickel Phosphorus .................. 110 5.2.1 Introduction ...........................................................
110
5.2-2 Experimental ............................................................ 1 1 1 5.2.3 ResuIts and Discussion ................................................. 113
5 .2.4 Conclusions ............................................................. 119 5.2.5 References ..............................................................
5.3 Magnetic Properties of Nanocrystalline Ni-Fe
120
........................ 121
5.3.1 Introduction ............................................................ 121
. . . 123
5.3.2 Experimental .....................................
5.3.3 Results and Discussion ................................................. 126
5.3-4 Conciusions ....................................................... 5.3.5 References ...................................... 5.4
128
.
129
,
......................... 130 5.4.1 Introduction ........................................................... 130
Magnetic Propenies of Electrodeposited Co-Fe
5.4.2 Experirnental .............................................................
134
5.4.3 Results and Discussion .................................................
136
.............................................................. 5.4.5 References ...............................................................
140
5.4.4 Conclusions
CHAPTERSIX
..................................................................
141 142
MAGNETOELECTROHYDROLYSISOF PURE NANOCRYSTALLINE Ni AlSiTl Co ..............................
142
6.1
Introdii~~ion..................................................................
142
6.2
Expenmental ...................................................................
144
6.3
Results and Discussions .......................................................
146
6.4
Conclusions
6.5
References
CEiAPTER SEVEN
.................................................................... ....................................................................
150 152
...................................................................
Introduction ....................................................................
Temperature Dependence of Ms and Tcin Nanocrystailine Ni and Ni-Fe 7.2.1 Experimental ............................................................ 7.2.2 Results and Discussion 7.2.3 Conclusions
...........................................
..............................................................
Effects of Annealing on Hc of Nanocrystailine Ni, Ni-P and Ni-2O%Fe 7.3.1 Experimental ...........................................................
7.3.2 Results and Discussion ................................................
7.3-3 Conclusions
References
.............................................................
.......................................................................
vii
ANALYSIS. TRENDS AM> APPLICATIONS FOR
NANOCRYSTALLINE MAGNETIC MATERIALS .......................
174
8.1
Magnetisrn of Nanocrystalline Femomagnets ................................... 174
8.2
Potential Applications for Nanocrystaiiine Soft Magnets ................... 181
8.3
References .......................................................................... 188
CHAFTER NINE
..........................................................................
190
CONCLUSIONS ................................................................... 190
List of Figures Figure 1.1:
Grain boundary design concepts......................................... 2
Figure 2.1:
Vanation of the spontaneous magnetization (MJof nickel within a domain as a finction of temperature according to the Weiss mean field model, after Weiss and Forrer .............................. 20
Figure 2.2:
Temperature dependence of spontaneous magnetization within a domain for iron, cobalt, and nickel compared with calculations baseci on the ferromagnetic Brillouin firnction.......................... 23
Figure 2.3:
The Slater-Paulingcurve, which gives the net magnetic moment per atom as a function of the number of 3d electrons per atom...... 24
Figure 2.4:
Arrangement of magnetic moments in an antiferromagnet at T=OK, a) in zero applied field and b) in a field, H, applied perpendicular to the moments.. ............................................................ 26
Figure 2.5:
Magnetic domain structure in the surface of iron.. ................... 29
Figure 2.6:
Magnetization curves for a single crystal of cubic iron and cubic nickel with applied field parallel to various crystailographic directions.................................................................... . 3 3
Figure 2.7:
Easy and hard rnagnetization cuntes for cobalt at room temperature.................................................................. 35
Figure 2.8:
Schematic diagram of M vs. H or hysteresis loop showing magnetic properties and domain structures during different stages . . of rnagnetization.. .......................................................... 39
Figure 2.9:
Hysteresis loop behaviour for soft iron and hardened steel.. ......... 42
Figure 2.10: Relative permeabilities and coercivities of ferromagnetic materials. .................................................................... .44 Figure 3.1:
Tetrakaidecahedra representing grain shape for equiaxed rnaterials. 5 1
Figure 3.2:
Schematic diagram of a nanocrystalline materiai....................... 52
Figure 3.3:
Intercrystailine volume fiaction as a function of grain size............ 53
Figure 3.4:
Schematic representation of gas-condensation charnber for the synthesis of nanocrystalline materials............................... 56
Figure 3.5:
Schematic diagrarn of electrodeposition apparatus.. ............. ...... 5 9
Figure 3.6:
Schematic representation of square-wave modulated curent.. ....... 60
Figure 3.7:
Hall-Petch plot for nanoc~ystahegad-condensed Pd, Cu and electrodeposited N-P. ......................................... ......... ....65
Figure 3.8:
Resistivities . . of gas condensed Pd as a finction of temperature and gram size.. . .............................. ........ ....................... 69
Figure 3.9:
Excess resistivity as a iùnction of intercrystalline volume fraction in electrodeposited M... ... . .. . .. ... ... ......... . ... . . . .. . . . .. . . . . .. . .. . . .. 70
Figure 3.10: Temperature coefficient of resistivity vs. grain sue in gas condensed Pd. ............................................................ ... 71
Figure 3.11: Coercivity vs. grain skze for various soft magnetic doys.. . .... ...... 73 Figure 4.1:
Schematic diagram of the vibrating sample magnetometer (VSM). 85
Figure 4.2:
Schematic diagrarn of the hybrid vibrating sarnple magnetometer.
Figure 4.3 :
Sample output fiom the hybnd VSM for an experiment using polycrystalline nickel (99.99% purity). ..................... . .. ..... ... 88
Figure 4.4:
Schematic diagrarn of sample configuration between pole pieces of electromagnet and induced moment piclaip coils.. ....... .............. 89
Figure 5.1 :
TEM darffield micrograph of nanocrystalline nickel h a h g anaveragegralnsizeofllnm..........................................
86
95
Figure 5.2:
X-ray diiacîion scans showing weak (200) texture, a) as plated material (20nm), b) amealed material (40nrn). .......... 96
Figure 5.3:
Brightfield transmission electron rnimograph of nanocrystalline Co electrodeposit with 1Snrn grain size.. .............. .. .. . .. .. ..... . ...96
Figure 5.4:
Scanning electron micrograph of electrodeposited polycrystalline Co with an average grain size of 5p.m.. ............ ...... .... ... ...... 97
Figure 5.5:
X-ray difiaction scans of electrodeposited Co, comparing nanocrystalline (1 5nm) and polycrystaliine (Spm) structures.. . .... 98
Figure 5.6:
Saturation magnetization of Ni as a fiindon of grain size........... 100
Figure 5.7:
Coercivity of Ni as a function of grain size.. . .. . ... ...... . .. ... ... .... 104
Figure 5.8:
Hysteresis ioops for a) nanocrystaihe Co and b) polycrystalline Co....................................................... 105
Figure 5.9:
a) Darkfield electron micrograph of electrodeposited NI-Pwith average grain size of IOnm. b) Corresponding electron dEaction Pattern........................................................................ I l 2
Figure 5.10: a) Brightfield electron micrograph of electrodeposited amorphous Ni-P, Corresponding electron deaction pattern.. .................... 112 Figure 5.11 : a) Saturation magnetization of nanocrystalline and amorphous N-P electrodeposits (Region 1: amorphous, Region II: nanocrystalline). b) Saturation magnetization of pure Nielectrodeposits as a function of@ size (section 5.1) .................................................. 114 Figure 5.12:
Effect of phosphorus content on saturation magnetization. (Region 1: amorphous, Region II: nanocrystalline)..................
1 15
Figure 5.13 : a) Coercivity of nanocrystalline and amorphousNi-P electrodeposits (Region 1: amorphous, Region II: nanocrystalline) b) Coercivity of pure Ni electrodeposits as a function of grain size. 1 17 Figure 5.14:
TEM micrographs for an electrodeposited nanocrystalline Ni-20wt%Fe alloy: a) brightfield, b) darkfield........................ 124
Figure 5.15 : Effect of iron content on the grain sùe of electrodeposited nanocrystalline Ni-Fe alloys. ............................................. 125 Figure 5.16:
Saturation magnetization of nanocrystalline Ni-Fe electrodeposits. 126
Figure 5.17:
Coercivity of nanocrystalhe Ni-Fe electrodeposits.. ................. 127
Figure 5.18:
Variation of saturation magnetiiion with composition in iron cobalt alloys after Weiss and Forrer.. ............................... 13 1
Figure 5.19 : Phase diagram of Co-Fe ailoy............................................. 132 Figure 5.20:
Phase transformation in iron-cobalt (50%Co) showing the high temperature f C.C.phase is non-magnetic. ............................. 13 3
Figure 5.2 1: Saturation mapetkition of Co-Fe alloys fiom Owt%-22wt%Fe.... 138 Figure 5.22:
Coercivity as a function of composition in Co-Fe electrodeposits.
139
Figure 6.1:
Apparatus developed for magnetoelectrohydrolysis of Ni-2O%Fe.
146
xi
Figure 6.2:
Effects of in-field plating on saturation magnetization................ 147
Figure 6.3:
Effects of in-field plating on wercivity.. ............................ .... 148
Figure 7.1:
vs. BIof typical soft magnetic materials.. ........................ 154
Figure 7.2:
Saturation magnetization behaviour as a fùnction of temperature.
161
Figure 7.3:
Determination of Cune temperature through graphical analysis.
163
Figure 7.4:
Coercivity as a fùnction of annealhg time for nanocrystalline Ni
167
Figure 7.5:
Coercivity of Ni and NiP electrodeposits after annealing............ 168
Figure 7.6:
Coercivity of annealed nanocrystalline Ni-20wt%Fe. .............. ... 169
Figure 8.1:
Magnetization in Co-W alloys as a function of tungsten concentration. Fiiied squares: theoretical calculations................ 176
Figure 8.2:
Magnetization in N-P doys as a function of phosphorus concentration. Fiiied squared: theoretical calculations................ 176
Figure 8.3:
Coercivity (Hc) vs. grain size @) for various soft magnetic alloys. Nanocrystaiiine materials include: FeNbSiB, FeCuVSiB, FeZrB and FeCoZr............................,.................................. ... 178
Figure 8.4:
Striped domain structure in nanocrystaliine Ni (1 hr @ 1OO°C). ..
Figure 8.5:
Magnetic force microscopy showing regular saip domains on nanocrystallineNi.. ........................................... ...... ........ 180
Figure 8.6:
Resistivity as a fûnction of temperature and grain size................ 183
Figure 8.7:
Relationship between Bsand 1 lcHz for NANOPERM~,the nanocrystalline Fe-SiB-Nb-Cu alloys and conventional soft magnets......... . ................... ....................... .. ... . .. ... ...... 186
180
List of Tables Table 2.1.
Curie points for various ferromagnetic elements..................... 20
Table 2.2.
Anisotropy constants for, nickel and cobalt ...........................
Table 5.1.
Saturation magnetization and coercivity of nickel.................... 99
Table 5.3:
Gralli sue. atornic percent phosphorus. volume ikaction and magnetic properties for NiP ............................................. 113
Table 5.4.
Materials requirements for magnetic recording heads................. 122
Table 5.5.
Composites vs. structure for Co-Fe electrodeposits.................... 137
Table 6.1:
Magnetoelectrohydrolysisof Cobalt...................................... 149
Table 7.1:
Mapetostridion and anisotropy minima in NI-Fe..................... 155
Table 7.2.
Grain growth temperatures and peak activation energies.............. 157
Table 7.3:
Cune temperatures of nanocrystalline Ni and NÏ-Fe................... 164
Table 8.1:
Cornparison of saturation induction and electrical resistivity......... 183
34
Glossary of Symbols Mean field constant Direction cosines of magnetic vector with respect to the applied field
Magnetic induction Remanent induction Saturation magnetic induction Direction cosines of direction of measurement with respect to the applied field Curie constant Velocity of light Susceptibility Initial susceptibility Maximum susceptibility Energy Electronic charge Exchange energy Anisotropy energy Magnetostatic energy Magnetoelastic energy Domain wali energy Spectroswpic splitting factor Magnetic field strength Planck's constant Coercivity Demagnetizing field Total atomic angular momentum quantum number Exchange constant
Anisotropy constant Boltwian's constant First anisotropy constant for cubic system Second anisotropy constant for cubic system Orbital angular momentum Length Domain wall thickness Magnetostriction
xiv
Magnetization Magnetic moment Electronic mass Remanent magnetization Saturation magnetizaiton Permeability Bohr magneton
Initial pemeability Maximum permeability Pemeability of fke space Nurnber of atorns per unit volume Demagnetizhg factor
Radius
Atomc spin angular momentum Reciprocal ffee volume Temperature Curie temperature Criticai temperature Neel temperature Angle Hysteresis loss
Number of electrons per atom
List of Units Symbol
Quantity
emu
Magnetic induction Remanent induction Saturation magnetic induction Susceptibœity Initial susceptibility Maximum susceptibiüt.
unitless unitless
unitless
emu oe-' emu cc")0e-' emu cc-3 ~ e - '
Energy Magnetic field strength Coercivity Demagnetizing field
Oe (Oerstead)
Oe Oe
Anisotropy constant Length
Magnetization Magnetic moment Rernanent magnetization Saturation rnagnetkaiton Permeability Initial permeability Maximum permeability Temperature Curie ternperahire Neel temperature
emu cë3 erg/G ernu cè3 emu cc-3
Selected Values of Physical Constants Velocity of iight in empty space
2.998 x 108 d s
Electronic charge
-1.602 x 1 0 - l ~C
Planck's constant
6.626 x
Boltzman's constant
1.381 x lua JK
Electronic rnass
9.109 x lu3'kg
Js
Bohr magneton
Permeability of f?ee space
1.257 x
lo4 W m
Chapter One INTRODUCTION
In recent years, the successfûl application of grain boundary engineering concepts to process rnaterials with improved physical and chemical properties has been clearly
dernonstrated [1,2]. In generai, grain boundary engineering c m be achieved by two approaches as show in Figure 1.1. The first approach makes use of structural differences obsented for dserent types of grain boundaries and tnple junctions, and the associateci variations in grain boundary and triple junction properties. Aust et al. [3] recently reviewed the structural models and
geometrical critena for special grain boundaries and triple junctions and assessed their importance on the basis of experimental studies of energetic, kinetic, physical, chemical and mechanical phenornena. Methods of controlling the grain boundary character
distribution in an effort to improve the overall performance of conventional polyciystalline materials have also been presented in the literature [3-61. The second approach to grain boundary engineering involves the increase in overall quantity of grain boundaries and triple junctions, which yield nanocrystaüine materials. It was first predicted by Gleiter [7]and has, since then, been experimentaily confirmeci that many properties of nanocrystalline materiais differ significantly fiom
those of conventionai materiais. This is a direct result of the high interfacial content inherent to nanocrystailine materials.
[Grain Boundary Engineering
Conveotiond hlaterids Gnin Boundary Churicter Distribution 'Special C Boundhes'
1
Materisils
Intercrystdl ine Volume Fraction
Magnetic Electrical Chemical Mechanical
Figure 1.1 :
Grain boundary design concepts
In a recent review, Gleiter [8] reported on several property enhancements through nanoprocessing, which has led to considerable international efforts to produce nanocrystals in large quantities. Nanoc~ystdlinematerials are produced by several methods including spark erosion, gas-condensation, and ball-milling[9]. However, these production methods yield nanocrystalline powders which require secondary consolidation processing such as cold isostatic pressing. Consequently the final product may contain considerable porosity. It
has been shown that even a small percentage of porosity cm have a large effect on the properties of nanocrystalline materials such as Young's modulus[l0,11].
In contrast, the p s t 15 years of research efforts by Erb et al. at Queen's
University, Kingston, Canada, have concentrat ed on electrochernicalproduction methods including conventional direct current plating, electroless plathg, cm-deposition, pulse plating and continuous drurn plating techniques to produce pure metals, alloys and composites in nanocrystalline fonn [12,13]. Nanocystalline electrodeposits are vimially pore-fiee and as such, represent the bulk materials properties of fully dense nanocrystalline materials. In the past decade, considerable attention has been devoted to the nanoprocessing
of magnetic materials to enhance specific magnetic properties[7,12,14-321. The nanocrystalline materials produced for these purposes range fiom bulk samples to granular films and compositionally modulated thin films. Many techniques have been used to synthesize these materials including gas condensatio~i17,149181, bal1 rnilling [19,20], chexnical processing [21-22], metal-organo complex routes as well as solution reduction and CO-precipitation[23], sol-gel methods [24], sputtering [24-261, rapid solidification [27], devitrification of amorphous alloys [28-3 11, and electrodeposition [13,24]. Potential applications of nanoprocessed materials with enhanced soft magnetic
properties have been disaissed in the literature [30-321. In addition, the applications of nanocomposites for magnetic refngeration [23,24], layered nanostnictured materials showing gant magnetoresistive effects [25,26]and hard magnets 1271 have been discussed.
The magnetic properties of polycrystalline ferromagnetic materials are strongly dependent on parameters such as crystallographic texture, intemal stress, grain shape
anisotropy, grain size distribution, etc. [33,341. For nanocrystalline materials in which
the grain size approaches the dimensions of the domain wali thickness of conventionai matends considerable changes in the magnetic behaviour are expected. With this in mind, the main objective of this thesis is to magneticaly characterize various ferromagnetic nanocrystalline pure met& and d o y s which were produced by electrodeposition over the past several years. These systems were synthesized by Eh et al. at Queen's University, Kingston, Canada and include pure nanocrystalline Nt and Co as well as alloys of Ni-P, Ni-Fe and Co-Fe. The rnagnetic measurements consist of hysteresis loops generated by vibrating sample magnetometers for the various electrodeposited nanocrystalline ferromagnetic matenals. Beginning in 1991, the first magnetic meauirements on
and N i P were taken
on a vibrating sample magnetometer located at the Royal Military College in Kingston, Ontario, Canada. Subsequently, a hybrid vibrating sample magnetometer was developed (see section 4.1) at Queen's University, Kingston, Ontario, Canada where the remaining magnetic measurernents on Ni-Fe, Co and Co-Fe were completed. It should be noted that in 1991, when these measurernents were initiated, relatively Iittle was known about the magnetic properties of nanocystalline ferromagnets.
For this reason, the literature review for the magnetic properties of nanocrystahe matenals that appears in section 3.5.4. reviews the state of scientific understanding until 1992 when this research was initiated. As the research was completed over the next seven years, new literature relevant to each system, that has become available fiom 1993 to present, has been added to the appropriate sections.
The thesis is organized in the following manner. Chapter two entitled: "Theory of Magnetism" presents magnetic theory relevant to this thesis. Chapter three provides an overview of nanocrystalline materials starting with the structure and synthesis and finishing with a current literature review of some of the various properties of nanocrystalline materials. It should be reiterated that the magnetic properties are only reviewed in this section up until this research was initiated in 1992. Chapters four, five, six, seven and eight represent the original research completed by the candidate. Chapter four discusses the development of a hybrid vibrating sample
magnetometer used to magnetically characterize samples in the majority of work completed in this thesis. Chapter five is a chronological presentation of al1 systems characterized in this thesis. The sections in chapter five are designed to stand alone, each with its own review of literature fiorn 1993 to present, experimental, results and discussion, conclusions and reference section. Chapter six discusses results of electroplating experirnents done in a magnetic field to assess the effects of a magnetic field d u ~ plating g on the properties of nanocrystalline Ni-Fe and Co. Chapter seven discusses the rewlts of thennomagnetic studies completed on Ni, Ni-P and Ni-Fe including the effects of temperature on the behaviour of sahiration magnetization and coercivity. In sorne cases the Curie temperature was measured for these materials. Chapter eight is a discussion chapter designed to give an andysis of results and present theories and trends for nanocrystalline magnetic materials in general. Chapter Nne States
the conclusions of this thesis and chapter ten provides recommendations for fuhire work.
References T. Watanabe, Res Mechanica, 11,47, (1984).
KT.Aua and G. Palumbo, Structure and Properties Relationshi~sfor Interfaces, eds. J.L.Walter, A.H. King, K. Tangri, ASM, 1, (199 1). KT.Aust, U. Erb, G. Palumbo, in Mechanical Properties and Deformation Behaviour of Matenals Havinp Ultra-Fine Microstructures, eds., M. Nastasi et al., Kluwer Academic Publ., 107,(1993). G. Palumbo, KT.Aust, U. Erb, P.J.King, AM. Brennenstuhl and P.C.
Lichtenberger, Phys. Stat. Sol.(a), 131,425, (1992).
G. Palumbo et al., in i
4s, . G.W.Bailey et al.,
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AM. El-Sherik, U. Eh,V. Krstic, B. Szpunar, M.J. Aus, G. Palumbo,and K.T. Aust, MRS Symp. Proc., 238, 173, (1993).
U. Erb, A.M. El-Sherik, G. Palumbo and K.T. Aust, Nanostr. Mat., 2,383, (1993). Hydroscope, 34,4, Feb, Ontario Hydro Corporation, Canada, (1997).
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M.R Fitsimmons, J.A. Eastman,R A . Robinson, A.C. Lawson, G.H. Kwei, K.E. Sickafus, M.A. Nastasi and E. Burkel, Nanostr. Mat., 3, 3 11, (1993). L. Daroczi, D.L.Beke, G. Posgay, G.F.Zhou and H. Bakker, Nanostr. Mat., 2,
515, (1993). M. Li, R. Bimnger, W.L.Johnson and R. S. Shull, Nanostr. Mat., 3,407, (1993).
G.N.Glavee, K.J.Klabunde, C.M.Sorenson, G.C. Hajipanayis, Z.X.Tang and L. Yïping, Nanostr. Mat., 3, 391, (1993).
T.D.Xiao, Y.D.Zhang, P.R.Strutt, J.I. Budnick, K. Mohan and K.E. Gonzalves, Nanostr. Mat., 2, 197, (1993).
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York, (1969).
Chapter Two THEORY OF MAGIVETISM IN CONVENTIONAL POLYCRYSTALLINE MATERIALS
2.1
Historieal Perspective
The first examples of magnetic materials were variations of magnetite or lodestone (loadstone) which consisted of various oxides of iron typically found in crystalhe igneous rock masses consisting of FeO, Fez03 and Fe04. These materials were abundant in an area called Magnus or Magnesia and were named magnets by early Greek philosophers[l]. Prior to this, magnetic materials including steels for tool
production were found in rneteorite deposits long before the art of steelmakhg was known. Ancient mariners made use of the magnetic properties of loadstones and compasses for navigation of their ships as early as 1000A.D.in China and 1100A.D.in the Western World. With regards to navigation by compass, Christopher Columbus, (Cnstoforo Colon) noted that the compass error changed fiom easterly to westerly upon crossing the meridian indicating a geomagnetic north differing fiom true north[l]. The first historical recording of magnetic materials in the Mediterranean dates back to the 13" century when Magister Petrus Peregrinus de Maharnecuna (Peter the
Pilgrim) sent a letter to a feUow mathematician regarding the development of an apparatus for studying magnetic bodies that he called a terella. This spherical sirnulacmm of the earth was noted to effect other such magnetic bodies. The various
magnetic interactions between the terella were studied by measurùig angles and distances[l 1. There were no major works on magnetism published until Gilbert's treatise on magnetism entitled De Magnete[2]. Gilbert was the first to note wrrectly that the south pole of a magnet pointed to the earth's geomagnetic north pole. Furthemore he suspended a magnetized needle horizontally on a fine wire which he called a versorium.
This apparatus was used to measure inclination and declination or the vertical component of magnetic fields. Much later, in 1715 the versorium was successfully used to measure the magnetic strengths of lodestones at distances up to nine feet. Gilbert's other
important scientific findimg was that iron would lose its magnetic properties when heated
and could regain thern upon cooling and hammering in an appropriately oriented magnetic field[l]. By 1600 it was also noted that keepers or sofi iron pieces on the ends of a lodestone or piece of magnetite would enhance the magnetic properties and "keep" the matenal magnetized[ 11. In 1820 there were three significant discoveries which finally established the
relationships between electricity and magnetism. A definite relationship was previously hypothesized, as ship's compasses would often point south after being stmck by lightening which Franklin had previously show to be electrical in nature. Firstly, Hans Christian Oersted notd that a current canying wire deflected a compass needle at right angles [3]. Subsequently, Arnpere noted in 1820 that a loop of wire canying cutrent behaved in the sarne fashion as a magnet thmst in a loop with an initiai zero current[4]. Finally in 1820, the first electromagnet was created by Dominique Francois Arago in a helid u>il(l]. As a point of interest, the first application of electromagnetism was
medical in nature and involved the removal of steei splinters fiom the eyes of needle sharpeners in 1822. In 1833, Gauss improved on the versorium for measuring field strength by constructing a torsional pendulum of magnetic material as a dynamic measurement technique[S]. The next major work on ferromagnetic materials arrived in 1878 when Ewing pubfished work in The Electrcian regarding how magnetic properties of materials affectecl transmission lines[l]. J.A. Ewing studied the ferromagnetic properties of iron as
they apply to transmission lines in Japan and paved the way for a succession of Japanese scientists with considerable expertise in ferromagnetic matenals[l]. Pnor to 1900, the only materials known to be measurably magnetic were of the ferromagnetic class and consisted of iron, cobalt, nickel, certain alloys and steels such as tungsten steel and iron dicon. M e r 1900, alternating current and transfomers were discovered to lower resistivity and transmission losses over long distances and as a result, core materials received a great deal of attention. It was found that insulated iron bundles or rods lowered core losses and the first core materiais used in A.C. transfomers were made of paper wrapped iron hat stiffeners. It was noted early on, that material with an increased resistivity would lower core losses in cyclic applications and the birth of Fe-Si aîloys occumed shortly &er 1900[1].
Modern commercial magnetic materials for hard and soft magnetic applications will be discussed in greater detail in section 2.7 after the relevant theoiy of rnagnetism is discussed in the forthcorning sections.
2.2
Atomic Magnetism
It is wel lmown that e l e d k charges in motion are responsible for the production of
magnetic fields. Hans Christian Oersted first discovered the field produced by current carrying wires in 1820[3]. The magnetic moment produced in spontaneously magnetized materials (permanent magnet materials) arises as a result of the motion of electnc charges (electrons) at the atomic level. Magnetic moments at the atomic level arise from three distinct sources: the spin of the electrons, their orbitai angular motion about the nucleus and the change in anpuiar motion as a result of an applied field. The electron spin and orbital angular motion act to increase the magnetic moment (paramagnetic and ferromagneticeffect) of the atom. The change in angular orbital motion as a result of an applied field acts to decrease the moment (diamagnetic effect) of the atom[6]. The most basic unit of magnetic moment is defined by quantum theory as the magnetic moment produced by the spin of one electron. This fundamental, naturai moment is calleci the Bohr magneton (pB)and is calculateci as follows [7]:
where e is the charge of one electron, m, is the mass of the electron, h is Planck's constant, and c is the velocity of light.
According to quantum theory, electrons may only have positive or negative spins (S= +/- %), which order themselves on the atomic level to minimize energy States. The
magnetic moments in energy levels, which contain an q u a 1 number of electrons with positive and negative spins, will act to cancel each other. For example, helium has two electrons, whkh spin and therefore orbit the nucleus in an opposing fashion. The end result of the paired electrons is that the heiium atom displays no net magnetic moment. Upon application of a magnetic field to helium, magnetic moment arises in the direction opposing the applied field only, which effectively explains why helium cm be classified
as a diamagnetic material. The case for atoms containing u d e d orbits is quite dserent and explains the presence of magnetic moments in the common ferromagnetic materials, iron., cobalt and nickel. In these atoms, it is the unfiued 3d shell which is responsible for the magnetic moment[7]. Electrostatic energy factors dictate that the electrons within a shell under certain conditions will favour parallel spin alignments within unfilled orbits. In the case
of iron with 26 electrons the 1s, 2s, Zp, 3s, 3p and 4s shells are filled whereas the 3d sheb contains five electrons with parallel spin and one electron opposing this spin. These four unpaired electrons are responsible for the ferromagneticbehaviour and the net magnetic moment present in iron. This net moment in the neutral iron atom is therefore equal to approximately 4 electrons (with spin = %) giwig two Bohr magnetons or 2 ~ ~ [ 7 ] .
There also exists a contribution to the net magnetic moment in a free iron atom as a result of angular orbital momentum. However, when iron atoms appear in a lattice, the
electron orbits become relatively fixed and are not influenced to a great extent by a magnetic field and are expected to cancel one another[7]. It is now possible to classify magnetic materials on the basis of their electronic
structure and inherent magnetic interactions.
2.3
Classification of Magnetic Materials
Certain types of materials contain atoms that may exhibit a magnetic moment. As previously stated, these moments orighate from three distinct sources: electron spin, electron orbital motion, and in the case of an applied magnetic field, the change in electron orbital motion. These moments may align in either unordered ways to form diamagnetic, paramagnetic or superparamagnetic structures or in the foliowing types of magnetidy ordered stmctures within a solid: ferromagnetic, antiferromagnetic, ferrimagnetic, and in special cases, helimagnetic. In the following brief explmations of magnetic structure, an effort was made to introduce the ciassical theory developed to explain magnetic phenornena and follow up with a cornparison to modem quantum
mechanical analysis. Susceptibility ( x ) is defined as the ratio of magnetic moment per volume (magnetization) to the applied magnetic field and it is possible to classi@ magnetic materials by their x and the variation of x with temperature 181.
2.3.1 Diamagnetism
Diamagnetism in general occurs in a material that has no permanent dipole moments with atoms in which a1l of the electron shelis are filled. Atoms with fùll electron shells do not exhibit magnetic contributions due to electron spin and electron orbital motion and as such do not exhibit net magnetic moments. The diamagnetic contribution to the magnetization occurs as a result of the change in orbital moment
induced by an applied magnetic field. Diarnagnetism results when an extemal magnetic field is applied to an atom, and electncai charges act to shield the atom interior[6]. Lenz's law states that when an electric circuit flux is changed, induced currents act in opposition to the direction of the flux change. The atom may be viewed as an electnc circuit with electrons orbiting the nucleus. Therefore, when a diamagnetic material is subjected to an extemal applied magnetic field, it exhibits an induced magnetization opposing the applied field, indicating a negative susceptibiity (x). The magnetization
behaviour is given by the Langevin theo~yof diamagnetism 191:
(eqn 2.2)
where N is the number of atoms per unit volume, Z is the number of electrons per atorn, e is the electronic charge, m, is the electronic mass and 3> is the root mean square atomic radius. It should be noted that the quantum mechanical analysis agrees with the classical Langevin result161. Diamagnetic materials include: Cu, Au, Bi, Be, B, S, noble gases, most organic compounds diatomics such as N2and Hz0 and iondsalts such as Na+ or Cr [8]. Materials that exhibit a positive susceptibility (pararnagnets and ferromapets) still have a diamagnetic effect or component, which is essentially independent of
temperature. However, this effect is quite small when compared to the effects of extemal fields on atoms which exhibit net spin or orbital angular momentum.
2.3.2 Paramagnetism
Paramagnetic stmctures develop in materials that exhibit permanent dipole moments but have no neighboring moment interactions. In these materials the electron spin and electron orbital angular momenturn contribute positively to the magnetization. This positive contribution leads to a positive susceptibility (x). In general, the dipole moments of a paramagnetic materid are oriented randomly in zero applied field[8]. Paramagnets becorne slightly magnetized in an applied magnetic field as they have a slightly positive susceptibility (x). Furthermore, this magnetization is inversely proportionai to the temperature. In terms of classicai theory rnaterials, in which the orbital moments are unbaianced due to unpaired electrons a permanent magnetic moment exists at the atornic level. The net magnetic moment is the sum of the orbital and spin components[8].
Boltzman statistics can be used to analyze the moment based on Langevin's assumptions that the neighbonng moments are non-interacting and that randomness increases with increasing thermal energy. Evaluating the probability function gives the Langevin equation for paramagnetic rnaterials, which can be arrangecl in terms of susceptibility as given in section 2.3.1 for diarnagnets:
(eqn 2.3)
where N is the number of atoms per unit volume,
is the pemeability of free space, m
is the net magnetic moment, kB is Boltunan's constant, and T is temperature. It can be
,
readily seen that this mode1 indicates that the susceptibility of these materials varies inversely with temperature. It should be noted that for materials with other magnetic stnichires, which transfonn to paramagnetic above a critical temperature (Curie Tempemaire) a modification of the above theory is required. The quantum theory of paramagnetisrn can be divided into single electron atoms, meaning atoms with one unpaireci electron and multiple electron atoms. In the case of single electron atoms, the expression for the bulk magnetization is also evaluated using Boltzman statistics. However, the initial expression for the magnetic moment due to a localized electron has changed to [IO]:
where g is the Lande or spectroscopic splitting factor, p~is the Bohr magneton, and J is the total angular rnomentum of the isolated atom. It should be noted that g is derived fiom orbital and spin angular moments in unfiiled electron shells as follows[lO]:
(eqn 2.5)
where S is the atomic spin momentum and L is the atomic orbital momentum.
Substituthg equation 2.4 into 2.3 yields the quantum theory of paramagnetism for single electron atoms [ 1O] :
(eqn 2.6)
The multi-electron atom situation is slightly more complex as the energy levels for multi-electron atoms are equal to 2J + 1 which changes the final result of the Curie law of paramagnetic susceptibilities to[lO]:
(eqn 2.7)
It can be seen that quantum mechanical treatrned yields similar results to the classical expression for susceptibilities of paramagnetic materials. For exploration of the quantum theory of diamagnets and paramagnets in greater detail the reader is referred to Chapter 14 of Kittel[6].
Paramagnetic materials which display random dipole moments with no neighboring moment interactions include: Al, Pt, Mn,Cr, the diatomic gases O2 and NO
as well as some rare earth metals and rare earth oxides[8]. In general, paramagnetic magnetism is stronger than diamagnetism, however the
overall magnetic effect is small as only a small percentage of atomic moments align in conveniently accessible extemal fields. Thermal effects act to oppose alignment of atomic moments in paramagnetic materials. Therefore the overall increase in magnetic field is small in paramagnetic materials.
2.3.3
Ferromagnetism
Ferromagnetic materials are the most usefùl class of magnetic materials and have the distinct ability to acquire and maintain large magnetic fields upon application of relatively small magnetic field@]. Ferromagnetic materials are a subclass of paramagnetic materials in which the atoms exhibit dipole moments and there exist magnetic interactions between neighboring atoms. As a result of these interactions, ferromagnetic materials tend to have regions containing atoms with sidarly aligned dipole moments called domains. It is possible for a ferromagnetic material to have many domains (large areas of cornmon aiignment) which sum vectorially to give an overall magnetization in the material of zero in zero applied field. On the other hand, when al1 magnetic moments in a material are similarly digned the material is magnetically saturated. Magnetic saturation occurs readily in most ferromagnets subjected to extemal magnetic fields of 1 Tesla or les. However, when the field is removed, the magnetization may attain any value fiom near saturation to zero. When a ferromagnetic material maintains a magnetic moment in the presence of zero applied field it is spontaneously magnetized. Weiss interaction theory or molecular (Weiss) field theory was developed to expiain this phenornena[ 101. Weiss postulated that atorns exhibiting localized dipole
moments interacted with neighboring atoms, which allowed for the formation of magnetic order below critical temperatures for which thermal agitation is not sufficient to randomize dipole alignments parametrically [IO].
.
If dl magnetic interactions within an atom are assurned to be equivalent, an expression basecl on a mean-field exchange interaction parameter may be utilized to express the Weiss field developed within a particular magnetic domain. Within any domain the dipole moments of the atoms are digned and the material is locally sahirated
'
which yields the following expression for the Weiss field[lO]:
(eqn 2.8)
where a is the mean-field interaction, Msis magnetic saturation and mi is the moment that an atom experiences as a result of the field produced by another atom.
Ln zero applied field the only field operating is the Weiss interaction field and using Langevin's analysis for paramagnetic materials, with no constraints on the direction of the magnetic moment (rn) the following result is obtained[lO]:
(eqn 2.9)
This lads to the description of magnetic dipole moment behaviour as a function of temperature as shown in Figure 2.1. This function shows that at absolute zero al1 moments are aligned and the material is rnagnetically saturated. As the temperature approaches the Curie Temperature, the spontaneous magnetization falls quickly to near
zero values. At this point there is a transition in the magnetic structure fiom ferromagnetic to paramagnetic as the thermal energy overcomes the exchange interaction. Table 2.1 shows the Curie temperatures for various ferromagnetic elements [IO].
.
Table 2.1:
Curie points for various ferromagnetic elements [IO].
Figure 2.1:
Variation of the spontaneous magnetization M. of nickel within a domain
as a fiinction of temperature according to the Weiss mean field mode], after Weiss and
Forrer (1929)[10].
When this interaction is analyzed in tems of the Langevin function for ferromagnetic materials above their critical temperature (afler undergohg transition to
paramagnetic) the magnetintion behaviour becomes [IO]:
(eqn 2.10)
It shouid be noted that this equation takes the form of the Curie-Weiss law, which
describes the magnetization behaviour above Curie Temperature as shown in the following equations[8] :
(eqn 2.11)
where C is the Curie constant given by [IO]:
and Tcis the Curie Temperature given by [SI:
(eqn 2.13)
21
The Curie-Weiss law is derived fiom the theories of ferromagnetism and demonstrates that below the critical temperature, paramagnetic behaviour ceases to exist. It should be noted that this mean-field approximation is only valid within domains, whereas a more accurate approximation involves the nearest neighbor magnetic interactions only. In this case the derivations are the sarne but a is replaced by its nearest neighbor (z), in combination with the localized interaction parameter J. Quantum mechanicd treatment of the problern yields the behaviour of the
magnetization in the case of single-electron thermodynamical statistics [1 O]:
(eqn 2.14)
and, in the case of multi-electron atoms [IO]:
(eqn 2.15)
where Bi is the Brillouin function which is the quantum mechanical analog to the Langevin function [7]. Sirnilarly to the classical treatment, at higher temperatures, the rnaterial becomes paramagnetic which yields the susceptibiiity in the following form of the Curie-Weiss Law [IO]:
(eqn 2.16)
Quantum theory cm also be used to describe the magnetic behaviour within a domain in a ferromagnetic material as a function of temperature.
Figure 2.2 shows the
daerences between the classicd derivations and the quantum mechanical derivations for the main ferromagnetic materials Co, Ni and Fe as a function of the temperature for
different total atomic angular momentum quantum numbers.
ABSOLWTE TEMPERATURE. f CURIE TEUPERATVRE 8
,
Figure 2.2:
Temperature dependence of spontaneous magnetization within a domain
for iron, cobalt and nickel compared with calculations based on the ferromagnetic Brillouin fùnction [1O].
.
As stated previously, ferromagnets comprise the group of materials that are the
most useful and widely used in most magnetic applications. The notable ferromagnetic metals are the transition metais Fe, Ni, Co (and their alloys) as well as rare earth metals with atomic numbers fiom 64 to 69.
Al1 of the nanocrystalline materials studied in this thesis are ferromagnetic in nature and are pure metals or alloys of the transition metals noted above. The reason these elements are ferromagnetic lies in their atornic structure. Iron, cobalt and nickel aII have unfilled 3d electron bands required for magnetic moment as well as a full 4s band which has paired electrons and does not change the magnetic properties. The Pauli exclusion principle and Hund's Rules as describeci by Kittel [6] define the electron occupations of Co, Fe and Ni. The non-integral moments observed on the Slater-Pauling curve (shown in Figure 2.3) of these 3d elements and their alloys can be explained with band theory also known as the itinerant electron theory.
Figure 23:
The Slater-Pauling curve, which gives the net magnetic moment per atom as a function of the number of 3d electrons per atom.
2.3.4
Antiferrornagnetism
Antiferromagnets were thought to be a subclass of paramagnets owing to their small positive susceptibilities. However, it was later found that these materiais have a definite arrangement of dipole moments, which is effected by nearest neighbor interactions. The spin configurations of dipole moments in antiferromagnetic materials
are antiparallel as the neighboring interactions are negative[8]. Antiferromagnetic materials obey Weiss type interaction behaviour above their critical ordering temperature known as the Neel temperature. Above this temperature, antifemomagnets may be analyzed as being compriseci of two sublattices of paramagnets with equal and opposite dipole moment configurations (antiparallel). If we assume that each lattice has the same Curie constant C as both lattices have the sarne magnetization
and take p as the mean-field interaction parameter then we cm solve for the susceptibility as follows 161:
This is simply the Curie-Weiss law in a form containing TNas the cntical temperature, which is positive, and it applies only for temperatures above the critical temperature. Therefore above the Neel temperature eqn 2.17 applies but below this temperature the susceptibility decreases with decreasing temperature in a more complex fashion.
Certain rare earth elements such as terbium or dysprosium display both Curie and Neel temperatures When the dipole moments are aligned parallel (and also antiparallel) to an.
extemai magnetic field, halfof them are aligned favorably and the other haif are held rigidly antiparallel.
Figure 2.4:
Arrangement of magnetic moments in an antifemomagnet at T=OY (a) in zero applied field and @) in a field, H, applied perpendicular to the
moments [8].
When the field is applied perpendicularly to the dipole moments, the moments
will rotate slightly out of their aiignment as shown in Figure 2.4 which will cause the susceptibility in this case to vary between the maximum at the Neel temperature and zero. Antiferromagnetic materials are ofien transition metal oxides of the NaCl type stnicture such as MnO,Co0 and Cf2O3,Feû, Fe203and NiO. Some compounds from the haiide group also form antiferromagnetic materials such as FeF2, FeC12, CoClz and NiC12.
Ferrimagnetism has components of ferrornagnetism and antiferromagnetism in that the material exhibits dipole moments with neighborhg interactions which are antiparailel. However, the magnitude of the dipole moments is not equal which gives nse
to the foUowing modifications to the Curie-Weiss law previously discussed for antiferromagnets. As the magnetic moments and Curie constants are different for the different atoms in each sublattice we introduce MA, CAto describe one lattice and Mg,Cg to describe the other. The ha1 result is presented here but Kittel provides a full treatment [dl.
(eqo 2.18)
In ferrimagnetic materials as with ferromagnetic materials above the Cune temperature the material becomes paramagnetic. Another similarity to ferromagnetic
materïals, is that the dipole moments of fenimagnetic materials self-organize into domains and exhibit hysteresis and saturation[lO].
The oldest magnetic material, magnetite or Fe304is femmagnetic in nature as are other mixtures of elements and bon oxides such as the cubic iron gamets, which take the
form M3Fe50i2. Hexagonal ferrites such as SrOa6(Fe203)and BaOe6(Fe203)femtes have been used as hard magnetic materials as a result of their large anisotropy and high thermal stability or critical temperatures between 500 and 800°C[10].
2.3.6
Superpararnagnetism and Hdimagnetism
Superpararnagnetism describes the magnetic behaviour of certain materials, which
contain very small particles. If the particles are extremely small, they may not have the dimensions required to maintain domain walls (See section 2.4.5). In addition, the thennal agitation may be high enough to overcome the anisotropy which wiil lead to
randomly aligned dipole moments that rotate spontaneously. These materials have no hysteresis and unusually high susceptibilities, which allow saturation in srnail extemai fields. The moment is much larger than single atom moments, which means that groups of the dipole moments are rotating in unison. These particles exhibit magnetism in which the net magnetism follows the Curie-BriUouin-Langevin derivations for pararnagnets. Helimagnetism describes a variant of ferromagnetic order that occurs in some elements. If the ferromagnetic order of successive base planes is inclined at some angle with respect to the next basal plane, helimagnetic ordering may result. When the angle of
the successive basal plane rotation is calculated, it may be used to derive an expression . for minirnized energy providing stable helimagnetic structures. This type of magnetic ordering can occur in terbium, dysprosium and holminium.
.
2.4
Energy Contributions to Magnetic Domain Formation
As stated in section 2.3.3. domains form in ferromagnetic materials. The structure of
the domain is a consequence of the various contributions to the energy of a ferromagnetic body. in a magnetic field, the overail energy as a result of the applied field is given by the following equation:
E=E,+E,+E,+E,+E, where
(eqn 2.19)
is the exchange energy, E, is the magnetostatic energy, Ei, is the anisotropy
energy, E,is the magnetoelastic energy, E,is the domain wall energy and EH is the magnetic field energy. The rninirnization of the sum of equation 2.19 helps us to theoretically determine domain structures, which have been experimentally observed. An example of domain structure is show in Figure 2.5.
Figure 2.5:
Magnetic domain structure in the sufice of iron [Il].
.
2.4.1
Exchange Energy
The exchange energy descnbes the spin alignment preference of electrons as a
result of nearest neighbor interactions. The expression for the Weiss field exchange interaction parameter was discussed in section 2.3.3. Therefore, the energy per magnetic moment due to the exchange interaction is given by the following equation [IO]:
(eqn 2.20)
where mi and mj are the neighboring moments, and since they are equal they may be replaced with an m2 tenn. Furthemore, this energy is dependent on the tum angle of the moments between the neighboring atoms. Therefore, a $ term describing the tum angle is added and the analysis is based on interaction in a linear chain of two interacting moments per atom which lads to the sum of interactions over a domain wall given by the following equation [1O]:
(Eqn 2.21)
Where n is the number of atoms or moments in the wall. This equation shows that the minimization of this energy requires a small angle between the neighboring moments,
which lads to wide domain wdls. The energy of domain walls is usually rneasured per length across the domain wall.
.
The exchange energy c m also be defined quantum mechanically. Assuming a is the nearest-neighbor interaction parameter, the quantum mechanical expression for
exchange energy (J)becornes [IO]:
E, = - p o d m Z = JS2
(eqn 2.22)
where S is the number of electron spins per atom and J is the exchange integral.
2.4.2
Magnetostatic Energy
The magnetostatic energy is sometimes calied the demagnetizing energy and it is the energy of the sample in its own field. When the atomic dipole moments are
perpendicular to a surface, demagnetizing energy exists which leads to the strong dependence of the magnetostatic or self-energy on geornetiy. This demagnetizing factor (Nd)is based on geometry and is created in a sample whenever magnetic dipole moments exist. The intemal field opposing magnetization in a magnetic material is proportional to the moment and may be represented by [IO]:
(eqn 2.23)
-
Furthemore, the energy of this demagnetizing field may be taken as the integral of the
field over the volume of the dipole of magnetization, which yields the following [IO]:
(eqn 2.24)
This is the self-energy of a particle with al1 moments aligned in one direction, or a single
domain particle. The demagnetinng energy is located at the surfaces or ends of a sarnple and does not extend through the sarnple as domains within the sample negate the
dernagnetizing effect. As such, the self-energy can be reduced fùrther if more domains are introduced in the material. This lowers the interna1 field opposing rnagnetization as
well as the demagnetizing factor. The domain wall size and hence number of domains in a sample is determined by the domain wall energy as discussed in section 2.4.5.
2.4.3
Anisotropy Energy
Anisotropy energy sometimes calleù magnetocrystaiiine energy &ses as a result
of the asymmetncal overlap of electron orbits in crystalline materials. This in turn causes the magnetization in ferromagnetic materials to be directed dong certain crystallographic .
axes. These axes of preferred magnetization are called the easy directions. Figure 2.6 shows the variation in magnetization behaviour with respect to various crystaIlographic
directions for Fe and Ni.
B, (gauss)
Figure 2.6:
Magnetization curves for a single crystal of cubic iron and cubic
nickel with the applied field parallel to various crystallographic directions.
Depending on the crystal structure and hence symmetry the anisotropy energy
will take various forms. In the case of iron and nickel we are dealing with a cubic system
.
which can be represented in the following manner. Cosines (a)of the angles between the direction of magnetization and the three cube axes are used to describe the magnetization vector such that [8]:
(eqn 2.25)
.
The energy of anisotropy can be determined fiom a relation that is expressed as a
power senes. The mathematical combinations required to satis& symmetry requirements in a cubic system indicate that only even powers are used as the magnetization is reversible and the second power term is siiply eqn. 2.25. Therefore the fourth and skth terms in the power senes give us the relation between magnetocrystalline energy and anisotropy constants[8]:
(eqn 2.26)
Cobalt has a hexagonal stmcture with an easy direction of magnetization along the basal plane. Since the magnetization in cobalt is uniaxial we only need to consider the angle
(8)between the direction of rnagnetization and the easy direction of rnagnetization. This leads to an even power series of sin 8:
E, = ~ , s i n ~ @ + K # n * B
(eqo 2.27)
The anisotropy constants for the ferromagnetic materials studied within the body of this work are listed by Table 2.2. Table 2.2:
Material
Anisotropy constants for iron, nickel and cobalt [IO].
Ki
K2
Iton
0.480
-0.050
Nickel
-0.045
0.023
Cobalt
4.1
1.O
As the anisotropy of cobalt is much higher than iron and nickel we expect that the.
magnetization differences between easy and hard directions in cobalt will be much pater. Experimentaliy,this has been shown to be true as illustrated in Figure 2.7.
Figure 2.7:
Easy and hard magnetization curves for cobalt at room temperature.
2.4.4
Magnetoelastic Energy
Magnetostatic energy, also called magnetostriction (A.), is the energy which arises fiom the interaction between magnetization and strain of the crystal lattice. Similar to anisotropy energy, when the lattice is strained, the electron orbitals of the electrons interact in an unsymmetrical way leading to magnetoelastic effects. In strained lattices, the magn~toelasticenergy E, increases linearly with
increasing lattice strain. Unstrained lattices have zero magnetoelastic energy. The magnetostriction constant (h)is defined as the fiactional change in length (1)[1 O]:
il= dl//
(eqn 2.28)
In cubic materials there are two magnetostriction constants: hioo, hi 11. Furthemore, strains as a result of the magnetoelastic effect are defined along the cubic matenal principle axes as follows[lO]:
where ai,az,as,are direction cosines of the field relative to the axis along which the moments are saturated and Pi,P2,P3 are direction coimes relative to the field direction measuring saturation magnetostnction. Materials with positive magnetostriction coefficients expand when they are magnetizd and magnetization is increased in the direction of the coefficient under an applied stress.
2.4.5
Domain Wall Energy
The domain wall energy is difncult to calculate as the spins of the dipole moment in the domain wall are directed away from the axis of the easy direction in the material.
The exchange energy and anisotropy energy are both associated with the energy of the domain wall. In addition if the lattice is stressed there will be a magnetoelastic effect leading to an,,E
term in the domain wall energy.
The domain wall enera is the difference in energy between dipole moments in a
dornain and those in a domain wail and is given as energy per unit a r a of the wall. The exchange energy must compete with the anisotropy energy to detemine the structure of the domain wall. If the thickness of the wall is defined as b, and a is the lattice parameter, the sum of these competing energies will determine the dornain wall energy expressed per unit area as follows:
(eqn 2.30)
Equation 2.30 shows that when the misotropy terni is dominant, a thinner domain wall will rninimize the domain wall energy. Conversely, if the exchange energy term is
dominant, a large domain wall where the change in dipole moment orientation takes place over possibly several hundreds of atomic layers will be favoured.
Domain walls in iron are about 4nm in width, which l a d s us to speculate that as we decrease the grain size of materiai into the nanocrystailine region, magnetic behaviour
will be altered due to energy considerations.
2.5 Magnetic Properties Derived from Hysteresis Loops
The magnetization / demagnetization curve for a aven ferromagnetic material provides a great deal of information during magnetic characterization. The full magnetization curve, shown in Figure 2.8, and known as a hysteresis loop, illustrates the methodology as well as which properties of the material under scrutiny are revealed. The hysteresis loop has the axes of applied field (H) aiid magnetization of the
material (M). Ofien the magnetic induction @) is displayed as a function of applied field (H) but B is related to M by the following equation:
(eqn 2.31)
where ~ iis,the pemeability of fiee space.
The hysteresis loop is generated in the following fashion as s h o w in Figure 2.8. At point (a) the material is in the virgin or demagnetized state and incurs an applied field (H) uniil the material achieves magnetic saturation at point (b). Frorn point @) the field is reduced to zero and the magnetization remaining in the material with zero applied field (remanence) is shown at point (c). From point (c) a negatiw field is applied to the
sarnple and the induced magnetization is reduced from (c) to (d) which corresponds to the arnount of field required to reach zero magnetization in a particular sarnple. Point (d) is known as the coercive point. The sample is negatively magnetized to point (e)
comsponding to magnetic saturation in the opposite direction. When the field is decreased to zero and the magnetization falls to the value at point (f). Lastly, the field is
applied in the positive direction to bnng the sample back to positive rnagnetic saturation
crosshg the applied field axis at point (g) and joining the initial saturation at point @). The area enclosed fkom point (b) wunterclockwise to point (b) is the hysteresis loop.
Figure 2.8:
Schematic diagram of M vs. H or hysteresis ioop showing magnetic properties and domain structures during difEerent stages of magnetization.
Hysteresis cornes from the Latin verb 'hyster', meaning 'to lag' . The area enclosed by the hysteresis loop represents a property h o w n as the energy loss ('WH) in cychg the
magnetic field. There are several rnitigating factors that determine the shape and size of the hysteresis loop and al1 of the properties reveaied by the hysteresis loop.
An important property for characterizing ferromagnetic materials and their subsequent application is determined in part by the susceptibility or permeability of a materiai. The susceptibility ( x ) is defined by the foiiowing:
(eqn 2.32)
where M is the magnetization and H is the applied field. Similady, the pemeability (p) is defined in the following equation:
(eqn 2.33)
The susceptibility and pemeability are defined as the slopes of M vs. H or B vs. H respective behaviour on hysteresis loops. As these slopes change throughout the magnetization process, the most usefùl information is the initial susceptibility (initial pemeability), xi (pi) which occun at point (a) in Figure 2.8 and the maximum susceptibility (pemeability) X , (bwhich ) ocairs at points (d) and Cg) wrresponding
to the coercive points. The susceptibilities give us important information regarding the
ease at which a material may be rnagnetized. We have already seen how anisotropy can
change the shape ofthe initial magnetization curves in the ferromagnetic transition elements shown previously in Figures 2.6 and 2.7. Saturation magnetkation (Ms) describes the limit of magnetization in any given material. This occurs when al1 of the dipole moments in the material are aligned in the direction of the applied magnetic field. Magnetic saturation occurs at point (b) in Figure 2.8. It should be noted the saturation magnetization depends only on the volume of
atoms being magnetized and is expected to be a structure insensitive property [IO]. Retentivity (MR) or remanent induction (BR)is the amount of magnetization
remaining in a material afler a saturating field has been removed. The retentivity is given by point (c) in Figure 2.8. Permanent magnet theoy makes use ofthis property
sometimes called the retention of magnetization to distinguish ferromagnetic materials fiom paramagnetic materials [ 1O]. The coercivity (He)is defined, as the applied field required to demagnetize or to bring the magnetization of a saturated material to zero. The coercivity or wercive point is given by point (d) in Figure 2.8. It should be noted that the width of the hysteresis loop is dehed as twice the coercivity and as such, the coercivity will determine to a great deal the magnetic properties of a material. The domain structure of the material is magnetically saturated in the direction of the applied field at point @) in Figure 2.8. The domain structure changes when the
magnetic field is applied in the opposite direction to have zero net magnetization at point (d) as shown in Figure 2.8. At point (e) in Figure 2.8, al1 dipole moments are aligned in
the direction of the applied field.
There are various contributions to the hysteresis of a ferromagnetic material. For instance, a cold worked specimen will increase the hysteresis core loss (WH) and
werciviîy (Hc) as show in Figure 2.9. Impurities, inclusions and dislocations in a matend act as pinning sites for domain wall movernent. This in tuxn provides an 'intemal fiction' during the magnetintion process that subsequently increases the hysteresis.
Figure 2.9:
Hysteresis loop behaviour for sofi iron and hardened steel [IO].
2.6
Hard and Soft Magnetic Materials
It is possible to broadly classiS, ferromagnetic materials as either hard or soft magnetic materials based on their permeabilities and coercivities. Figure 2.10 iliustrates the relative permeabilities and coercivities dong with the year of discovery for both hard
and soft magnetic materials. A hard magnetic material is one in which the matenal, once rnagnetized is difncult to demagnetize. Materials, which f d into this category, are used for permanent magnets and magnetic data storage or recording media. Hard magnetic materials typically have coercivities greater than 10kAlm[10]. Permanent magnetic materials are chosen for applications based on their 20d quadrant hysteresis loop properties or what is known as the demagnetization curve. The magnetic properties of permanent magnets are determined not ody by their composition but heat treatments and processing treatments during fabrication. One important pemanent magnet class, which was discovered in 1984, is the neodymium-iron-boron class of magnets. The Nd-Fe-B magnet is known for its extraordinarily hi&
coercivity,
which can reach values of 1.12 x 1o6 A/m. The other important magnetic property that determines the suitabiiity of a
permanent magnet material is the maximum energy product Pb).For Nd-Fe-B magnets the energy product cm be as high as 320 x 10~1/rn~[10]. Recordiing media
require high remanence and coercivity not unlie permanent magnet materials. Another requirement of matends for recordmg is a square hysteresis loop, which means that the
matenal has a high remanence and coercivity but is able to switch frorn one state to
another quickly. T y p i d materiais used for recording media are oxides of iron, chromium and cobalt. Soft magnetic materials find use in entirely dierent applications. Typical sofl
magnetic applications include: transformer cores, rdays, recording heads, electric motors
and electromagnets. Sofi magnetic materials are materials, which can be easily magnetized and demagnetized with minimal coercivity (Hc) and core loss (WH) in cyclic applications. Sofl magnetic materials typically have coercivities lower than I M m .
Electromagnets must have a high permbility and saturation but low coercivity so that the field direction may be easily reversed. Soft iron and alloys of cobalt and iron
are used almost exclusively which provide a saturation magnetization as high as 1.7 x 106
A,mand 1.95 x 106Alm respectively. As transfomers operate in an altemating current environment they must minimize core losses which are generated as a result of magnetic hysteresis and eddy current formation. Laminated structures cornpnsing, grain oriented iron-silicon alloys are usually used for transfomers as they have a high magnetization combinai with a low conductivity (due to Si) which reduces losses due to eddy currents. 1o7 @
a
mgm (inri
106
i
i
I
i
10'
10 . 1
,
1
i
!.
i
I .
j .
10.'
t
10.'
10" 10"
.
1
!XS ( j s m O. Alnko V ( 1 W e
yK (1911)
1
i
~~n-Putinum(rrw) SmCo!l%r)
j
10 *c
1
.
1v
to'
N&f*8(1U) I
10'
10J
Figure 2.10: Relative permeabilities and coercivities of ferromagnetic materials [1 O].
2.7
References
[1 ]
L.W. McKeehan, Mamets, D. Van Nostrand, Canada, (1 967).
123
W. Gilbert, De Mapnete, Londii Excvdebat Petru Short Anno MDC, (1600) as
.
cited in [l]. [3]
H. C. Oersted, Bibliotheque Universelle des Sciences, Belleslettres, et Arts 14, 274, (1 820), as cited in [l].
[4]
A.M. h p e r e , Annales de Chem. et Phys.,(2), 15,93, (1 820) as cited in [Il.
[5]
K.F.Gauss, Annalen der Physik, (2), 28,241, (1833) as cited in [Il.
[6]
C. Kittel, Introduction to Solid State Physics 7h Edition, John Wiley & Sons, New York, 1996.
[7]
R.M.Bozorth, Ferromagnetism, D. Van Nostrand Co. Inc., New York, (1% 1).
[8]
J.P. Jakubovics, Magnetism and Magnetic Matenals, The Institute of Metals,
London, (1987). [9]
P. Langevin, h a l e s de Chem. et Phys.,5, 70, (1905).
[IO] D.Jiles, Introduction to Magnetism and Magnetic Materials, Chapman and Hall, London, (1991). Bozorth and W. Shockley, Phys. Rev., 75, 155, (1949) as [Il] H.J. Williams, R.M. cited in [10].
Chapter Three OVERVIEW OF NANOCRYSTALLINE MATERIALS 3.1
Grain Boundary Engineering in Conventional Polycrystals
3.1.1
Introduction to Grain Boundaries -1.-
The structure of grain boundaries has been a source of debate since scientists
decided upon the crystalline nature of metais in the late 19' century[ l,2]. The first theones proposed delineated a structure of atomicaiiy ordered grains separated by an
amorphous 'boundary cementY[3].Subsequently, theories were proposed which stated that grains meeting in a well-defined manner would contain a 'transitional lanice' which was a reproducible transition nom one grain to its adjoining neighbor[4]. Theories based on dislocations and coincident sites of atoms were developed to explain the periodicity of the grain boundary structure. Measurements of grain boundary energy along with the development of x-ray and high-resolution rnicroscopy wnfirmed the non-amorphous structure of grain boundaries[S-71. In contrast, it has long been known that some sintered ceramics, which are consolidated powders with a high volume fraaion of 'binder phase', wntain amorphous boundaries[8,9]. In addition, some polycrystalline (or nanocrystalline) materials produced fiom amorphous preairsor materials will contain residual amorphous grain boundaries[lO- 121. However, aside fkorn these two examples, it is generally assumeci that the stiucture of grain boundaries is ordered as shown by both theory and experiment. In fact, a great ded of research in the last nfteen years has been in the area of grain boundary design to improve various properties of a material. This thrust area
manifests itself in two ways: (1) by increasing the interfacial content of the materials through reduction in grain size (nanocrystaliine materials [13]) or (2) by controlling the
actual types of interfaces present in a polycrystalline materiai, to improve various properties[14].
3.1.2
Special Boundaries
In 1959, Aust and Rutter published key findings based on the studies of individual grain boundary mobilities in zone refined lead with various additions of tin which were in
amounts that formed a solid solution with lead[l5]. The addition of tin to the lead had a dramatic effect on the grain boundary mobility and Aust and Rutter were able to distinguish "special orientation relationships" which yielded fast moving boundaries. More specifically, the grain interfaces that were rotated: 36-42" about the < I l 1>, 23' about the 4 11> or 26-28' about the direction were not as effected by solute additions in the 0.002 to 0.006wt% concentration range of tin and were d l e d "special boundaries". The reason for the differences in the mobility of "special boundaries" as
compared to general boundaries was the exceptional lattice matching across the "special boundaries" which would not favour the accommodation of solute atoms[l5]. It should be noted that this was the first result, which described special properties of hi&coincidence interfaces or a low reciprocal volume density (low Ç) of coïncidence site
lattice sites.
.
3.1.3
Property Differences between Low C and General Grain Boundarie.
It is well known that grain boundanes with values of C< 29 display special properties[l6]. The following list is reprinted frorn a recent review by Aust and Palumbo and shows some of the properties of low C boundaries as compared to high Z, or general
Less susceptibiiity to solute segregation Lower energy in pure metals Greater mobility with specific solutes in a certain concentration range Smaller diffisivity Greater resistance to grain boundary sliding, fiacture and cavitation Lower intrinsic resistivity Greater resistance to localized corrosion
3.1.4
Grain Boundary Design
Watanabe first introduced grain boundary design concepts in l984[14]. Improvements in various properties of polycrystalIine materials were obtained through effective application of grah boundary design and interface control. For example, by increasing the low L boundary density in a material beneficial effects can be gained in the areas of: intergranular fiacture, embrittlement associated with high temperature creep,
intergranular cavitation, liquid embrittlement and high temperature sliding and
fiacture[l7]. Palumbo determhed that certain boundaries are more susceptible to intergranular fracture and derived a probabiiity fùnction for crack amest based on the fractions of susceptible and non-susceptible interface@ 81. Furthemore, Palumbo
determined that crack arrest would occur when al1 possible crack continuation paths (interfaces) were coherent twins[l8]. The 'twin lirnited' stnicture proposed by Palumbo represents a CSL distribution consisting entirely of Iow Z boundaries as a result of twinning operations.
This grain boundary distribution could contain a maximum of 213
annealing twins (X3 boundaries) where the remaining boundaries are twin related variants[19]. This grain boundary engineered structure is of great importance to
materials, which require intergrmulu degradation resistance and perhaps the best illustration of grain boundary design to date.
The previous description of grain boundary engineering illustrates how properties of a material may be altered or tailored by controhg the actual types of interfaces present in a conventiond polycrystalline material. The other method of grain boundary engineering involves increasing the overall interfacial volume fiaction through the reduction in grain size. In the following sections, the development of nanocrystalline materials as a distinct form of grain boundary engineering will be discussed.
3.2
Nanocrystalline Materials
3.2.1
Introduction to Nanocrystalline Materiaïs
In 198 1, Gleiter initially proposed the synthesis of nanocrystalline materials containing large volume fiactions of interfaces[20]. An extensive review of nanocrystalline matenals was published in 1989 by Gleiter, which outlines some of the properties displayed by nanocrystalline materials as a consequeme of their ultra-fine
stmcture[l3]. The increase in overall quantity of grain boundaries (intercystalline volume fiaction) in nanocrystalline materials can affect the physical, mechanical, and chemical properties. There are several emerging applications for nanocrystallime materials including: corrosion and Wear resistant coatings[21], catalysts[22], in-situ repair
of failed s t e m generator tubing in the nuclear industry[23,24], transformer cores[25,26], DC converters[27], recording heads[26,28], magnetic sensors[26], giant-magnetoresistive layers[29], low temperature sintenng[30], optical[29] and computing devices[29].
3.3 Structure of Nanocrystalline Matenals
Nanocrystalline materials represent a novel class of materials with grain sizes less
than l O O m resulting in significant inmeases in grain boundary and triple junction volume fiactions[13]. In three dimensions, grains have been represented by cubes or more accurately by fourteen sided polygons (tetrakaidecahedra) as s h o w in Figure 3.1. As previously discussed, grain boundaries can be characterized according to their
geometry and it has been shown that they can affect the physical, mechanical, chemical, electrical and magnetic properties associated with the bulk material.
Figure 3.1 :
Tetrakaidecahedra representing grain shape for equiaxed matenals.
Over the past decade, the atornic structure of nanocrystalline matenals has been revealed by several methods including: x-ray difEaction[311, extended x-ray absorption
fine stnicture (EXAFS)[32,33], transmission electron microscopy (TEM)[34,35], high
resolution transmission electron microscopy (HRTEM)[36-381, Mossbauer spectroscopy[39], positron annihilation[40], and neutron difiactionr411. The microstructure consists of well ordered crystallites with iarger fiee volumes
or les-ordered areas associated with the grain boundaries. Figure 3.2 is a schematic of nanocrystalline material showing a large fraction of atoms in the grain boundaries and triple junctions[34]. It should be noted that the boundary atoms in white are show in lattice positions but would relax to more energetidly favourable positions.
Figure 3.2:
Schematic diagram of a nanocrystalline materiai[34].
Many of the studies showed strong agreement with the notion that grain boundaries in nanocrystalline materials are not unlike those in polycrystalline materials.
Notably, the grain boundary width in nanocrystalline materials was observed to be approximately lm which corresponds well with what had been observed previously in polycrystaliine materials[361.
The primary differencebetween polycrystalline materials and nanocrystalline materials is the increase in the volume fractions of grain boundaries and triple junctions in the later. This interaystalline volume hction has been quantified by Palumbo et al. using a three dimensional treatment involving tetrakaidecahedral grains[42]. These
calculations have shown that the intercrystallinevolume fraction increases fiom a value of 0.3% at lOOOnrn to more than 50%at grain sizes less than Snm, assuming a grain
boundary thickness of Inm. Figure 3.3 shows the behaviour of the volume fiaction of the intercrystaiiine components as a function of grain size. Furthemore, it has been
observed that the tiple junction volume fraction displays a greater size dependence than the grain boundary volume fraction such that they become equivalent at a grain size of approximately 2m[42].
Grain Size (nm)
Figure 3.3:
Intercrystalline volume fraction as a function of grain sue[42].
The grain boundaries in nanocrystaüine materials have been shown to be similar
in structure to their polycrystalline counterparts. However, very recently, some researchers have proposed that certain high-energy grain boundaries in simulated polycqstalline and nanocrystalline siiicon can exist in the arnorphous state and are, in fact, thennodynamically prefened[43-45]. It should be noted, however, that even though these simulations yielded thermodynamicdly stable arnorphous boundaries in silicon, they have yet to be observed by electron microscopy.
3.4 Synthesis of Nanomaterials
3.4.1
Introduction to Nanoprocessing
German researchers were the first to make a concerted effort to produce materials
with a grain size in the nanometer range in the early 1980s[20]. These materials with grain sizes typically less than lOOm were synthesized by the gas condensation method. Since this t h e a plethora of other synthesis techniques have been developed. Over the last fifteen years inert gas condensation[13,20,46-481,mechanical alloying[49-541, amorphous precursor processing[12,55-741, electrochemical deposition[75-841, spray conversion[85], themochemical processing[86], plasma processing[87], borohydride reduction[88], laser deposition[89], sputter deposition[90-921, electron bearn vapour deposition[93], rapid solidification[94] and sol-gel processing[95,96] techniques have been successfully used to produce nanocrystalline materiais. The final form of the matenais produced by different methods can Vary considerably in terms of structure and properties. In the following, four of the most popular
fabrication methods will be discussed. The gas condensation and mechanical dloybg or
ball miliing methods are common methods for producing nanocrystalline matenals or matenals with ultra-fine particles. However, these methods require at least one consolidation step to yield the nanocrystalline materials in their final form. In contrast, the morphous precursor and electrodeposition methods yield nanocrystalline materiais directly. An effort will be made in the following subsections to examine key differences in the final matenal produced by each of the four methods dimssed.
3.4.2
Nanoprocessing Techniques Requiring Consolidation
3.4.2.1 hert Gas Condensatiou
Nanomatends were first produced in 1981 by Gleiter using the inert gas evaporation method[20]. Figure 3.4 illustrates the typical schematic for a gas-condensing chamber for the synthesis of nanocrystalline materials[l3]. These materials are produced in an ultra high vacuum chamber, which is typically evacuated to l0dPa and then back-
füled with an inert gas such as helium or argon to a pressure of lKPa Then the matenal
to be processeci is evaporated ushg Joule heating in a metal refiactory boat, typically made of tungsten. A 'coldfinger' or substrate chilled intemally with liquid nitrogen or
liquid helium is introduced into the chamber. Atoms of the evaporated metal gas molecules of interest collides with the inert gas causing a loss of energy and the nucleation of metal clusters from the gas phase ocuirs. The evaporated metal
subsequently condenses on the 'coldfinger' which provides maximum nucleation and
minimum growth of grains. Periodicdy the condensate on the chilleci substrate is mechanically stripped. The resultant ultrafine particles are consolidated in a press at pressures between one and five
GPa. These consolidated materials typically take the fonn of a penny shaped disk a few miIlimeters thick with diameters close to one centimeter. The processing parameters that control the final particle size of the nanomaterial are evaporation rate and inert gas pressure. It should be noted that the density of
materials produced by this method is severely reduced when compared to their fuliy dense polycrystailine counterparts. The density is îypically 90% but may drop to as low
as 70% of the polycrystalline material.
Figure 3.4:
Schematic representation of a gas-condensation chamber for the synthesis
of nanocrystalline materi&[ 131.
3.4.2.2 Mechanical Alloying
Initiai materials for mechanical alioying or high-energy ball milhg are usually
coarse grained powders which are heady deformed to produce nanomatenals. These materials are synthesized in a bd-miil using various müling schedules to tailor properties
of the resultant nanomaterial which is later compacted. The initial deformation simultaneously reduces grain size through mechanical attrition while nucleating nanocrystalline grains at shear bands in the deformed material. Longer milling times typically lead to an extremely fine structure or even the formation of amorphous materials. initially, materials with a face-centred cubic structure were not easily milled into ultrafuie particles as they are sofl and tend to conglornerate or sinter into larger particles over time. This obstacle has been overcome by miiling in a hydrogen atmosphere[50]. The hydrogen is believed to embrittle the material and hinder plastic flow allowing effective milling of face centred cubic materials. As a result of hydrogen atmosphere ball
Mlling, Al, Cu, Ni, Pd, Rh and Ir have been nanoprocesseci. Materials with body centered cubic structures such as W, Nb and Cr as weU as materials with hexagonal structures such as Co, Ru, Hf and Zr are effectively milled in inert atmospheres such as
neon or argon. It is believed that these materials are able to store high amounts of energy and damage allowing them to be ball-milied in an ineri atmosphere. It should be noted that nanomateriais produced by hi&-energy bal1 rnilling also
require a post-processing consolidation step sVnilar to the steps taken in the production of nanomatenals by inert gas condensation.
3.4.3
Nanoprocessing Techniques Yielding Nanocrystalline Materials Directly
3.4.3.1 Arnorphous Precursors
This method of nanoprocessing is initiated by producing a rnatenal with an amorphous structure. The amorphous materials are typically produced fiom a melt by meit-spinning, rapid quenching or wire drawing. In some cases amorphous materials cm be obtained by electrodeposition[e.g. 341. Regardless of the initial processing route, the amorphous material is subsequently annealed at temperatures that induce ciystallization.
This anneaiing treatment yields a stnicture with nanostmctured grains, which Vary in size depending on t h e and temperature of the heat treatrnent. Nanocrystalline materials produced by annealing of amorphous precursors c m
lead to the formation of nanocrystalline grains surrounded by a matnx of amorphous material. This structure is observed most ofien in specialty soft magnetic alloys such as Fe-Cu-Nb-Si-B called ~inemets~~[55]. The discovery of ~ i n e r n etype t ~ ailoys which
are ferromagnetic metallic glasses that exhibit excellent magnetic properties has driven
this method into the spotlight over the last decade[55]. These iron-based alloys, typically incorporate Cu, Nb Si and B into the d o y for very specific reasons. The dual phase formed upon annealhg is an Fe-rich phase surrounded by a Cu-Nbnch phase. Grain size of the Fe-rich phase is limited during the a ~ e aby l the fact that the Cu-Nb nch phase has
a higher crystallization temperature that the Fe-nch phase. Furthemore, as Cu segregates to the Fe-rich boundaries, nucleation is enhancerl.
.
Electrodeposition has successfullyyielded many nanomaterials over the last 15 years and is the synthesis route for materials studied in this thesis. Both, direct
current[75,76] and pulsed cument[77] plating methods have been used to yield materials
with nanometre grain sizes. A schematic diagram of a basic experimental setup for producing electrodeposited
nanocrystais is shown in Figure 3.5. Electrodeposition involves a bath containhg a
solvent, bath additions as well as the ionic species to be electrodeposited.
Power Supply/
Figure 3.5:
Heating plate/
Schematic diagram of electrodepositionapparatus.
The main features of this method are the bath, electrodes, and a power-supply
which can provide either direct current or pulsed current with various waveforms as shown in Figure 3.6. There are usually two electrodes, an anode which may be soluble or dirnensionally stable, and a cathodic substrate ont0 which the material of interest is
electrodeposited. By varying, bath composition, pH, temperature and current characteristics such
as current density and the duty cycle in pulsed deposition, materials with varying grain sizes and compositions can be produced. It should be noted that these materials usually do not require post-processing and may be plated in varying thickness and shapes.
Figure 3.6:
Schematic representation of square-wave modulateci current [97].
3.4.4
Cornparison of Nanoprocessing Techniques
The main dierence between the materials requiring post processing consolidation and those that are produced diiectly is the porosity remaining in the consolidated materiais. It has been shown elsewhere that srnail changes in porosity c m lead to large differences in materiais properties such as Young's modulusp8].
The gas evaporation method readily produces pure materials, alloys, ceramics, and composites in high purity. This method of nanoprocessing is very capital intensive and initially had a smaii yidd of a few kg/day. It has also been noted that this material
can have a final density of only 70% after compaction, making porosity a severe issue in materiais produced by this method. In addition, this material is prone to irnpurity pickup or oxidation when exposed to regular atmosphere (as a result of porosity). Hïgh-energy bal1 miiling provides a low-cost alternative to the gas evaporation method and yields many kg/day of nanomaterials. However, in addition to porosity in the nanocrystaliine compacts, there exist many irnpurities fiom the milling balls[99]. The amorphous precursor method typically results in very thin films or ribbons with thicknesses in the micrometer range. Therefore these materials may need to be stacked or layered to achieve the desired structure. Although these materials are fùliy dense, in some cases, residual amorphous gmin boundary regions may alter properties. Aside from being a less expensive route for nanoprocessing, electrodeposition is very versatile, in that complex shapes in various sizes, inctuding: thin films, thick films, fieestanding foil, sheet produas and thick electrofoms, may be produced.
3.5
Properties of Nanocrystaiiine Materiais
3.5.1
Physical Properties
It should be noted that many of the eariy sindies on nanocrystalline materiais were conducted on materials produced by gas-condensation and many of these results disagree with more recent measurements on nanomaterials produced by methods not requiring consolidation steps.
3.5.1.1 Density
Gleiter first stated in a review that the density of nanomaterials produced by the gas evaporation rnethod varies nom 75% to 90% as compared to fully dense
polycrystalline materials[l3]. This variation in density could be controlled to some extent by varying the composition of the materiai being consolidated and the subsequent
annealhg treatrnents. In fact, the theoretical 100% density of polycrystalline matends could be achieved with appropriate a ~ e a l i n g(sinte~g)treatments leading to lirnited
grain growth. Gleiter stated that the decrease in density was a direct result of reduced atornic density at the grain boundaries[l3]. There was considerable variation in reportai results regarding porosity levels of these gas condensed nanomaterials[ 131. Direct metallographic examination revealed levels of closed volume porosity fiom 1-8%[13]. Techniques such as gas permeation, BET rneasurements and positron annihilation were used to study the densification behaviour of nanocrystals. In one study conventional titania (TiO2) with a mean particle size of 1-3pmwas observed to sinter
only at temperatures higher than 1100°C. In contrast, nanocrystalhe titania was obsewed to sinter or densify at temperatures as low as 600°C with grain sizes stable up t o 1000°C. Pore drag was the mechanism proposed for the enhanced thermal stability with
regards to grain growth retardation[l3]. Nanocrystalhe materials that are fabricated in bulk fom directly do not require additional consolidation leadhg to large observed differences in properties including density. In a study on nanocrystalline nickel produced by electrodeposition, the density
is reporteci to be close to 100% of the density observed in conventional polycrystaUine nicke1[100].
3.5.1.2 Hardness
It is well known that the only way to simultaneously increase the strength and toughness of engineering steels is by decreasing their grain size[lOl]. The Hall-Petch relation for yield stress (a,) as a fûnction of grain sire (d) is well established for micrometer and larger grain s
i metals and ceramics[102,103]:
q = o, + kd-"
(eqn 3.1)
where 0. is the stress of fiction and k is a constant representing the stress required to
move dislocations from yielded grains into adjacent grains. Therefore, it is expected that the grain size reduction in nanocrystalline materials wiil have a profound effect on the mechanical properties of hardness and Wear resistance.
In gas condensed nanocrystalline titania, the hardness of Ti02 was found to be greater than its polycrystailine counterparts. The authors amibuted this increase in hardness to decreased particle size as a result of rapid sintering times and noted that the sintering occu~edat 400 to 600°C lower than with conventional particles in the micrometer region[l04]. For nanocrystallirne Cu and Pd produced by inert gas condensation with grain sizes from 3 - SOnm, hardness values increased by factors ranging between two and five[ 1051.
The density of the nanocrystallinemet& was between 72 and 97% of its coarse grained
polycrystalline counterpart. The authors postulate that the ultrafine grain size decreases both the generation of dislocations and their mobility and indicate that residuai stress may be a factor in the increased hardness observed[l OS]. This increme in hardness followed
regular HaU-Petch behaviour, which is in agreement with other research on ball-miiied and electrodeposited nanomaterials[ 106,1071 At extremely smaIi grain sues, some materials have shown a deviation fiom Hall-
Petch behaviour and the material sofiens somewhat which is in disagreement with conventional Hall-Petch behaviour. Figure 3.7 shows this inverse W-Petch behaviour for both gas condensai Pd, Cu and electrodeposited Ni-P[lOS]. Simiiar obsenrations have been made for electrodeposited Ni[109]. Grain boundary sliding[ 11O], diffusional creep[l 101 and increased triple junction volume fiaction[l 1 l] are mechanisms suggested to account for the deviation from regular Hall-Petch behaviour in extremely small grained sarnpies.
Figure 3.7:
HallPetch plot for nanocrystahe gas-condenseci Pd, Cu and electrodeposited Ni-P[IO8].
Despite these deviations fiom Hall-Petch behaviour, the hardness of fully dense nanocrystauine Ni electrodepositsincreases by a factor of five for a grain sue of lOnm as compared to a 100pm polycrystdline sample[l09].
3.5.2
Chemical Properties
3.5.2.1 Corrosion
Under certain electrochemical conditions, localized corrosion in polycrystalline materials is enhanced at grain boundaries and triple junctions[112]. As a result of the increased intercrystalline volume fraction associated with an extremely refined grain size, it seems likely that corrosion resistance in nanocrystalline materials will be reduced. One of the first studies on the corrosion behaviour of nanocrystalline materials was completed on melt-spun arnorphous alioys that were crystallized to form
nanostructured Metglass alloys[ll3]. For Metglass Alloy 2826 (40Fe-40Ni- 16P-4B) there was little difference between the polarization curves of the nanocrystaliine and amorphous alloys. However in the case of Metglass alloy 2826A which contains 15Cr with a reduction to 25Fe the passivation current density as well as the active current density were lower for the nanocrystaliine alloy[l13]. Hashimoto et ai. [114] completed studies on arnorphous pre-cursor Fe4 OCr- l3P7C which led them to conclude that the nanocrystalline alloy was much less stable than the amorphous d o y of the same composition. Upon examination of the corroded surface of the nanocrystaliine samples, it was observai that preferential general corrosion occurred at crystal defects[ 1 141.
Further studies on the anodic polarization behaviour of FqzNi36Cri4PizBsin acidic chloride-sulphate media showed that the nanocrystalline alloy displayed corrosion resistance superior to the amorphous d o y of the same composition[i 151. This was attributed to increased Cr mobility leading to a thicker passivation layer in the
nanocystaliine alloy[ll5]. The potentiodynamic studies of FenSi1oB&~3 and Fe&loBIs alloys produced in the same manner yielded slightly diierent results. The Fe&i&i&ri3 alloy in the nanocrystalline form showed a lower passivation current density than the amorphous doy[ll6]. However the FenSi1&5 alloy did not passivate
at any current which was attributed to the composition of the alloy which contained no chromium[l16]. The potentiodynamic anodic polarization behaviour was studied for sputterdeposited nanocsystalline 304 stainless steel[] 171. The lodized corrosion in this stainless steel film was greatly reduced for the nanostnictured materid. The authon attributed this reduction in l o d u e d corrosion to the enhanced distribution of chforide ions at the metal surface as a result of increased defect density in the passive layer[i 171.
In addition the passive film was more stable than the passive films on regular 304-type
stainless steel[l17]. The anodic polarization behaviour of nanocrystalline nickel electrodeposits has recently been studied by potentiodynamic and potentiostatic testing in a 2M HzS04 solution[1181. Nanocrystalline nickel exhibited the same active-passive-transpassive behaviour as observed in normal aystalline nickel. However, the current density in the
passive range for nanocrystalline matenals is higher than for conventional nickel. This is believed to be due to the increased defect concentration in the passive layer on the nanocrystaline material. In contrast to conventional nickel, however, the nanocrystalhe
material did not show localized corrosion at the gain boundaies and triple junctions.
Electrodeposited nanocrystalline nickel has also been subjected to industrial sait spray resistance tests[80]. It was s h o w that both the nanocrystalline and polycrystalline
nickel coating provided the same corrosion protection for the steel substrate[80].
In general, the overall corrosion rate seems to be slightly enhanced for most nanocrystalline materials although the corrosion is less localized and more uniform.
3.5.3
Electrical Properties
When a fiee electron in a metal meets defects such as a density fluctuation, impurities, vacancies, dislocation lines, interstitial atoms, thermal vibrations or grain boundaries it is scattered and does not move freely[119,125]. Grain boundary scattering occurs as reflections at the boundary or transmission with an associateci Ioss of energy. In either case the essential hypothesis is the same: ifthere is an increase in grain boundary width or volume hction then the electron scattering effect becomes more prevdent . Considerable increases in the overall electî-ical resistivity of nanocrystalline
materials have been observed. Examples of this include nanocrystalline Fe, Cu and Pd produced by the inert gas condensation technique[l3], evaporated Fe-Co fiIms[l21], sputtered Pt filmsjl221, and electrodeposited Ni[123,124], Ni-Fe[125], and Co[K25]. Gleiter[l3] States that metallic impurity content and porosity cannot account for such changes in resistivity. Figure 3.10 shows that the room temperature resistivity of electrodeposited nanocrystatline Pd with a grain size of lOnm is greater than that of
polycrystahe Pd by a factor of ten[l3]. At lower temperatures this factor increases to
approximately hventy five tirnes. To examine the effects of increased intercrystallme
volume fraction in nanocrystalline Ni, Figure 3.1 1 shows the excess resistivity as a function of interphase volume fniction. The excess resistivity is defined here as the total resistivity minus the resistivity of conventional poiycrystalline Nifor which the intercrysdiine volume fraction is negligible[l24].
Figure 3.8:
Resistivities of gas condensed Pd as a function of temperature and grain size[l3].
Not d grain boundaries have the same electron scattering effect on the overail
electrical resistivity in the metal. In recent studies, using bi-crystal samples with stnlcturally weli characterized interfaces, it has been s h o w that the electrical resistivity
is even dependent on the type of grain boundary[126,127]. Certain low X boundaries
such as the twin boundary were observed to have lower resistivity than general high Z boundaries. The authors explained this behaviour in terms of electron scattering at the core of grain boundary dislocations.
O
5
10
15
20
25
30
Volume Fraction (%)
Figure 3.9:
Excess Resistivity as a function of intercrystalline volume fiaction in electrodepositedNi [1241.
Figure 3.10: Temperature coefficient of resistivity vs. grain site in gas condensed Pd[13].
In addition, as the grain size of a metal approaches or becomes smaller than the electron mean fiee path, the temperature coefficient of resistivity is expected to decrease.
Experimental evidence for this effect has been presenteâ for Pd[13] in Figure 3.12 and for Pt[122], Ni[l23,124], Ni-Fe[125] and Co[125].
3.5.4
Magnetic Properties of Nanocrystalline Materials
3.5.4.1 Coercivity
It should be reiterated at this point that this section is a literature review covering the progress in the understanding of the magnetic properties of nanocrystalline materials
up until the measurements for this thesis were initiated in 1992. The scientSc understanding has changed considerably throughout the course of t his thesis and as such the literature review has been stnictured to reflect this. The magnetic properties of polyaystalline ferromagnetic materials are strongly dependent on parameters such as composition, crystallographic texture, intemal stress,
grain shape anisotropy, and grain size distribution[128- 13O]. For nanocrystalline materials in which the grain sine approaches the dimensions of the domain wail thickness or the ferromagnetic exchange length in conventional materials, considerable changes in the magnetic behaviour have been obsewed[l3 1-1411. A nanocrystdine system which has generated world-wide research interest is
devitrified amorphous F~.5Sii3.5BsNl&utype dloys where it has been shown that grain size has a strong effect on the soft magnetic behaviour [l2,59,13 11. These nanocrystalline alloys have been given special consideration for sofi magnetic applications such as transformer cores, recording heads and saturation reactors. In these
materials, the grain size of a-Fe-Si grains is controlled in the nanocrystalline range by Cu and Nb additions which act as nucleation agent and grain growth inhibitors respectively[ 12,601. Figure 3.13 shows the behaviour of coercivity as a function of grain size for alloys of this type, as compared to other soft magnetic materials[l3 11. Low
coercivitiesare desired for soft magnetic materials and in the past, materials with large
grain sizes were produced to achieve this. However, in grain sues in the nanometer range, the coercivity is also observed to decrease dramatically with decreasing grain size
with a much stronger grain size dependence than in the polycrystalline region. Best soft magnetic properties were noted in alloys of this type with grain sizes smaller than 20m. This behaviour was explained in terms of the random anisotropy model[l3 11. To elaborate, as the grain size becornes smaller than the ferromagnetic exchange length,
local magneto-crystailineanisotropies are averaged out which gives similar coercivities
Figure 3.11: Coercivity vs. grain size for various soft magnetic alloys[l3 11.
Strong effects of particle size on coercivity have also been observed in other studies[139,140]. For exarnple, for ultrafine particles (1 0-50nm)of NI, Co and Fe, Gong et d.[1391 reported maxima in coercivity which according to their calculations, correspond to a critical particle size at which the transition to single domain occurs. The coercivity was observed to decrease at grain sizes below the critical grain size for single domain formation. This was explained to be the result of the ferromagnetic to superparamagnetictransition. Gangopadhyay et al.[140] also observed a transition to superparamagneticbehaviour for ultrafine Fe and Ni particles in the 5 to 20 nm range.
3.5.4.2 Saturation Magnetizntion
Strong effects of grain size and particle size on saturation magnetization and coercivity have been observed in the study of nanocrystalhe materials and ultrafine particies. The saturation magnetization of nanocrystailine materials and ultrafine particles is found to be reduced as compared to polycrystalline rnatenals[l3,139- 1411. Gleiter fint reported a 40% decrease in saturation magnetization for 6nm i o n as compared to conventional a-iron[l3]. This decrease was attnbuted to differences in the magnetic microstructure between nanocrystalline and conventional polyciystalline iron. In the case of ultrafine particles (1 0-50nm)of Ni, Co and Fe prepared by inert gas condensation, Gong et ai.[139] observed a rapid decrease in saturation magnetktion with decreasing grain size. This decrease was attnbuted to antiferromagnetic oxide iayers on the ultrafine metal particles. In another study on ultrafine matenals it was
found that as particle diameter decreased, the nomalized rnagnetization ratio decreased[l40]. The reduction in rnagnetization was l i e d to surface effects (oxide layers) which are increased as a result of the smail particle size. Schaefer et ai.[14 11 noted a decrease in Ms in nanocystalhe Ni produced by inert gas condensation, which they explaineci in ternis of the stmctural disorder of,theintenaces. The atoms located in the interfaces were calculated to posses only haifthe magnetic moment of the atoms located within the grahs[l41].
3.5.4.3 Magnetic Microstructure
The magnetic microstructure of nanocrystalhe materials is suggested to dSer fiom that of conventional materials[13,132- 1351. Gleiter[ 131 proposed that every crystallite in nanocrystaliîne iron represents a single ferromagnetic domain. Curent theory states that grains larger than a certain size exhibit multi-domain structures, whereas grains, which are smaller than this critid size result in a single domain
structure. Early studies using the Bitter Technique, Lorentz electron microscopy and Kerr micoscopy have not reveaied domain structures in nanocrystalliie materials[lî]. As a result of this, much research is being conducted in the ara of determinhg
the properties of verifiable sub-micron single domain grains or particles. The critical size for single magnetic domains has been estimated by Kittel as eariy as 1946 who stated that thin films, wires and pariides should have single domains[1361. Calculations of critical particle size bas4 o n energy considerations, for nickel for instance, estimate the formation of single domains below 40nm grain size[137].
Wagner et al. studied the magnetic rnicrostmcture of 7 m nanocrystalline iron by
small angle neutron scatterhg[l38]. They proposed that the materid consists of ferromagnetic grains separateci by a non-magnetic or weakly magnetic interface component with a density of about 40% of that of the crystallites.
3.5.4.4 Magnetic Transitions
Magnetic phase transitions are also a c u m t topic of interest. A sîudy on
nanocrystalline Er in the 10-70nmrange has been carrieci and it was noted that three magnetic phase transitions of erbium norrnally present have vanished at these grain sizes and a new superparamagnetictransition occurs at low temperatures[ 1341. For larger particle sized nanocrystalline Er the three transitions occur but at different temperatures in addition to the new transition. The Neel temperature (transition fiom paramagnetic to antiferromagnetic) has been studied in detail in the FeF system by Mossbauer spectroscopy[l3]. In cornparison
to the narrow (2K)transition temperature range that occurs in sigle crystais, the nanocrystalline FeF has a wide range of transition temperatures. Gleiter reporteci a reduction in the Curie temperature of 40°C in iron with a grain size of 6nm[l3]. In a similar study completed on nanocrystalline Ni, Schaefer et al.[l411 States that the apparent lowering of the Curie temperature is due to reduced Curie temperature of the grain boundary regions.
'
3.6
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AM.El-Sherik, U. Erb, G. Palumbo and KT.Aust, ScriptaMetall. Mater., 27, 1185, (1992).
[Il01 A.H. Choski, A. Rosen, J. Karch and H. Gleiter, Scripta Metail. Mater., 23, 1679, (1989). [ l 111 G. Palumbo, S.J. Thorpe and KT.Aust, 24, 1347, (1990). [Il21 M.G. Fontana and N.D. Greene, Corrosion Engineering, Mc-Graw Hill, Inc.,
(1967). [2 133
R.B.Diegle and J.E. Slater, Corrosion, 32, 155, (1976).
[114] K. Hashirnoto, K.Osada, T. Masumoto and S. Shimodaria, Corrosion Sci., 16, 71, (1976). [Ils] S.J. Thorpe, B. Rarnaswarni and KT.Aust, J. Electrochem. Soc., 135,2162, (1988). [116] P. Bragagnola, Y. Waseda, G. Palumbo and KT.Aust, Corrosion/Coating of Adv.
Mat., Mat. Res. Soc. Int'l. on Adv. Mats., 4,469, (1989). [117] R.B. Inturi, 2.Szklarska-Smralowska, Corrosion, 48,398, (1992). [Il81 R. Rofagha, R. Langer, A.M. El-Sherik, U. Erb, G. Palumbo, and K.T.Aust, Scripta Metall. Mater., 25, 2867, (1991). [119] C. Kittel, Introduction to Solid State Physics,'7 ed., John Wiley & Sons, Toronto, (1996). [120] M. Springford, Electrons at the Fermi Surface, Cambridge University Press,
(1 980).
[121] J.P. Riviere, P. Bouillaud, J.F. Dinhut, and J. Delafond, Thin Solid Films, 176, L161,(1989). [122]
X.Mei, M. Tao, H. Tan, Y. Han, and W. Tao, Mat. Res. Symp. Proc., 286, 179, (1993).
[123] I. Bakonyi, E. Toth-Kadar, T.Tarnoczi, L.K.Varga, A. Cziraki, 1. Gerocs, and B.
ForIwaw,
-
M.J. Aus, B. Szpunar, U. Erb, AM. El-Sherik, G. Palumbo, and KT.Aust, J. Appl. Phys., 75,3632, (1994).
J. McCrea, MA.Sc thesis, University of Toronto, Canada, in progress, (1998). 1. Nakamichi and T. Kino, Proc. Of JIMIS 4, Suppl. Tram Jap. Inst. Met., 1013,
(1986).
I. Nakamichi, J. Sci. Hiroshima Univ., Ser 4 54,49, (1990). E.C. Stoner, Magnetism and Matter, Methun and Co. Ltd., London, (1934).
A.E. Berkowitz, E. Kneller, Magnetism and Metallurq, Academic Press, New York, (1969).
RM. Bozorth, Ferromametisîq IEEEPress, New York, (1 978). G. Herzer, Mat. Sci. Eng., A133, 1, (1991).
U. Admon, M.P.Dariel ,E. Gninbaum, and J.C. Lodder, J. Appl. Phys., 62, (S), (1987).
J.F. Smyth,S. Schultz, D.Kem, H. Schmid, and D. Yee, J. Appl. Phys., 63, (81, (1988).
A Cowen, B.Stolmiam, R.A. Averback, and H. Hahn,J. Appl. Phys., 61,33 17, (1987).
E.E.Anderson, S. Arajs, and N. Amin, J. Appl. Phys., 69, (8),
(199 1).
C. Kittel, Phys. Rev., 70, 965, (1946).
D. files, Introduction to Magnetism and Mametic Materials, Chapman and Hall, New York, 1991.
W.Wagner, A Wiedenmann, W. Petry, A Geibel and H. Gleiter, J. Mat. Res., 6, 2305, (1991).
W.Gong, H. Li, 2.Zhao and J. Chen, J. Appl. Phys., 69, (1 99 1). S. Gangopadhyay, G.C. Hadjipanaysis, B. Dale, CM.Sorenson and K.J. Klabunde, Nanostr. Mat., 1, 77, (1992).
H.E.Shaefer, H.Kisker, H. Kronmuller and R. Wunchum, Nanostr. Mat., 1,523, (1992).
Chapter Four DEVELOPMENT OF APPARATUS
4.1
Measurement of Magnetic Properties using a Hybrid VSM
The dipole moment induced in a sample placed in a uniform magnetic field is proportional to the product of the applied field and the sample susceptibility. If the sample is vibrating in a uniform magnetic field, an electrical signal proportional to the
magnetic induction will be induced in a pair of suitabiy placeû coils as a result of the magnetic flw changes near the sample. The generated electncal signal will be an AC signal with a frequency determined by sample oscillation. The vibrating sample magnetometer, first described by Foner [l] makes use of these properties by mechanically osciilating (vibrating) a sample in a uniform magnetic field with a @id rod comected to a mechanical transducer assembly as s h o w in Figure 4.1. The mechanical transducer may take the forrn of a synchronous motor assernbly or
even a loudspeaker. In modem systems, (such as the Lakeshore Mode1 7300 [2]),a vibrating diierential capacitor (one plate fixed) is used to provide a reference for amplitude and frequency. The ac signal induced in the capacitor plates (as shown in Figure 4 . 9 provide a control signal for modulating the transducer. The pickup coils are located adjacent to the space the sample vibrates in and are centred in the gap to provide
for maximum field homogeneity.
.
Mechanical OsciIlator
1
Refertnct
Y
Base Plate \
Magntt Pole
\
\
i
Rotstable Housing
-
Rcfercnce Sarnple
Vibrating Rud or Shaft
Sample
Figure 4.1:
Schematic diagram of the vibrating sarnple magnetometer (VSM).
The magnetic field applied by the electromagnet is controlled by a feedback loop incorporating the bi-polar power supply and a Hall sensor located between the pole
pieces. The first magnetic measurements on Ni and NLP were completed on an EG&G Parc 458A vibrating sample magnetometer located at Royal Military College, Kingston, Ontario. During the t h e these measurernents were being taken, the following apparatus was under development.
4.2
Development of the Hybrid Vibrating Sample Magnetometer
The vibrating sample magnetometer (VSM)that was constructed for the purposes of this research is the result of the hybridization of a modern data acquisition package,
mechanical transducer, amplifier/controller with a large iaboratory electromagnet as shown in Figure 4.2.
Sample Cup
Figure 4.2:
\ ~ i c kUp Coiis
Schematic diagram of hybrid the vibrating sample magnetometer.
The basic system consists of an 0.43rnVarian electrornagnet mated with a custom 168 amp, 48 volt bi-polar power supply fkom hverpower Corporation. This
magnetlpower supply configuration is used in conjunction with a Lakeshore mode1 7300 vibrating sample magnetometer and controller. The equipment in this configuration can measure magnetic moments as low as 5 x 104iIm2 and with the 70 mm gap can achieve
an applied field of 1.1 x 106Alm. The entire system is controlled by a powerful data acquisition software r u h g on a high end desktop cornputer. The 70 mm gap must house pickup mils as well as the sample placing limitations on sarnple geometry. Despite this geometrical constraint, the sensitivity of the VSM allows accurate measurement of sample quantitities of less than 1 mm).
In addition, a high temperature micro-fûrnace can be installed between the pole pieces of the electromagnet. This allows measurement of magnetic properties at temperatures between 293K4273K as well as other themo-magnetic measurements such
as the determination of transition temperatures or to induce uniaxiai anisotropies through annealing in a magnetic field. The data analysis and presentation tools allow the display and manipulation of measured data. The applied field is represented on the x-axis, and the y-axis is directly proportionai to the magnetic moment of the sarnple. Figure 4.3 shows an achiai output
fiom the hybrid VSM running the Lakeshore CryotroNcs 730s data acquisition package.
Or I entak ion
Per-pend I cu 1ar rbss
Sam
le
ID
NI&! Test
Uo1ume
0.1 cc
tmum Field
k n t
4.97 EraJ
Figure 4.3:
F~eld
6 2 0 Oer
-0 Oer #%gnet1z a t Ion 5.16
EIP)
Sample output fiom the hybrid VSM for an experiment using
polycrystalline nickel (99.9g0/0 purity).
4.3
Calibration of Hybrid VSM
A cylindrical nickel sample was provided by Ldceshore Cryotronics for
calibration purposes. The mass of the sample was 0.066g and 99.99% pure. The magnetic moment of the sample was measured using two testing methods which provided magnetic moments of 3.64 emu (electromagnetic units) when the matenal was magneticaliy saturated at 400 kA/m. The length-to-diameter ratio of the right cylinder was 1.
The hybrid VSM controiler was caübrated for both the centraiized position of the
sample within the pickup coii assembly and the output of the magnetic moment for the
standard nickel sample. Samples measured within the body of this work were 3mm disks punched fkom electrodeposited sheet and stacked to form right cylinders with a length to diameter ratio of 1, as s h o w in Figure 4.4.
Figure 4.4:
Schematic diagram of sarnple configuration between pole pieces of electromagnet and induced moment pickup coils.
4.4
Advantages and Disadvantages of Technique:
The vibrating sampIe magnetometer is one of the most versatile magnetic measurement systems available today. In tems of measurernent capabilities, full magnetization curves (Mvs. H)for almost any material may be measured which yield: saturation magnetization, coercivity, initial permeability, and remnant magnetization. Anisotropy measurements can be conducted as the whole assembly (rod and motor) has
the provision for 360' rotation relative to the base plate (see Figure 4.1). The vibrating sample magnetometer cm rneasure full magnetization vs. temperature ranging fiorn 12733 to 4.OK with the addition of a high temperature oven and appropriate sample holders and/or the addition of a iiquid He cryostat. With these
.
measurements the magnetic transition temperatures such as the Curie point may be determined. Lastly, magnetization vs. time behaviour cm be obtained which ailows for the determination time dependent phenornena such as flux creep.
In addition to its versatility, its sensitivity permits the vibrating sample magnetometer to measure the magnetic moment of weakiy magnetic materAs. In addition, small sample size requirements in the mg range or lm3range make this a very attractive method for magnetically characterizhg materials which are produced in lirnited quantities. The accuracy of the magnetic moment measurement is better than 2% and provides reproducibility better than 2%. It can also be used in a zero field environment to measure samples with i n t ~ s i magnetic c moments as the VSM only responds to the flux
due to the magnetic moment of the sample.
Although the vibrating simple magnetometer generates magnetization vs. magnetic field data, it is not easiiy suited to the determination of induction W. extemal magnetic field. This is due to the fact that sarnple size is short which leads to large
demagnetizing effects as a result of sarnple geometry. In practice, calibration charts have been calculated for various sample geometries which allow the determination of
induction with well known geometries [3-51.By examining sample orientation and mass or volume with a known density, it is possible to determine a correction factor specific to a certain sample geometq. This is tirne consurning, and to date, there is no software which keeps a fle of the demagnetizing factors for well-known geometries. As such, this system displays M vs. H behaviour as opposed to the direct display of B vs. H. Lastly, the h e d pole piece width limits the sarnple's siie to less than 2.5".
4.5 References
[1]
S. Foner, Rev. Sci. Inst., 30, 548, (1959).
[2]
Lakeshore Ciyogenics, Hardware 7300S,Reference Manual, (1993).
[3]
E.C. Stoner, Phil. Mag.,(7), 36,803,(1945).
[4]
J.A Osborn, Phys. Rev.,67,351, (1945).
[5]
D. Chen, J. A. Bmg and R. Goldfarb, IEEE. Trans. Mag., (4), 27, 3 60 1 ( 199 1).
Chapter Five MAGNETIC PROPERTIES OF
NANOCRYSTALLINE TRANSITION METALS
5.1
Magnetic Properties of Pure Nanocrystalline Ni and Co
5.1.1
Introduction
Early studies on the magnetic properties indicated substantial decreases of
saturation magnetization with decreasing grain sime as indicated in section 3.5.5.2[14].
As saturation magnetization has previously been determined to be dependent only on the magnetic moment and nurnber of atoms per unit volume of a matenal it has long been thought to be unrelated to the microstructure of the material[5,6]. Therefore, as a result of reported lower saturation magnetizations for nanocrystalline materials[l-41 this study was initiated with the aims of determinïng saturation magnetization behaviour for pure
nanocrystalline materials produced by electrodeposition.
5.1.2
Experimental
Nanocrystalline nickel of 99.9% purity was produced by A.M.El-Sherik, Queen's University, Kingston, Canada, using a recently developed electroplating procedure[7]. Specimens were electroplated fiom a modifieci Watt's bath ont0 a Ti substrate to a thickness of approximately 0.3rnm, and subsequently mechanically stnpped fiom the
substrate. The electroplating pararneters were adjusted to produce Ni electrodeposits in two grain siies, 10 and 20nm. As previously shown,this method of electroplating produces pore free samples 171. Some of the 1Onm samples were then annealed at 300°C in an argon atmosphere for 15, 30 and 60 minutes to produce grain sizes of 2Onm, 30nm and 40nm respectively by grain growth. Sarnples with IOnm, 120nm, and 5OOnm grain sizes were obtained in the as-plated condition. Nanocrystalline and polycrystaihe cobalt electrodeposits of 99.9% purity were produced by E. Chung, Queen's University, Kingston, Canada[8]. Specirnens were electroplated fiom a bath containing cobalt sulphate ont0 a Ti substrate to a thickness of approximately O. lmm, a subsequently mechanically stnpped fiom the substrate. The electroplating parameters were adjusted to produce Co electrodeposits in two grain sizes,
15nm and 5pm. The grain size of the as-plated Ni and Co electrodeposits and the annealed Ni
deposits were detemineci directly fiom 250 grah s i z h u n t measurements on dark field transmission electron micrographs. Thin foils for TEM examination were prepared by electropolishing using an electrolyte comprising 6% perchionc acid, 15% methanol and 79% acetic acid at a temperature of-10°C and a voltage of 15 volts DC. Figure 5.1
shows a TEM dark field micrograph of the 1lnm as-plated M. The materiai shows an
equiaxed grain structure and essentially no porosity.
.
Figure 5.1:
TEM dark field micrograph of nanocrystailine nickel having an average grain sue of 1lm.
Upon annealing, the grain size increased unifody without major changes in the
crystallographic texture. For example, Figure 5.2 a,b shows x-ray diffraction scans of an
as-plated sample (grain size 20nm) and an annealed sample with a grain sue of 4Onm as determined by a 250+ grain count fiom a trammission electron micrograph giving an
accuracy of +/- lm. Both x-ray diction scans show a weak (200) texture meaning the ratio of (200) to (1 1 1) peaks is slightly hi*
than for standard powder difiaction
peak ratios (JCPDS- Powder Difhction File) for nickel with a random texture. Figure 5.3 shows a TEM bright field micrograph of the electrodeposited nanocrystalline Co. The rnaterial displays an equiaxed grain structure with an average grain size of 1S m . Figure 5.4 shows a SEM micrograph of the electrodeposited
polycrystalline Co with an average grain size of 5pm.
Figure 5.2:
X-ray difiaction sans showing weak (200)texture, a) as plated matenai (20nrn), b) anneded materiai (40nrn).
Figure 5.3:
Bright field transmission electron micrograph of nanociystalline Co
electrodeposit with 15nm grain size.
Figure 5.4:
Scanning electron micrograph of electrodeposited polycrystalline Co with
an average grain size of 5 pm.
X-ray *dion
scans of the polycrystalline and nanocrystalline cobalt samples
are presented in Figure 5.5. This figure shows that the nanocrystalline cobalt has a strong direction [14]. In the present study the samples were mounted in such a way that
the 4 11> direction was perpendicular to the magnetizing field. Therefore we have 6
easy directions within 20' of the plane of measurement. It has been shown for material with single domain uniaxial anisotropy that as the angle between the easy direction and magnetic field is decreasing the coercivity d i increase significantIy[l5]. it is difficult to use this theory for cubic systems but the sharp increase in coercivity observed in Figure 5.13 at 3nm grain size, coincides well with the (1 11) texture evolution just before the
transition into the amorphous region[l]. In addition, the coercivity may be affectecl by stresses in the matenal. It is weU
known that as the intemal stress increases in nickel the coercivity will increase [Hl. As the phosphorous content increases, the grain size of the nickel phosphorus deposit decreases (see Table 5.3). As previously proposed[lL this increase in phosphorus
concentration iikely strains the nickel lattice, which may lead to an increase in coercivity.
5.2.4
Conclusions
The saturation magnetization of nanocrystaiiine Ni-P electrodeposits was found to decrease with increasing phosphorus content and decreasing grain sue in the deposit. It
is proposed that this decrease is mainly due to the phosphorus content in the deposit and
to a lesser degree to the grain siie. Furthemore the ferromagnetic to paramagnetic transition at approximately 16 at% P was found to coincide with the structural transition fiom nanocrystaüine to amorphous.
The coercivity was observed to be strongly affected by grain sïze/phosphorous content with a maximum occurringjust before the transition nom nanocrystailine to
amorphous. Clearly the interpretation of the mercivity data is cornplex, as texture, grain size, phosphorus content and intemal stresses are ail contnbuting effects. The crystallographic texture, intemal stresses and domain structure must be studied in more detail before this behaviour can be understood.
References
G. McMahon and U. Erb, Microstr. Sci., 17,447, (1989).
D. Ostrander, 1993, MSc Thesis, Queen's University, Kingston, Ontario, Canada. P.A. Albert, 2.Kovac, H.R.Lilenthal, T.R.McGuire and Y. Nakamura, J. Appl.
Phys., 38 (3), 1258, (1967). 1. Bakonyi, L.K.Varga, A. Lovas, E. Toth-Kadar, and A. Solyom, J. Mag. And Magn. Mat., 50, 111, (1985). 1. Bakonyi, A Burgstaller, W. Socher, J. Voitlander, E. Toth-Kadar, A. Lovas, H.
Ebert, E. Wachtel, N. Willman and H.H. Liebermm Phys. Rev. B., 47,961, (1993).
G.Palumbo, S.J. Thorpe, and KT.Aust, Scripta Metall., 24, 1347, (1990). M. Hansen, Constitution of Binant Allovs, McGraw W, Toronto (1958). A. Burgstaller, W. Socher, J. Voitlander, 1. Bakonyi, E. Toth-Kadar, and A. Lovas, J. Mag. And Magn. Mat., 109, 117, (1992).
G.McMahon and U.Erb, J. Mat. Sci. Lett., 8,865, (1989). B. Szpunar, R Zugic, U. Erb, and L.J.Lewis,Cm. Met. Quart., 34, 28 1, (1995). B. Szpunar, W.E.Wallace and P. Strange, J. L e s - C o m o n Met., 123,37, (1986). C. Kittel, Phys. Rev., 70, 965, (1946).
L.Neil, C.R (Paris), 224, 1488, (1946). AE. Berkowitz, and E. Kneller,
1,New York, 377,
(1969).
E.C.Stoner and E.P.Wolfarth. Phil. Trans. Roy. Soc., 240, 599, (1948).
5.3 Magnetic Properties of Electrodeposited NaoocrystaIlineNi-Fe 5.3.1
Introduction
The nickel-iron systern contains many technologically relevant alloys such a s Pennailoy, Hymu,Mu-metal and Supermalloy. These ailoys consist mainly of nickel and iron and have been used for soft magnetic applications for many years[l]. These aüoys contain varying amounts of Cu, Mo and Cr in various concentrations depending on the material property to be optimized[2]. One of these is the use of Mo to raise the eiectrical resistivity in transformer core material to reduce eddy current losses. In a similar fashion, Cu is added to these aiioys to irnprove fomabiiïty during fabrication. Permalloy has been used for many years as the standard soft magnetic material for cornparison of other soit magnetic materials. Permalloy is a binary ailoy, which has
an extremely high pemeability as a result of low anisotropy constants and magnetostriction tending to zero in the composition of interest (Ni-20%Fe). Low anisotropy leads to lowering of domain wall energy and near zero magnetostriction reduces the effects of pinning sites ami residuai stresses during domain rotation or domain waii movement[3]. The result of these effects is a very hi&
permeabaty
wupled with a low coercivity making it ideal for soft magnetic applications. Jagelinski has recently summarized the materials requirements for future high performance inductive magnetic recording heads[4]. These are reprodud in Table 5.4,
which also lists the reasons for each of the property requirements. Jagelinski pointed out that conventionaiiy processed Permalloy, Sendust and amorphous alloys currently used in
recording heads will not meet new materials requirements[4].
Table 5.4:
Materials requirements for magnetic recording heads (adapted fiom ref 4)
Property Large saturation magnetization
Reason
I Large gap
Hîgh permeability at ali frequencies
I
High efficiency over wide fiequency range *
Srnail coercivity with low hysteresis loss
Low thermal noise
SrnaIl but non-zero uniavial anisotropy
ControI of domain structure
Low magnetostriction
Low media contact noise
High resistivity
Minimized eddy current losses 1
Wear Resistance
Long life
Corrosion Resistant
Long Me
Good thermal and time stabiiity
Reliability
Low f o d g effêct
Easy and reliable manufachiring process -
As s h o w in d o n 5.1, nanoprocesseci nickel displayed constant saturation
magnetization. Nanoprocessed nickel aiso displays substantial increases in hardness[S], and electricai resistivity[6]. However, through the addition of iron, the saturation magnetization could be incread to levels of commercial importance. At the same t h e , the coercivity should decrease due to the anisotropy and magnetostriction decreases at the Perxnalloy composition[l]. This reasoning fomed the impetus for the development of the nanocqstalline Ni-Fe electrodeposits studied in this section. 5.3.2
Esperirnental
Nanocrystalline samples of nickel-iron doys with Fe contents up to 30%Fe were electrodeposited by C. Cheung, Queen's University, Kingston, Canada[7]. The propnetary solution used in the present study was a modified Watts nickel bath containing additions of iron chloride (source of iron), sodium citrate (complexing agent), and saccharin (stress reliever / grain refinement agent). The cathode matenal was
titanium wMe the anode was electrolytic nickel. Electrolysis was canied out at 50°C, without bath agitation, at a current density of 0.2 /Wcm2 D.C.. Electrodeposits of 1 cm x
lcm x 300 pn were produced and subsequently mechanically stripped off the titanium substrate. The composition of the deposits was determined by energy dispersive X-ray
spectroscopy (EDS)in a conventional scanning electron microscope (SEM). An average of 10 analyses were perforrned on each deposit. Grain sizes were calculated using the Scherrer formula for the (1 11) diffraction line obtaineâ by using a standard 8-20 x-ray diaactometer and compared to the standard line width of a nickel powder sample with a grain size larger than 1 Pm. For some samples, the grain size was also determined directly from transmission electron micrographs. Thin foils for transmission electron Mcroscopy (TEM) examination were prepared usimg jet-polishing in an electrolyte comprising 75% acetic acid, 15% rnethanol, and lû% perchloric acid at a temperature of -10°C and a voltage of 14 W C . Figure 5.14a,b shows the TEM bnght field and dark field micrographs, for a nanocrystalhne 80%Ni-20%Fe electrodeposit.
.
Figure 5.14: TEM micrographs for an electrodeposited nanocrystalline Ni-20wt%Fe doy: a) bright field, b) dark field[7].
Figure 5.15 shows how the grain size decreases with increasing iron content in the
electrodeposit. It has been stated elsewhere[7] that codeposited uon has a profound grain refining effect. The grain size of the pure nickel electrodeposited fiom this plating bath was 2 lm whereas the grain size of the electrodeposit containing 29%Fe reached a
minimum of 11m.
Iron Content (wt%)
Figure 5.15: Effect of iron content on the grain size of electrodeposited nanocrystalline
Ni-Fe doys[7].
The Lakeshore Cryogenics Mode1 7000 hybrid vibrating sarnple magnetometer detded in chapter four was used to generate hysteresis loops of the nanocrystallineNi-Fe electrodepositsat room temperature. For each composition several 3mm disks were stacked on top of each other to obtain samples of a constant weight and length to
diameter ratio.
5.3.3
Results and Discussion
Figure 5.16 shows the saturation magnetization of nanocrystaUine N-Fe electrodeposits for iron contents up to 30%. It can be seen that the saturation
magnetization increases linearly with increasing iron content. This is due to the increased contribution of the iron atoms to the overail magnetic moment observed in each sample. Thus, as for the case of nanocrystalline Ni and Ni-P (sections 5.1 and 5.2
respectively) the grain size of the material has M e effect on saturation magnetization of
Ni-Fe ailoys. The saturation magnetization of the nanocrystdineNi-Fe at the technologically relevant Permalloy (78%Nig22%Fe)composition is 98% of the iiterature vaiue
h
VI
O F C
E
3 C:
-O w
a -e N
C
z
0)
C
O .-
C
2 3
w
Ca
V)
O
5
10
15
20
25
30
Fe Content (wt%) Figure S. 16: Saturation magnetization of nanocrystallineNi-Fe electrodeposits.
35
Figure 5.17 shows the behaviour of coercivity as a function of iron content. It cari
be seen that the coercivity decreases with increasing iron content or decreasing grain size. This is expected as the Ni-Fe alloys in the 20.30%
Fe range exhibit low coercivities as a
result of low or zero magnetostriction and low values of anisotropy constants[l].
5
1O
15
20
25
Fe Content ( w t % )
Figure S. 17: Coercivity of nanocrystalline NoFe electrodeposits.
A good soft magnetic material is said to have a coercivity below 500A/m and it
can be seen that the nanocrystalline NLFe in the range 15%-30% Fe satisfies this demand. It remains to be seen whether thermal treatments could lower this coercivity value further making this a technicaily feasible replacement for existing soft magnetic applications which require high hardness and Wear resistance.
5,3.4
Conclusions
The saturation magnetization of electrodeposited nanocrystalline Ni-Fe hcreases linearly with increasing iron content up to 3OwtOhFe and is relatively independent of
grain size. Nanocrystalline Ni-22wt%Fe displays approximately 97% of the saturation magnetization of comrnercially used 22wt%Permdoy. The coercivity of the electrodeposited nanocrystahe Ni-Fe decreases with increasing iron content and decreasing grain size. It is believed that Iow magnetostriction
wmbined with low values of anisotropy constants, can account for this decrease in wercivity with increasing Fe concentration in these NoFe alloys. Overall, these materials represent a large improvement over nanocrystalline nickel
for sofi magnetic applications. The saturation rnagnetization is increased and the wercivity is decreased for nanoprocessed material with a composition close to conventional permalloy.
5.4.5
References
[l]
D. mes, Introduction to Marnetism and Mametic Materials Chapman and Hd,
New York, (1991). [2]
Metals Handbook Properties and Selection: Non-Ferrous and Special Purpose Mateds, 1orn edition, Volume 2, ASM International, (1 990).
[3]
R.M. Bozorth, Ferromametism, D. Van Nostrand Co., New York, 195 1.
[4]
T.Jagelinski, MRS Bulletin, 36,March 1990.
[5]
A.M. El-Sherik, U.Erb, G. Palumbo and K T . Aust, Scripta Metall. et Mater., 27, 1185, (1992).
[6]
M.J. Aus, B. Szpunar, U. Erb, AM. El-Sherik, G. Palumbo and K.T.Aust, J. Appl. Phys., 75,3632,(1 994).
[7]
C. Cheung, F. Djuanda, U. Erb and G. Palumbo, Nanostr. Mat., (5), 5, 513, (1995).
5.4
Mapetic Properties of Electrodeposited Nanocrystalline Co-F'e Alloys
4 1 Introduction
Iron and cobalt are the only elernents which,when added together, produce a conventional polycrystalline alloy with saturation magnetization higher than that of iron alone. Figure 5.18 shows the synergistic effects on saturation magnetization in ironcobalt alloys. This phenornenon was first observed in 19 12 by Preuss and Weiss[l]. The
composition with the maximum saturation rnagnetization contains about 35% Co and reaches 1.95MA/m (219 emu/g)[2]. This alloy was narned Hiperco and was used in applications requiring extrernely high saturation magnetization. However, iron-cobalt at
this composition is brittle and extremely difficdt to work with. In addition, Hiperco suffers fiom low permeability and is infenor to iron-alloys, nickel iron alloys (Permdloy and Hypernik) as well as nickel iron alloys with Mo and Cu (Supermalloy and Mumetal).
It turns out that the synergistic effect on saturation magnetization is entirely due to the uicrease in moment of the iron[2,3]. The moment of Co remains constant over the entire composition range of Fe-Co deys as the moment of Co is independent of its atomic environment. It should be aoted that the Curie point of Fe is much lower than that of Co even thougb the moment of Fe is higher. This is due to the number of nearest neighbors in the f C.C. and h.c.p. lattice being 12 as opposed to the 8 which ouxir in the b.c.c. lattice
of iron[3].
I
20
1
1
f
40 60 80 Weight Percent Co
IGa
Figure 5.18: Vanation of saturation magnetization with composition in bon-cobalt
alloys after Weiss and Forrer (1929) [Z].
Eimen invented Permendur in 1932, which contains 5Ph iron and 50%Co. This alloy has a saturation magnetization only slightly lower than Hiperco. However, the pronounced increases in permeability after annealhg (850°C or 1000°C) at the 50%FeSO%Co composition can be attributed to the fact that at 4S%Co, the magnetoaystdline
anisotropy coefficient (K) 0s zero[4]. Further improvements on this alloy were made in 1939 when Wahl and White were granteci a patent on Supennendur for an FeCo 49%
alloy which wntained 2% vanadium. The vanadium served to increase both the resistivity and penneability as weii as the workabifity of the previously brittle alloy. This addition of vanadium eliminated many of the more costly and elaborate procedures for
producing the iron-cobalt alloys in their final fonn[5]. Figure 5.19 shows the phase diagrarn of the Co-Fe system produced by Ellis and
Grenier[6]. It is interesting to note that Co-Fe aiioys with Co compositions ranging fiom 30%-70% undergo an order-disorder transformation which has a temperature maximum
near the 50%cobaft-50%ironcomposition[l]. From a magnetic standpoint, this alloy is the most usefil and therefore this transition is of some interest. The disordered structure
is b.c.c.which transforms to f C.C.at the transformation temperature. CENT COBALT SO 60
ATOMIC PER '10
20
30
40
1
70
80
90
I
1
1
1
M ELT
-
CO
1495-
4 7 +'MELT
-
7 (FACE-CENTERED) 1115~ MAGNETIC
910-
.-,-R,,MAGNETIC
,a-c
-'-
770°
(BODY-CE~TERED)
TRANSFORMATION TRANSFORMATION
CG
DJSORDERED
(d)/œ-- ORDERE>--
-
I
?
10
20
30
50 60 PER CENT COBALT
40
Figure 5.19: Phase diagram of Co-Fe alloy[6].
70
Another feature of interest on the phase diagram shown in Figure 5.19 is the magnetic transformation which occurs upon heating fkom a ferromagneticb.c.c. phase to a non-magnetic Ec-c.phase. Figure 5.20 shows the change in induction during the phase transformation in SO%Fe-Co. The transfomation temperature at this composition is 980°C.
This transformation occurs at temperatures lower than the Cune temperature
which makes deterrnination of the "ferromagnetic to paramagnetic" transition temperature difficult. Foner has extrapolated the Bs vs T curves that he called "virtud"
Curie points that existed above the non-magnetic a to y transition temperature[7].
0
Figure 5.20:
200 400 600 800 IO00 1200 TEMPERATURE IN DEGREES CENTIGRADE
Phase transformation in ironcobalt (SO%Co) showing the high temperature Ec.c. phase is non-magnetic[l].
The region of interest pertaining to this report lies in the pure Co to Co-22%Fe
range. At room temperature, for pure cobalt and low iron compositions the phase present is hexagonal. Upon adding iron in quantities greater than 5%-7%, the d o y transforms into the face centered y phase and at approximately 25%Fe the alloy transforms into the body centered cubic structure, a,which is characteristic of iron. In between these single phase areas are 2-phase regions, which were detected by analysis of x-ray difiaction scans on the Co-Fe alloys. Previous sections in this chapter have shown that the saturation rnagnetization of Ni, Co, Ni-P and Ni-Fe is independent of grain s k down to 1On.m. Results presented in section 5.1 have shown that the coercivity of Co was lowered by nanoprocessing. Although only a few of the Co-Fe samples produced for the purposes of this study are nanocrystalline, the previous work (sections 5.1-5.3) forms a basis for hypotliesis. The purpose of the present work is to study the magnetic properties of the Co-Fe with iron concentrations up to 22%. The properties addresseci in this report are as follows: crystal structure, saturation magnetization, and coercivity.
Co-Fe alloys were produced by E. Greenberg and C. Cheung, Queen's University, Kingston, Canada, via electrodeposition using D.C. plating in a bath wntaining cobalt sulphate[8]. Iron was added in 5gR. increments in the fom of iron sulphate. The g was 39'43OC. The pH during plating was kept temperature range used d u ~ plating
between 3.4 and 4.5. The specimens were electroplated ont0 a Ti substrate to a thickness
.
of approlemately 3 0 0 and ~ subsequently mechanically strïpped from the substrate. Optical microscopy, SEM and TEM were perforrned on these samples after stripping and ultrasonic cleaning to determine the morphology and grain size[l].
EDS analysis revealed that CO-depositionof iron and cobalt occurred depending on the arnount of iron sulphate added to the bath and that the samples produced had compositions ranging fiom pure cobalt to cobaltœ22%iron[9].The samples used for magnetic study were al1 plated with a current density of 100A./cm2to maximize mass and thickness[9]. SEM and x-ray dIfltiaction performed by C. Cheung on the Co-Fe electrodeposits revealed decreasing grain sue with increasing iron content[9].
The morphology of this range of samples fds into three distinct regions. For the pure cobalt and alloys with low Fe concentration, grain sizes ranged fiom 1 to 5 microns with a "leaf like" morphology. As the iron content was increased nodules formed, and the grains were refuied to the range of 100 nanometers to 2 microns. The electrodeposits containhg the highest amount of iron displayed a "pyramid like" structure with grain sizes ranping fiom less that 1 micron to 20nm[8]. It should be noted that, in ref.[8], no reasons for these three distinct regions were given. The Lakeshore Cryogenics Mode1 7000 hybrid vibmting sample magnetometer
(VSM)was used to generate hysteresis loops at room temperature. Each electrodeposited coupon was punched into 3mm disks. These disks were stacked on top of each other to form a cylinder for use in the VSM. Hysteresis loops were generated using a sweep time of 20 minutes and a maximum field of 960kNm. At least three hysteresis loops were generated for each sample. The magnetometer was calibrateci using a 99.999% pure Ni
cylinder, aligned in the centre of the pickup coils.
.
5.4.4
Results and Discussion
Three distinct phases exist at room temperature in the compositional range O25Wt%Co as s h o w by the phase diagram by Ellis and Greiner[6]. Co-Fe alloys ranghg
fiom pure Co to Swt%Fe display a hexagonal close packed stmcture at room temperahires. Aiioys ranging fiom Swtohiron to 1lwt%iron exist in the f.c.c. phase whereas a dual phase stmcture exists from 1lwt% iron to 24wt% iron in iron-cobalt alloys. Above 24wt%Fe, a body centered cubic structure with the lattice parameter corresponding to b.c.c. iron is present. It should be noted that the phase diagrams for electrodeposited material are ofken much different than for material produced by different methods. Analysis of x-ray diffraction scans for the electrodeposited Co-Fe alloys show that they undergo the structural transitions predicted by the phase diagram. The alloy compositions for the three distinct morphologies agree weU with daerent phases present in the Co-Fe system as shown in Table 5.5. The oniy difference apparent in this system is
that the phases appear to be shifted towards higher Co compositions. For example, in the phase diagram shown in Figure 5.19, the b.c.c. phase transition occurs at 25%Fe. However for the electrodeposited alloys this transition has occurred already at 22%Fe.
Table 5.5: Composition vs. Structure for Co-FeElectrodeposits
I
Composition
Morphology(8J
Structure
3.4 wt%Fe
Leaf
h.c.p.
6.6 wt%Fe
Leaf
h.c.p. + fC.C.
7.4 wt% Fe 10.5 wî% Fe
I
Nodule
Noduf e
~c.c.
15.0 wî% Fe
Nodule
fc.c. + b.c.c.
17.5 wt% Fe
nodule
fkc. + b.c.c.
Figure 5.2 1 illustrates the saturation magnetization (Ms)behavior of the electrodeposits. The saturation magnetization data initially follows previously measured data displaying an increasing magnetic moment with increasing iron content. Then the
Ms drops between 6wt0hiron and
15 wt% iron. At 15 wt% and above, the Ms again
reaches levels previously measured. In view of the srnail
ofgrain size on
saturation magnetization observed in NI Co, Ni-P and Ni-Fe it is unlikely that the observed drop in Ms for Co-Feis due to grain size effects.
*+
Weiss and Forrer(1929)
- -V- Weiss and Forrer(1929) -El-
Present Study (1 998)
Figure 5.21. Saturation magnetization of Co-Fedoys nom Wh-22wt%Fe.
Upon examination of Figure 5.21, it is evident that the data for pure Co and Co6.6%Fefall within the h.c.p. phase and agree well with the results of Weiss and
Forrer[lO]. Weiss and Forrer did not have data for Co-Fe doys between 5% and 13%Fe,
which corresponds to the range in which an unexpected drop in saturation magnetization was noted in this study. The drop is unlikely the result of grain site but more likely
linked to a phase transition. Normally, for cobalt, the fC.C. phase occurs at higher temperatures, however it is believed that this metastable structure created through electrodeposition is responsible for the decrease in the measured magnetic moment. The
reduced saturation magnetization noted in this alloy system can be attributed to the
'
presence of a weakly magnetic phase or weakly magnetic duai phase structure and not a result of variations in grain size or the presence of oxides. Figure 5.22 displays the coercivity behaviour as a function of composition for the
Co-Feelectrodeposits. In the h.c.p. or cobalt-rich transition to the fC.C. phase, the coercivity data takes on a parabolic arch type behavior. Then upon further increase of iron content, the coercivity fdls in what corresponds to the fC.C. region. A similar behavior is noted in the region where the d o y undergoes the f C.C. to b.c.c.transition. These patterns suggest that the transition regions between pure phases (dual phase regions) have either more pinning sites or are more stressed intemally as both coercivity
and retentivity increase in these regions. This intuitively makes sense as the presence of a second phase would provide more pinning sites which would make domain wall
movement more difncult. This analysis explains the apparent magnetic hardening of the
matenal in these regions. IO
-
h c ~
fcc
I
I
fa-sbcc
I
I
bcc
c.
E
s
2. .>
.-
E O
O
Composition (Co-WhFe)
Figure 5.22. Coercivity as a function of composition in Co-Fe electrodeposits.
5.4.4 Conclusions
The alloys examined in this sîudy had wmpositions which ranged fiom pure Co to CoW22wt%Fe.Both previous SEM studies and the current anaiysis of x-ray difiaction data reveaied that the structure of the electrodepositsagree with the equilibrium phase diagram for the system found in the literature. The h.c.p. cobalt, f C.C. cobalt and b.c.c. iron phases were al1 identifieci within the composition range ofalloys studied in this report. The saturation magnetization of the alloys agreed well with the literature in the
b.c.c. (alpha-iron) and the h.c.p. cobalt phases. However, a marked decrease in magnetic moment was observed in the f C.C. cobait region. This decrease could be attributed to a demeased magnetic moment of the f cc. phase at these compositions.
The coercivity of the alloys indicated that at the phase transition zones, the samples were either very stresseci due to the transformation of the Iattice or the formation
of intermediate, dual phase structures containhg more pinnllig sites served to make domain wdl movement and rotation more difficult. This translateci into increased
magnetic hardness in the dual phase region.
5.4.5
References
Bozorth, R.M., Ferromamaism, IEEE Press, New York 1978. Jiies, D., Introduction to Mametkm and Mametic, Chapman and Hall, Great
Britain, 1991.
Berkowitz, A., E. Kneiler, Mametisrn and Metallur~~ Vol l., Academic Press, 1969.
Cr& D.J.,R.S. Tebble, Ferromametism and Ferromametic Domains John Wiey and Sons, New York, 1965.
Chen, C.W., Sofi Maeentic Materials, North Hoiland Publishing, 1977. Ellis, W.C., Greiner, E.S.,Trans. Am. Soc. Metals 29,415, (1 941).
Forrer, R, J. phys. Radium, [7] 1,49, (1930). Greenberg, E*,BSc.E. Thesis, Queen's University, Kingston, Canada, 1996.
C.Cheung, Ph.D. thesis, in progress, Queen's University, Kingston, Canada, (1 998).
P. Weiss and R. Forrer, Compt. Rend., 189, 789, (1929) as cited in [l].
Chapter Six MAGNETOELECTROHYDROLYSIS OF
NANOCRYSTALLINE NI-FE AND CO
6.1
Introduction
As discussed in previous sections, conventional polycrystalline stmdures are unlikely
to meet the demands for new advanced soft magnetic materais. The magnetic properties required for many of these applications are as follows: high saturation magnetktion, low coercivity, high electrical resistivity and high Wear resistance[l]. Conventional soft magnetic matends with large grain site are relatively sofl mechanically, and as such, display a low Wear resistance.
Nanocrystalline materials would seem to be ideal for many soft magnetic applications given their high hardness and high electrical resistivity cornbined with a saturation
magnetization that is not significantly reduced as cornpared with theu polycrystaiiine counterparts if they are hlly dense as is the case for electrodeposited nanocrystals. Furthemore, the coercivity has been observeci to be reduced for nanocrystalline Finemet type materials as a result of the averaging of anisotropy constants and the vanishing of
magnetostnction [2,3]. In the first studies completed on the magnetic properties of electrodeposited Ni and
Ni-P outlined in sections 5.1 and 5.2 the saturation rnagnetization was not as high as required for many soft magnetic applications. Therefore it was suggested to add Fe to the Nckel electroplating bath to produce deposits with the Permalloy composition being Ni-
20%Fe. As previously diswssed, the magnetic anisotropy constants and magnetostndion approach zero at this composition yielding a highly permeable, and therefore excellent soft magnetic material with a low coercivity. Therefore, as detailed in section 5.3 Ni-Fe ailoys were produced in nanocrystahe fom[4]. It is well known that thermomagnetic anisotropy can be introduced by anneaüng in a magnetic field[5]. Ifmagnetization is applied parailel to the sample it is possible to "square up" the hysteresis loop as the sample is magnetized dong its easy direction and becomes difiïcult to demagnetire according to conventional Stoner-Wolfarth theory[6]. Conversely if the field is applied perpendicular to the sample it is possible to make the loop nearly linear. Applications of linear type hysteresis ioops would be for items which need to change magnetization rapidly with minimal core loss such as recording heads or transformers. Amplifiers and switching or digital storage elements make use of the 'square loop' characteristic in a magnetic material. The rnateial being nanocrystalline in nature cannot retain its srnall, metastable grain sue at traditional heat treatrnent temperatures used for large grainai materials and as a
result the analogy of thermomagnetic annealing was extended to electroplating in a magnetic field and provides the basis for this investigation. The rnethod of magnetoelectrohydrolysis has been used tu alter the magnetic
properties of electrodeposits in previous work[7]. As a result of discussions with a recording head manufacturer, regarding the fhrication of modem Pemalloy recording heads it was leamed that these heads are often electroplated in a magnetic field under proprietary conditions. It was hypothesized that the action of in-field plating wnied out during industrial fabrication served to alter the texture of the electrodeposit resulting in
some improvement in magnetic properties via induced ahisotropy or easy directions, which are more easily magnetized than other crystallographic directions according to
Stoner-WoIfgrth theory[8]. This assumption foms the basis for the foflowing magnetoelectrohydrolysis stuàies on Ni-20%Fe and Co. It should be noted that Ni20%Fe has a cubic structure whereas Co has a hexagonal close packed structure. As a
result of symmetry, Ui the cubic system, induced anisotropy as a result of in-field plating is expected to play a smaller role in altering magnetic behaviour when cornpared to the
hexagonal close packed system for Co which has a uniaxial easy direction of magnetization dong the basal plane.
C. Cheung developed proprietary electrolytes for electroplating Ni-20%Fe and Co in nanocrystailine form, which were used for this experiment[4]. The bath for Ni-20%Fe was a modiied nickel Watts bath, which uses iron chloride as a bath addition for the
source of iron in the electrodeposit. M. Fiynn, Queen's University, Kingston, Canada[P] designeci an electroplating ceil for plating in an extemal magnetic field. The bath was built to hold a volume of one liter and fit between pole pieces of an electromagnet as shown in Figure 6.1. The bath temperature was held at 4O0C+/-2°C with no agitation
during electrodeposition. Furthemore, electroplating was carried out in DC mode ushg a aiment density of 0.2 A/cm2. The pH of the bath was maintainecl between 3.8 and 4.1
to maintain constant plating conditions yielding samples of optimum quality. &Fe
samples were plated ont0 a titaniurn substrate and subsequently mechanically stnpped for
further evaluation. Using the above electroplating pakneters without an extemal field during plating has been show to yield Ni-2OWe deposits[lO].
DC plating was used to produce a polycrystalline cobalt sample with a grain size of 5pm whereas pulsed-cunent deposition was used to produce the nanocxystalline electrodeposits with a grain size of 15nm, fiom cobalt sulphate based baths. The electromagnet used to apply the magnetic field during plating was a 12"
Varian V4004 variable gap electromagnet. A RFL Mode1 904 gaussmeter was used to detennine the applied magnetic field strength dunng electroplating. The plating runs for the Ni-20%Fewere completed at field strengths of 0,75, 150 and 260 kA/rn. The plating
runs for Co were completed at applied field strengths of O and 260kA/m. However, samples that were plated in the perpendicular orientation had the cobalt replenishing electrode between the pole pieces of the magnet which effectively reduced the gap width and raised the field within the gap to 390 kAh.
Grain size of the Ni-Fe electrodeposits was deterrnined using the Schemer formula in conjunction with the Full Width Half Maximum
of the (1 11) peak in the NI-
Fe solid solution a s compared to the FWHM of the (1 1 1) peak for a nickel powder standard.
Magnetic measurements for this study were camied out on the Lakeshore mode1 7000 hybrid vibrating sample magnetometer. The electrodeposits were punched into 3rnm disks and stacked to form cylinders of constant volume for relative cornparison.
The applied field was 790kAlm to ensure cornplete saturation of the electrodeposit and the scan rate was 20 minutes.
.
Figure 6.1:
Apparatus developed for magnetoelectrohydrolysisof Ni-20%Fe.
6.3 Results and Discussion
Figure 6.2 displays the saturation magnetization behaviour as a function of the
applied field during electrodepositionof the Ni-Fe simples. It cm be seen that the
saturation magnetization for both perpendicular and parallel orientations increases with
increasing field strengths. In chapter 5.3 it was concluded that sahiration magnetization
of Ni-Fe is structure-independentbut closely related to the composition. Therefore the increase in moment is iikely due to an increase in the ion transport properties of iron in the bath as iron has a higher magnetic moment than nickel. As a result, increasing iron contents appear in the electrodeposit with increasing applied field during
electrodeposition. Unfortunately, the samples were destroyed in subsequent testing so
that no x-ray spectroscopy could be canied out to ver@ iron compositions.
Applied Field ( k N m ) O
Figure 6.2:
Substrate Parallel to Field Substrate Perpendicular to Field
Effects of in-field plating on saturation magnetization.
Figure 6.3 shows the behaviour of coercivity as a fùnction of applied field during electrodeposition of the Ni-Fe electrodeposits. There is no clear trend as far as coercivity is concemed. The texture, as examined by x-ray diaaction[9], for these electrodeposits does not change significantly during electrodepositionas such changes are most likely related to residual stresses in the sarnple as opposed to crystdographic shape or magnetic anisotropy.
Applied Field (kA/m) O
Figure 6.3:
Substrate Parallei to Field Substrate Perpendicular to Field
Effects of in-field plating on coercivity.
The results of the cobalt plating runs are summarized in Table 6.1 for both polycrystalline and nanocrystdhe Co the saturation magnetization rernained both constant and in agreement with the literature value of 1.4 x 10' A/m[ll] with and without the application of the extemal field.
The coercivity dropped with applied magnetic field in both the nanocrystahe .-_*. .
and polycrystalline electrodeposited cobalt. Furthemore it should be noted that the
coercivity dropped to a greater extent for samples plated in a magnetic field applied parallel to the substrate in both polycrystalline and nanocrystalline electrodeposits. The decrease in coercivity by nearly a factor of ten observed for nanocrystalline materials as opposed to polycrystalline materials is Iiely a result of crystaliographic texture differences as describeci in section 5.1. It can be seen that the n a n o c r y ~ t ~ ~ ~ e cobalt has a strong fiber texture as compared to the polycrystal1'ie sample electrodeposited fiom the sample bath which provides a basis for the expected lower coercivities observed in nanocrystalline cobalt.
Table 6.1:
Magnetoelectrohydrolysisof Cobalt
Plating Conditions
Nanocrystalline
Nanocrystalline
Polycrystalline
Ms ( N m * IO')
Hc (kA1m)
Hc (kA/m)
No Field
1.40
0.985
9.05
Pardel
I .40
0.880
7.69
Perpendicular
1.41
0.9425
8.52
6.4 Conclusions
There is a positive correlation between saturation magnetization and applied field
during magnetoelectrohydrolysis of Ni-Fe. The maximum saturation magnetization measured occurs for the sample plated in a field of 150Alm with the substrate parallel to the field. This increase in saturation magnetization is most Iücely due to the increased deposition of the iron atoms with increasing field strengths. It should be noted that at al1
applied field strengths, the samples electroplated with the substrate parailel to the applied field displayed higher saturation magnetization than those plated in the perpendicular configuration. There is no clear trend observeci in the coercivity behaviour for Ni-Fe electrodeposits in relation to the applied field strength. There is no effect on saturation magnetization with respect to applied field or
grain sue in the pure cobalt samples. Furthemore the saturation magnetization observed is in accordance with the literature value for polycrystaiiine Co[l l] and is in agreement with previous results presented in chapter five. Coercivity reductions in nanoprocessed Co samples are kely due to changes in crystaiiographic texture outlined in section 5.1. The coercivity decreases by a factor of nearly ten at al1 field strengths and configurations for the nanocrystalline electrodeposits
as cornpareci to the polycrystalline electrodeposits. The coercivity is lower for sarnples plated on a substrate parallel to the applied field. The lowest coercivity observed was
880Ahn for a nanocrystalline sample electroplated in the substrate pardel configuration. It should be noted that for both the polycrystalline and nanocrystailine Co similar reductions in coercivity were noteâ for both perpendicular and parallel substrate
configurations. In the case of the parallel substrate configuration the decrease in coercivity was on the order of 12%+/- 1% and for the perpendicular configuration the decrease in coercivity was on the order of 5% +/- 1%. This decrease is small when
compared to the reduction of coercivi~,for nanocrystdhe Co as compared to polycrystalline Co.
These results indicate that electroplating in a magnetic field wi improve the magnetic properties to a certain extent. Furthemore, non-cubic materials with a uniaxial easy direction of magnetization are likely to show greater changes in magnetic properties that relate to crystdlographic texture.
'
References
T.Jagelinski, MRS BuIletin, (3), 15,36, 1990.
G.Hener, Mat. Sci. Eng.,A133, 1, (1991). G. Herzer and H.Warhont, Encyclopedia of Materials Science and Engineering, R.W. Cahn, (ed.), Suppl. Vo1.3, Pergamon Press, 18 15, (1993). C. Cheung, Ph.D.thesis, in progress, Queen's University, Kingston, Canada.
R.M.Bozorth, Ferrornapnetism, D. Van Nostrand Co., New York, (1951). E.C. Stoner and E.P.Wolfarth, Phil. Trans. Roy. Soc., A240,599,(1948).
T.Z.Fahidy, J. App. Electrochem., 13, 553, (1983). R.A. McCume, Ferromagnetic Materials, E.P.WoIfarth, (ed.), Vol. 3, North HoIIand, 166, (1 982). M. Flynn, BScE. Thesis, Queen's University, Kingston, Canada, (1994).
C.Cheung, F. Djuanda, U.Erb and G. Palumbo, Nanostr. Mat., (5),5,513, (1 995).
D. Jiies, Introduction to Mametism and Mametic Materials Chapman and Hail, New York (19911.
Chapter Seven THERMOMAGNETIC STUDIES
7.1
Introduction
Recently, there has been a great deal of emphasis placed on the thermal treatment of amorphous and nanocrystalline alloys towards irnproving their magnetic properties[l171.
The system which has generated the most interest is the nanocrystallization of
arnorphous pre-cursor Fe[l], FeB[2-61, FeSiB ( ~ i n e m e -t type)[7-1 ~ S],and CoSiFe[l617ltype alloys with added Cu, Nb, Zr or Mo for soft magnetic applications. It is well
known that the rnagnetic properties of low alloy steels doys with added Cu, Nb have been improved upon during annealhg by foming a dispersion ofnanocrystalline (-1 Onm) F e 4 in a residual amorphous matrix[ 181. In conjunction with these results, the
soft magnetic properties are also enhanced as a result of the new microstmcture. The smail grain size causes the overall anisotropy constants to average and the magnetostricitive constants to tend towards zero, forming soft rnagnetic materials with properties comparable to traditional large grallied Permalloy type alloys[ 191. These studies are fundamental towards understanding the enhanced soft magnetic properties observed in nanocrystalline materiais since traditional soft magnetic materiais
have been large grainedI8-91. A recent review suinmarizes progress in these soft magnetic materials as shown in Figure 7.1 which compares the materials in terms of their l kHz pemeability and their saturation induction (Bs)[19].
One group of materials which is notably absent fiom this review are the Co-Fe base-alloys. Cobalt and iron are the only materials which when combined display a saturation magnetization that is higher than each of its constituents does. The synergistic action of this alloying should be exploited in future generations of soft magnetic
materials which make use of the highest saturation magnetization at 65%Fe-35%Co reaching 1.95 x 106 A/m as wmpared to the saturation mapetkation for pure iron of 1.71 x 106 A/rn[20].
It should be noted at this t h e that conventional large-grained Penndoy @ONi20Fe) has excellent sofk magnetic properties as a result of opposing r n a g n e t o î ~ s t ~ n e
anisotropies and magnetostriction for aiioys close to this composition[Z11. Further l o w e ~ of g the effective anisotropy has been shown to be theoretically possible through the reduction of grain size[22]. Assuming the ferromagnetic exchange length is constant
for decreasing grain size, the greater the number of grains involved in the ferromagnetic exchange length, the lower the overall effective anisotropy down to the amorphous
The low anisotropy acts to lower the domain waü energy and the low magnetostriction acts to reduce the pinning effects of defects and residual stresses. Table 7.1 explains the permeability maximum noted for 78% Ni in the Ni-Fe system[20].
Table 7.1:
Magnetostriction and anisotropy minima in Ni-Fe[l9]
Constant Minima
Composition
The overall effect of composition combined with appropriate heat treatments in Permdoy lead to a materiai with high initiai and maximum pemeabilities and low wercivity which, are the properties desired in a soft magnetic material.
Although the greatest thnist in this area is currently in the area of nanocrystalline materials produced by amorphous precursors there have been studies on nanocrystalline materiais produced by other methods as detailed in the magnetic literature review section (3.3.5). Initial studies on nanocrystahe materials indicated that both the magnetization
and Curie temperature for these materials were significantly reduced[23,24]. Chapter five has ciearty shown that the magnetization process is tittle dependent on grain size for fùlly dense electrodeposited rnaterials. For this chapter, annealing studies are designed to
reveal the following for nanocrystalline materials produced by electrodeposition:
The temperature dependence of Ms in nanocrystahe Niand Ni-Fe alloys. The Curie temperature of nanocrystallineNIand Ni-Fe alloys The effects of annealing on Hc in nanoctystalline Ni, Ni-P and Ni-Fe
It is well known that the magnetization process for ferromagnetic materials is strongly
dependent on temperature as the magnetization process involves the spin interactions between neighborhg atoms. In a ferromagnetic materiai, the overall magnetization of the matend is decreased with hcreasing temperature as the thermal energy acts to counter the magnetic ordering energy. When the thermal energy becomes large enough to overcome the magnetic ordering energy, the material undergoes a ferromagnetic to paramagnetic magnetic structure transition. This transition occurs at the Curie point as outlined in Chapter two. Nanocrystalline materials have a higher driving force for grain growth as compared to their polycrystalline wunterparts due to increased fiee energy associateci with larger
intercrystalline volume fiaction making them inherently metastable. In fact, the activation energies for grain growth in nanocrystallineNifor instance are similar to previous studies measuring grain boundaty self-difision in conventional matenals[25271. Table 7.2 shows grain growth peak temperatures and activation energies for grain
growth processes.
Table 7.2:
Peak grain growth temperatures and activation energies(251
Grain Growth Peak Temp. (OC)
lSm Ni
Activation
S°C/rnin
10°C/min
20°C/min
40°C/min
Energy
293
307
318
336
1.42 eV/atom
From Table 7.2 it is interesthg to note that w i t h this range of grain sizes (1030nm)it appears that the lower grain size has a higher thermal stability. Table 7.2 is included to show that grain growth in nanocrystalline NIand Ni-Fe occurs below the Curie temperatures of conventional materials examined in this study. For example, for
15nm Ni, with a heating rate of 5OClmin, the grain growth peak temperature i s given as
293°C which indicates that gain growth is occumng below the Curie temperature of 357OC for conventional polycrystalline Nckel given in the literature[Z 11.
Since the Curie temperature for conventional nickel and nickel iron alloys is above the point at which grain growth occurs it may be a very difncult problem to
rneaswe the Curie temperatures for nanocrystalliie materials at a given grain size. In deteminhg the Cune temperature, it follows fiom Table 7.2 and grain growth kinetics that the most accurate magnetization vs. temperature data will be obtained at the highest
temperature ramping rate possible as the materiai will spend the least amount of time at a temperature where grain growth occurs. Table 7.2 also indicates a large dEerence between the themal stability of nanocrystailine Ni and Ni-Fe. The suggestion of low mobility of an ordered Ni3Fephase has been made to explain these results[25]. It is possible to obtain ordered structures in
conventional Ni-Fe alloys with suitable heat treatments[28]. In fact these ordered structures display higher conductivity, saturation magnetization as well as a lower Curie temperature[21,29]. Since the lattice constants of &Fe and the solid solution Ni-Fe in the Ni-70.80% range and the atomic scattering factors for Ni and Fe are so sidar, x-ray dittiaction techniques cannot discem the structural dEerences[30]. One must rely on neutron difEaction[3 1,321, electrical resistivity[29] or magnetic studies[211 to deternllne the existence of an ordered phase. Thennomagnetic studies using in-situ hysteresis measurements while annealing have been perfonned on Finemet type alloys[33-361 as well as nanocrystalline Ni[37,38] and ultrathin Ni-Fe films[39]. Attempts were made in the Finemet type doys to determine the thermal onset of nanocrystallization fiom the amorphous preairsor materiai and to match the magnetic behaviour with the microstmcture[33-361. A thennomagnetic study
on nanocrystalline Ni produced through large deformation determined that high
coercivities are obtaind in the strained state and subsequently removed through a~ealiig[3 73.
The structure of Fiemet type nanocrystdiîne alloys upon annealing is nanocrystal1'me iron-rich grains in a residual amorphous matrix. In electrodeposited
nanocrystals, structural changes upon annealing above the grain growth temperabire wi take one of two forms: either a uniforni increase in grain size occurs or abnormal grain growth with large grains growing preferentiaily in a small grained matrix as previously shown[40]. Therefore, special consideration must be given to the resultant microstructure when ameaiing electrodeposited nanocrystalline matenals.
7.2 Temperature Dependence of Ms and Tc in Nanocrystalline Ni and Ni-Fe Alloys
7.2.1
Experimental
AM. El-Sherik produced nanocrystalline nickel samples of 99% punty, (1 Inm
grain size) at Queen's University, Kingston, Canada by electrodeposition as described in chapter five, section one. Nanocqstalline samples of nickel-iron ailoys in the O%Fe-
-
16.5%Fe range (2 1 14nm)were electrodeposited by C. Cheung, at Queen's University,
Kingston, Canada, using methods described in chapter five, section three. Energy dispersive x-ray spectroscopy (EDS) and x-ray difiction scans as described in the experimental sections of chapter five determined the composition and grain SURSof the deposits. The Lakeshore Cryogenics Mode1 7000 hybrid vibrating sample magnetometer was used to make magnetic measurements on the nanocrystallineNi and nanocrystalline
Ni-Fe electrodeposits. For pure nanocrystalline NI and each ofthe three compositions of nanocystalline Ni-Fe (Ni-3.5%Fe, Ni- 10.5%Fe and Ni- 16.5%Fe), several 3mm disks were stacked on top of each other to obtain sarnples of a constant weight and length to diameter ratio. The magnetic saturation of the samples was measured with respect to temperature in a microdurnace placed between the pole pieces of the rnag.net under a vacuum of approxhately 10" torr. An applied field of 790 W m was used for al1 measurernents to ensure magnetic saturation.
To minimize the effects ofgrain growth when obtaining magnetition behaviour with respect to temperature, the maximum heating rate ofthe oven, which is approxhately 50 degreedrnin, was used. It should be noted that table 7.2 gives peak
grain growth temperatures, and that the onset of grain growth occurs at approximately 1530°C below this temperature indicating cleariy that grain growth is occuning even at the highest temperature rarnping rate possible with this fumace.
7.2.2
Results and Discussion
Figure 7.2 shows the saturation magnetization behaviour for pure nanucrystaihe Ni and the nanocrystalline Ni-Fe alloys exarnined in this study. It should be noted that the
magnetization data has been converteci From electromagneticunits (emu) to Alm in accordance with SI units. It is seen that in ail cases, increases in temperature lead to decreases in the saturation magnetization. This is to be expected as the thermal energy is
acting in opposition to the magnetic wupling or exchange energy between neighboring
atoms. In ali cases,at temperatures close to the Curie temperatures for wnventional
.
polycrystailine materials the saturation magnefion of the nanocrystalline
electrodeposits is decreased to less than 1% of its magnetic saturation value at room temperature as show in Figure 7.2 .
Temperature (OC)
Figure 7.2:
Saturation magnetization behaviour as a fùnction of temperature
The saturation magnetization of the pure nickel sarnple decreases most rapidly and with increasing iron content the Ni-Fe curves are shified to higher temperatures and higher sahiration magnetization values which is to be expected, as both the Curie
temperature and magnetic moment of iron is pater than that for nickel.
As a result of the thermal instability of the nanocrystalline electrodeposits, the
Curie temperatures of these materials are impossible to measure precisely and must be approximated at best. When the temperature nears the Curie temperature the magnetization behaviour slowly tends to minimal values and as such the Curie temperature must be extrapolated fiom lower temperatures[41]. This is a result of the effects of the applied field (HJ on Ms. At low temperatures there is iittle eEect of the applied field, but at temperatures close to the Curie temperature the variation of Ms with
H becornes significant and as such the magnetization will never reach a zero value. One way to estimate the Curie temperature more amrately is to make use of the fact that Ms varies proportionally with (T- TC^'^ at temperatures close to the Curie temperature. This means that the plot shown in Figure 7.2 should be parabolic at these temperatures and by squaring the saturation magnetization values, and replotting with respect to temperature a straight line is expected. This line may be extrapolated to the temperature axis to give the Curie temperature. Figure 7.3 shows the results of squarhg the saturation magnetktion terms. It should be noted that in Figure 7.3, all data for which the applied field is sigmficantly c u ~ n away g fiom the straight-line portion of the extrapolation axis has been removed. Table 7.3 shows the calculated Curie temperatures as compared to the literature values of their polycrystalline counterparts[42]. It should be noted that in al1 cases the calculated value exceeds the literature value by approximately 3%. This is a result of the thermal mass of the sample and the disparity between the thermal couple and the sample. Extremely slow heating rates negate the effects of this thermal hysteresis, however it was important in these tests to have the maximum ramp rate possible to limit grain growth.
.
+Pure Ni
- -0- - Ni-3.5%Fe +Ni-lO.S%Fe -IP
O
100
200
300
-
Ni-16S%Fe
500
400
Temperature (OC)
Figure 7.3:
Determination of Curie tempemture through graphical analysis.
Repeat experiments showed no dinerence in saturation magnetization behaviour indicating that unless grain growth was extremely rapid, there is little t o no structural effect associated with magnetization behaviour as concluded in chapter five.
Table 7.3:
I
Curie Temperatures of Nanocrystalline Ni and Ni-Fe Calculated (OC)
Sample
LiteratureI421 ( O C )
Pure Ni
3 57
361
Ni-3.5%Fe
3 89
400
Ni-IO.S%Fe
480
494
Ni46,5%Fe
521
525
I
600
It should be noted that the Curie temperatures are within a few percent and the
error is attributed to the thermal rnass of the sample. Therefore there is no large change in
Curie temperature for nanocrystaiiine materials as previously reported. These results agree well with results obtained by McCrea et al. who shidied the temperature
dependence of saturation magnetization in electrodeposited nicke1[43]. The Cune temperature is based on the effects of thermal energy on the exchange energy. The thermal energy is dependent on the temperature and the exchange energy
has been s h o w to be independent of disorder[41]. Therefore it is expected that the Cune temperature will not be affectecl by an increase in disorder or a decrease in grain size.
7.2.3
Conclusions
The saturation magnetization decreases in all samplu with incressing temperature until the thermal energy exceeds the magnetic ordering or the exchange energy and the
material undergoes a ferromagnetic to paramagnetic transition. This is typical behaviour for a ferromagnetic material and repeat runs on the sarne sample show no difference in the dependence of saturation magnetization with temperature indicating that grain growth does not play a factor in this magnetic behaviour. Furthemore the magnetic saturation increases with increasing iron content as expected given that iron displays a higher magnetic moment per atom than nickel.
The Curie temperature for the nanocrystalline Ni and NoFe alloys tested in this study does not dser significantly from that observed in polycrystdine materials as previously reported. The small differences between the literahire values for polycrystalline materials and Curie temperatures obtained in this study are believed to be due to thermal hysteresis as al1 temperatures were slightly higher than literature values. This can be attnbuted to the large temperature ramping rate used in an effort to minirnize grain growth. Further experimentation at slower ramp rates is needed to afiirm these results.
7.3
Effects of Annuiling on Coercivity of Nanocrystalline Ni, Ni-P and Ni-20%Fe
7.3.1 Experirnental
AM. El-Sherik produced nanocrystaiiine nickel samples of 99% purity (1lm grain size), at Queen's University, Kingston, Canada by elechodeposition as described in
chapter five, section one. Nanocrystalline NP was obtained from Ontario Hydro in the f o m of an electrodeposit produced in-situ on the inside of a stem generator tube using a proprietary plating solution as part of the ~lectrosleeve~ nuclear stem generator repair process[44]. The bath contains proprietary arnounts of nickel sulphate, nickel carbonate, phosphorous acid, phosphoric acid and saccharin. The sampies for magnetic measurements consisted of 10 disks, which were cut using EDM from the electrosleeved portion of a stearn generator tube. Nanocrystalline samplu of nickel-iron alloy with the
composition of Ni-20%Fewere electrodeposited by C. Cheung, at Queen's University, Kingston, Canada, usiig rnethods described in chapter five, section three. The nanocrystalline Nckel deposits were annealeci in an argon atmosphere at a temperature of 2 5 0 ' ~for 20,60 and 120min. The ~lectrosleeve~ electrodeposits were
annealed in an argon-filleci stainless steel bag in a molten sait bath at a temperature of 360°Cfor 60 minutes. The nanocrystalline Ni-20%Fesamples were annealed in a
protective stainless steel bag in a molten salt bath at 220°C and 400°C for tirnes of 30, 60 and 120 minutes. Ali magnetic rneasurements were completed on the Lakeshore mode1 7000 hybnd
vibrating sample magnetometer (VSM). Fuli hysteresis Ioops were generating using an applied field of 790 W m to ensure complete magnetic saturation. Coercivities were extracted fiom the hysteresis loops and converted to metnc units (kA/m).
7.3.2
Results and Discussion
Figure 7.4 shows the coercivity behaviour of the electrodeposited nanoaystallùie nickel dter annealing @ 250°C. The coercivity decreases by a factor of 3 after annealing for only 20 minutes. The coercivity of sofl magnetic materials has been shown to decrease with increasing grain size as dixïissed previously[8,9]. However, this explanation for the observed decreases in coercivity can be discarded as this material has
been shown to be t h e d l y stable at a temperature of 250°C (see Table 7.2).
There are several causes of residual stresses in electrodeposited matends including co-deposited, impurities and hydrogen. The observai decrease in coercivity
can be attributed to stress relieving effects such as the reduction of entrained hydrogen in the sarnple during the 250°C treatment.
O
20
40
60
80
1O0
120
140
Annealing Tirne (minutes)
Figure 7.4:
Coercivity as a function of annealing time for nanocrystalline nickel.
Figure 7.5. shows the coercivity behaviour of electrosleeved NiP as a fiinction of
annealing time. The coercivity again decreases by a factor of 2 d e r annealing for 60
minutes at 350°C. It has been show previously that this material is themally stable at this ternperatureL451, thus the decreases in coercivity cannot be explained in terms of increased grah size. Again it is expected that stress relief due to lowering of hydrogen is
responsible for the severe decrease in coercivity. Figure 7.5 also shows the behaviour of as-plated nanocrystalline Ni and annealed nanocrystalline Nifor comparative purposes.
Tirne (minutes)
60
-
Nanocrystalline Ni 250%
0Nanocrystalline Ni-P - 350°C
Figure 7.5:
Coercivity of Ni and Ni-P electrodeposits before and after annealing.
Figure 7.6 shows the behaviour of coercivity of electrodeposited nanocrystalline Ni-20wt%Fein the as plated condition and after annealing at 220°C and 400°C for 30, 60 and 120 minutes. The coercivity drops by a factor of about 1.2 during the annealing
process. NanocrystaliineNi-20%Fe aruided at 220°C showed no grain growth
behaviour whereas the sarnples annealed at 400°C had an initial grain size of approximately 15nm and a final grain size af€erannealing of 40nrn[46].
20
40
60
80
1O0
Annealing Tirne (minutes)
Figure 7.6:
Coercivity of annealeci nanocrystalline Ni-20wt%Fe.
120
140
This decrease in coercivity upon initial annealhg treatments both above and below the grain growth temperature is believed to be due to stress relief in the materiai after plating. In addition to the removal of residual stresses, it has been show that
annealing at temperatures above 200°C is effective for the removal of entraineci hydrogen, which can act as both an embrittling mechanism and a stress raiser. This behaviour is not only obseived in nanocrystalline electrodeposits but also in the nanocrystalline materials produced fiom amorphous precurson, which are currently
receiving worldwide attentionr33-391.
7.3.3
Conclusions
The coercivity of nanocrystalline Ni, Ni-P and Ni-20%Fe electrodeposits is observeci to decrease upon a stress reliehg anneal at temperatures high enough to remove entrainecl hydrogen but not likely high enough to cause significant grain
growth[40,45]. These preliminary thennomagnetic studies provide a good basis for the exploration of heat treatments on the magnetic properties of nanocrystalline transition metais and ailoys.
7.4
References:
[1]
N.X. Sun, K.Zhang, X.H. Zhang, X.D. Liu and K. Lu, Nanostr. Mat., (6), 7,637, (1996).
[2]
A. Makino, A. Inoue and T. Masumoto, Mat. Trans., JIM, (7), 36,924, (1.995).
[3]
N. Shiga, F. K o g h and M. Yukumoto, Mat. Trans., J M , (7), 36,939, (1995).
[4]
J. Y.Park, S. J. Suh, K.Y. Kim and T.H.Noh, IEEE Trans. Mag., (5),33,3 799, (1997).
151
B.Idzikowski, J. Baszynski, 1. Skovanek, K.-H. Muller and D. Eckert, J. Magn. Magn. Mater., 177-181,94 1, (1998). K-S Kim, S-C Yu, Y-MMoon, and K.V. Rao, 5. Magn. Magn. Mater., 177-181, 968, (1998).
Y. Yoshizawa, S. Oguma and K. Yamauchi, J. Appl. Phys., 64,6044, (1988). G. Herzer, IEEE.Trans.Mag., (5),25,3327. (1989).
G. Hemr, Mat. Sci. Eng., A133, 1, (1991).
X-Y.Xiong and K-YHo, J. Appl. Phys., (S), 77,2094, (1994). A.R. Yavari and 0. Drbohlav, Mat. Trans.,JIM, (7), 36, 896, (1995). S. Surinach, A. Otero, M.D.Baro, AM. Tonejc and D. Bagovic, Nanostr. Mat., 6, 46 1, (1995).
J. Arcas, C. Gomez-Polo, A Zhukov, M. Vazquez, V. Larin and A. Hemando, Nanostr. Mat., (8), 7, 823, (1996).
T.Liu, T.D. Hu,Y.N. Xe, Z.T.Zhao and RZ. Ma, Nanostr. Mat., (7),8,909, (1997).
K.G. Efihimiadis, S.C. Chadjivasiliou, E.K.Polychroniadis, M. Ozer, G.A. Sterigoudis and I.A. Tsoukalas, J. Magn. Magn. Mater., 185, 187, (1998).
A. Serebryakov, V. Sedykh, V. Stelrnukh and N. Novokhatskaya, Nanostr. Mat.,
(9,7,5 19, (1996). H. Koshiba, A. Inoue and A. Makino, Nanostr. Mat.,@), 8,997, (1997). AR Yavari and D. Negri,Nanostr. Mat., (8), 8,969, (1997). K. Yamauchi and Y. Yoshizawa, Nanostr. Mat., 6,247, (1995).
D. Jiles, Introduction to Mametism and Mametic Materials Chapman and Hall, New York, (1991).
RM. Bomrth, Ferromametism, D. Van Nostrand, New York, (1% 1).
H. Hofian, Thin Solid Films, 58,223, (1979) as cited in 171. H. Gleiter, Prog. Mat. Sci., 33,223, (1989).
R.Z.Vaiiev, Ya.D. Vishnyakov, R.R Mulyukov, and G.S.Fainshtelli, phys. Stat. Sol. (a), 117, 549, (1990).
T. Turi, Ph.D. Thesis, Queen's University, Kingston, Canada, (1997). A.R. Wazzan, J. Appl. Phys., 36,222, (1965) as cited in [25]. 1. Kaur, W. Gust, and L.Koma, eds., Handbook of Grain Boundary and
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O.Dahl, 2.Metdkde, 28, 133, (1936). R.J. Wakelin, and E.L.Yates, JOURNAL MISSING, 221, (1952). F.E.Haworth, Phys. Rev., 54, 693, (1938). F.C. Nix, H.G. Beyer and J.R Dunning, Phys. Rev., 58, 1031, (1940). C. G. Shull and M.K.Wilkinson, Phys. Rev., (2), 97, 304, (1955). R. Grossinger, R.S. Turteiii, V.H. Duong and C. Kuss, in Mametism Materials and their Apdications, eds. F. Leccabue and V. Sagredo, (Cod Proc. 3" Latin
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L.K.Varga, E. Kisdi-Koszo, V. Strom and K.V.Rao, J. Magn. Magn. Mater., 159, L X 1, (1996). A. Hernando, P. Main, M. Vaquez and G. Hener, J. Magn.Magn. Mater., 177-
181,959, (1998).
Kh.Ya. Mulyukov, G.F.Korznikova, R.Z. Abdulov and R.Z. Valiev, J. Magn.
Magn.Mater., 89,207, (1990). Y.D. Yao, Y.Y.Chen, C.M.Hsu, H.M. Lin,C.Y. Tung, M.F. Tai, D.H.Wang,
KT.Wu and C.T. Suo, Nanostr. Mat., 6,933, (1995).
C.Hou. H.Fujiwara, T.J.KJemmer, R.M.Metzger and W.D.Doyle, IEEE Trans. Mag., (5), 33,3625, (1997).
U. Klement, U. Erb and KT.Aust, Nanostr. Mat., 6, 58 1, (1995). J.P. Jakubovics, Mametism and Mametic Materids,
L.J.Swartzendruber, V.P. Itkh and C.B.Alcock, I. Phase Equil., 12,288, (1991). J. McCrea, Ph.D. thesis, in progress, University of Toronto, Toronto, Canada, (1 998).
G. Palumbo,P.C.Lichtenberger, F. Gotualez and AM. Brennenstuhl, "Process and Apparatus for In-situ Electroforming a Structural Layer of Metal Bonded to
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F. Gonzalez, AM. Bremenstuhi, G. Palumbo, U. Erb and P.C.Lichtenberger, Mat. Sci. Forum,225-227,831, (1996). D. Ueno, BSc.E. Thesis, Queen's University, Kingston, Ontario, (1994).
Chapter Eight ANALYSIS, TRENDS AND APPLICATIONS FOR
NANOCRYSTALLINE MAGNETIC MATERIALS 8.1
Magnetism of Nanocrystalline Soft Ferromagnets
The eEeds of grain boundary stmctural disorder on the saturation magnetization and the local magnetic moment of atoms within a grain boundary has been a constant source of debate for nanocrystalhe materials since their discovery in the early 1980's. It was initially expected that nanocrystalline materials would exhibit magnetic properties, which are diierent than their polycrystaîiine counterparts, as a result of the large increases in intercrystalline volume fraction due to the extremely refined dimensions of ordered grains in nanocrystalline materials[2]. In addition, as the grain size of a material approaches the critical dimensions of exchange length and domain size, further changes in magnetic properties for nanocrystaliine materials were hypothesiied.
In early studies, there was a great deal of discrepancy between reported results for nanocrystallùie materiais. For example, a reduction in 40% in saturation magnetization was observed for 6nm nanociystalline iron produced by inert gas condensation[l]. Early
small angle neutron scattering studies indicated that the intercrystalline cornponent (i.e. grain boundaries and triple junctions) in consolidateci gas condensed nanocrystalline
materials exhibit either a weakly or non-magnetic behaviour[f]. Research on ultrafine particles reported significant decreases in saturation magnetization in conjunction with a reduction in grain size[4-6]. The formation of antiferromagneticoxide layers on the
surfaces of the consolidatecl particles were the proposed mechanisrn for the reduction in saturation magnetization. In contrast to -dies
complet& on materials requiing a post-processing
consolidation step, (i.e. bd-rnilIing, or inert gas condensation) results presented in the body of this work for fhly dense, electrodeposited nanocrystalline systems have shown that there is no, or very little reduction in saturation magnetization with decreasing grain size. This structural independence has been observed for electrodeposited Ni, Ni-P, NiFe, Co and Co-Fe.
Recent theoretical studies using the application of Tight Binding Liear-MuffinTin-Orbital Atomic-Sphere-Approximation(TB-LMTO-ASA) have andyzed the local magnetic moment for structures representing grain boundaries of varying degrees of disorder[7]. The main conclusion of this work stated, although there exists some variation in magnetic moment for various grain boundaries and grain boundary sites, the average magnetic moment is relatively insensitive to stnicture[7].
Furthexmore, these
calculations have been extended to account for alloys containing non-magnetic elements such as P and W of various concentrations and it has been observed that the magnetic saturation is much more iduenced by compositional rather than structural disorder[$]. Saturation magnetization results for morphous, nanocrystalline, microcrystalline and conventional polyc~stailinematerials fkom the present study and elsewhere[9] were
compared with electronic structure caiculations using TB-LMïO-ASA as shown for Ni-P and Co-W in figures 8.1 and 8.2 respectively[8]. The saturation magnetization values observeû for fully dense electrodeposited nanocrystalline matenals agree well with calculated values for buk materials with the same composition.
Present Calculations Aus [26] Admon e t al. 193 Luborsky e t al. [IO] Farcas e t al. [ I I ]
5
O
15
10
20
Atornic Z Tungsten
Figure 8.1:
Magnetization in Co-W ailoys as a function of tungsten concentration.
Filled squares: theoretical caicrifations[7], Experimental[9-111. Present Calculations
a o v A
O
O
5
1O
Atomic
Figure 8.2:
Aus 126.41 Albert et al. [12] Bakonyi e t al. [13] Sonnberger e t al. [14] Huller et al. [lS] Berrada e t al. [16]
15
20
25
30
Z Phosphorus
Magnetization in NLP alloys as a fùnction of phosphoms concentration. Filled squares: theoretical calculations[7], Expenmental[12-161.
While the magnetic moment and saturation magnetization are structure-insensitive properties other magnetic properties are strongly affectecl by structure namely coercivity, hysteresis and magnetic domains.
In early work on nanocrystailine matenals, it was believed that below a critical grain size, each grain represented a single domain as described by the classical Stoner-
Wolfarth model of rotational hysteresis[l7]. For example, reported increases in Hc in ultrafine particles with grain sizes below a critical grain size (e.g. -40nm for Ni), based on energy considerations for domain formation, have been attributed to the transition
from multi-domain to s i i e domain particles[4,6]. For grain sizes considerabiy smaller than the size of domain walls in conventional materials significant decreases in coercivity have been reported for Finemet type sofk magnetic alloys[ 18-20]. The coercivity is observed to increase with decreasing grain size at grain sizes above the nanocystalline range and to decrease sharply to levels observed
in amorphous matenals at extremely fme grain sizes as show in Figure 8.3[20]. The D~ dependence of coercivity on grain size for grain sizes below 50nm has been explained by the random anisotropy model which accounts for the averaging of anisotropy as a result
of the exchange interaction or grain coupling when the grain size is below the ferromagnetic exchange length[20,2 11.
O' O'
50N;Fe
a perni-
olloy
G r a i n Size
Figun 8.3:
O
Coercivity (Hc) vs. grain size @) for various soft magnetic ailoys.
Nanocrystalhne materials include: FeNbSiB, FeCuVSiB, FeZrB, and FeCoZr[20].
In materials where the grain size is srnaller than the ferromagnetic exchange length OH), there is some disagreement on the formation or even the existence of
magnetic domains. According to the randorn anisotropy model, reductions in grain size in the nanometer region lads to the lowering of the effective anisotropy (Km) which is responsible for magnetic domain formation. With the lowering of r70pi2cm)would reduce core losses significantly in devices such as D.C. motors and transformer cores. Eddy current loss prevention is a
high priority when choosing transformer core material.
.
Early research indicated that nanocrystalline materials wouid be excellent candidates for a new generation of soft magnetic materials. As a result of the research h s t towards this end in the past decade, high saturation flux density, hi&
pemeabiiity
zero magnetostriction magnetic alloys for use in toroidal cores, choke coiis and p~lse transfomers have been developed[43]. The relationship between saturation induction
flux density (Bs)and lkHz permeabiliiy (p) for these soft magnetic materials are shown in Figure 8.7. The nanoprocesseci alioys discussed in the body of this research and those shown
in Figure 8.7 represent technologically significant advancements in rnaterials available
for soft magnetic applications. Nanoprocessed soft magnetic d o y s with high saturation magnetization, electrical resistivity, Wear resistance and low coercivities represent excellent venues for future applications such as high performance transfomiers, motors, amplifiers and field sensors.
Figure 8.7:
Relationship between Bs and p at llcHz for NANOPERM~,the nanocqstalline Fe-Si-B-Nb-Cu alloys and conventional soft magnets[43].
In contrast, recent research into nanocrystalline composites for hard magnetic applications is yielding promising resuIts[44-491. In materials traditionally used for hi&energy permanent magnetic applications (e.g. NdFeB), nanoprocessing has been show to enhance the remanence through exchange coupling[46-481. In addition, composites
formed with sofi and hard magnetic phases show large saturation induction combined with increases in coercivity which are required magnetic properties for permanent magnet
materials[46-48].
The abdity to tailor application specific magnetic properties combined with other advantages of nanoprocessing will continue to reveal novel applications for future soft
and hard magnetic matends.
8.3
References
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K. Huller, G. Die@ R. Hausmann and K.Kolpiin, J. Magn. Magn. Mater., 53, 103, (1 985).
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Chapter Nine CONCLUSIONS
(1)
The saturation magnetization and coercivity of pore-fiee electroplated bulk nanocrystaliine nickel was reported for the first tirne. For grain sizes down to
lOnm the saturation magnetization is relatively little affecteci by grain size which is in good agreement with magnetic theory which States that saturation magnetization is a stmcture-insensitive property. The coercivity initidy decreases with grain size and regches a minimum at a grain s k of 40nrn and increases to a maximum at a further refined grain size of 10m. (2)
The saturation magnetkation was observeci to be independent of grain size for electrodeposited cobalt in contrast to nanostructures produced as unwmpacted or compacted powders containing porosity.
The differences in saturation
magnetization between electrodeposited cobalt and powder or compacted particle
cobalt are attributed to formation of antiferromagnetic surfàce layers on the latter.
The effects of texture on coercivity are more pronounced in this system as compared to cubic systems as a result of the uniaxial anisotropy in h.c.p. structured materials. The coercivity is r e d u d by a factor of ten in the nanocrystdine electrodeposit as compared to the polycrystalline as a result of the strong
basal plane crystallographic texture in the nanocrystailine electrodeposit
comesponding to the easy direction of magnetization for cobalt.
(3)
The saturation magnetization of the electrodeposited nickel phosphoms was found to decrease with increasing phosphoms content which is in good agreement with
suppression of ferromagnetic character previously reported for polycrystalline NiP alloys with increasing amounts of phosphonis. In addition, the materiai underwent
a ferromagnetic to paramagnetic transition at approxhately 16at%P that coincided with the structurai transition of the aiioy fiom nanocrystaliine to amorphous. The
wercivity was strongly affecteci by the gmUi s k with a large maximum occurring
at a grain size of 3m and composition of 14.5aW~corresponding to previously noted sharp crystallographic texture change to cl1l> at this grain sue.
(4)
The saturation magnetiiation of electrodeposited nanocrystalline Ni-Fe alloys, in the compositional range, MFe-28wt%Fe, increased linearly with increasing iron content.
Furthemore, the saturation magnetization of the electrodeposits is
structure-insensitive displaying 97% of the literature value for saturation magnetization.
The coercivity of the electrodeposited nanocrystalline Ni-Fe
deneases with increasing iron content in good agreement with changes in magnetostriction and anisotropy constants for diEerent compositions of the alloy.
(5)
The saturation magnetization and coercivity of pore-& electroplated bulk cobaltiron alloys with iron concentrations up to 22wt%Fe were reported.
They
displayed three diierent stnictures as identifid by x-ray difEaction analysis. The
h.c.p. cobalt phase was observed for low iron contents whereas the f C.C. cobalt
phase was observed at higher iron concentrations, and deposits with the highest
iron concentrations displayed a b.c.c. structure. The saturation magnetization for the alloys which displayed the h.c.p. cobalt and b.c.c. iron structures agreed well with values reported in the Iiterature. However a marked decrease in magnetic moment was observed in the f CA. region. The coercivity was markedly increased at phase transition zones (i.e. zones displayhg two phases).
This was attributed to
increases in stress during structural transformation of the sample as well as the formation of intemediate dual phase structures leading to an increase in domain
w d pinning effects.
(6)
The saturation magnetization and mercivity were measured for nanocrystalhe Ni2OwtOAFe samples produced by magnetohydrolysis with the substrate both parallel
and perpendicular to the applied field. There was a positive correlation between the saturation magnetization and the applied field during plating. This increase in saturation magnetization was attributed to increased deposition of iron during plating due to increased effects of field on iron, which has a higher magnetic
moment than nickel. Furthemore, samples that were electroplated in a field with the substrate parailel to the applied field displayed increased saturation magnetization and hence an increased iron concentration than those plated in the perpendicular configuration. Althouth, there were no clear trends in coercivity behaviour with regards to field strength or substnite orientation, the large observed
fluctuations in coercivity are most likely related to intemal stresses in conjunction with the stress sensitivity of Ni-Fe alloys.
'
(7)
The saniration magnetization and coercivity were rneasured for nanocrystalline Co
samples produced by magnetohydrolysis with the substrate both parallel and perpendicular to the applid field. There was no effect on saturation magnetization
with respect to applied field, grain size or substrate orientation with al1 samples displaying 100% of the literature value for saturation magnetization for cobalt. The coercivity is reduced by the development of a strong texture during electrodeposition of the nanocrystalline Co. The coercivity decreased by a factor
of ten for the nanocrystalline electrodeposits as compared to the polycrystailine electrodeposits. The coercivity was lower for those materials plated with the substrate parailel to the applied field as opposed to those plated with the substrate perpendicular to the applied field. These results indicate that both nanoprocessing
and electroplating in a magnetic field can improve the soft magnetic properties by lowering the coercivity.
(8)
The Curie temperatures and temperature dependence of saturation magnetization were measured for nanociystalline electrodeposited NieFe in the range M % F e to
16.5wt%Fee.The saturation magnetization increases with increasing iron contents at ail temperatures. The Curie temperature for the nanocrystalline Ni-Fe alloys does not diEer from that observed in polycrystalline materials in contnist to previous reports. (9)
The coercivity of nanocrystalline Ni, Ni-Iwt%P and NitOwt%Fe electrodeposits was rneasured in samples before and after annealing. The coercivity was observed
to decrease in al1 cases upon annealing below the grain growth temperature. This
decrease in coercivity is believed to be due to stress relieving effects via the removai of hydrogen f?om the electrodeposit.
Chapter Ten FUTURE WORK
(1)
Further work is required in order have absolute data for plotting hysteresis loops. More specifically, the demagnetizing âctors for the geometries of the shon cylinders or discs used in this study have not been calculated. Therefore the only data which is correct în the absolute sense is the coercivity and saturation magnetization. With correcteci hysteresis loops, the Uiitial, relative and maximum pemeability as well as retentivity could be accurately calculated.
(2)
A better understanding of the coercivity behaviour would be possible if the
experiments were designed to isolate the effects of residual stress, crystallographic texture, composition and grain sue. In this iàshion, al1 contributions to the
complex behaviour of coercivity could be accounted for and a predictive mode1 of the hysteresis behaviour as a funaion of grain size could be developed.
(3)
In cases where texture played a significant role in &ve
behaviour, attempts to
quant@ the effects of anisotropy as a function of texture would be useful in predicting coercive behaviour. New nanocrystalline samples with different crystaUographic texture and the sarne grain s i i and composition would have to be produced for this future work
(4)
An effort should be made to ascertain whether or not samples have a dual phase
structure upon annealhg or electrodeposition in the case of Co-Feand to quanti@ both the amount and effect of these phases on the magnetic properties of the material. In the case of Co-Fe the decrease in magnetic moment as a result of the
fC.C. phase could then be calculated.
.
