A NEW GENERATION OF Fe35%Co SOFT MAGNETIC ALLOYS CONTAINING ADDITIONS OF Nb, Ta AND Ni

A thesis presented for the Degree of Doctor of Philosophy of the University of London and for the Diploma of Imperial College.

by

Simon A. Hall, BSc.(Hons.)f A.R.C.S.

Department of Materials Imperial College of Science, Technology and Medicine London, SW7.

December 1990. Abstract

It has recently been found that additions of 0.2 to 0.3 % Nb or Ta to Fe35%Co produces an alloy which can be fabricated industrially. A full investigation of this new generation of soft magnetic alloys is now required.

As with the established Fe50%Co base alloys, a rapid quench is still neccessary prior to cold forming as the 35% alloys can also order sufficiently to produce a brittle material. It is shown that x-ray superlattice measurements are not appropriate for measuring ordering kinetics, and that the lattice expansion upon ordering in the 50% alloys does not occur in the 35% alloys. To glean information about the ordering kinetics of the 35%Co alloy, a new technique for measuring the relaxation of the frozen-in disorder is developed using DSC.

The relation between ternary additions and grain size is investigated for both 50.and 35%Co base alloys. In general ternary additions produce a material with smaller grains. A ductile to brittle transition occurs in the 35%Co alloys at ambient temperatures, whereas for 50%Co alloys this transition is at 0°C. This DBTT is dependent on grain size and hence ternary addition* It is also found through tensile tests that the disordered 35%Co alloys are more ductile than the 50%Co alloys.

The 35%Co alloys harden with heat treating at temperatures in the range 400 to 600°C due to the formation and dissolution of a fine precipitate. This is likely to be the cause of the grain refinement assosciated with ternary additions. Furthermore, the precipitate helps in the formation of a {100|<110> texture. This discovery was used with varying success, to enhance the magnetic susceptibility of all the alloys. In some Fe50%Co base alloys the anisotropy was such that the coercive force was reduced by more than 50% by a suitable choice of magnetising direction with respect to the rolling direction. CONTENTS PAGE

Introduction...... 4

CHAPTER 1: Review of the literature.

1.1.0 Phase Equilibria and the Effect of Ternary Additions...... 4 1.1.1 The Iron Cobalt Phase Diagram...... 4 1.1.2 Ternary additions to Binary iron cobalt...... 8 1.1.3 Ordering of Fe50%Co and Fe35%Co...... 8 1.1.4 Mechanism of ordering in undeformed alloys...... 11 1.1.5 Ordering of cold worked material...... 14 1.1.6 Ordering Energies...... 14 1.1.7 Disordering Energy...... 17 1.1.8 Methods of measuring long range order...... 17

1.2.0 The Atomic Structure of Binary Iron Cobalt Alloys...... 19 1.2.1 Variation of magnetic moments with iron to cobalt ratio...... 19 1.2.2 Variation of the mean magnetic moment upon ordering...... 22 1.2.3 Variation of the iron hyperfine field upon ordering...... 22 1.2.4 Variation of the cobalt hyperfine field upon ordering...... 24 1.2.5 Effect of Ternary Additions on Electronic Structure...... 24

1.3.0 The 550 peak in Fe Co binary alloys...... 27 1.3.1 History...... 27 1.3.2 Formation of the peak...... 27 1.3.3 Effect of precipitation in ternary alloys on the '550 Anomoly'...... 29 1.3.4 Application of the '550 effect’ to ternary alloys...... 29

1.4.0 Recovery and recrystallisation in cold worked alloys...... 31 1.4.1 Iron cobalt binary alloys...... 31 1.4.2 Recrystallisation in Iron Cobalt alloys with ternary additions...... 35

1.5.0 Precipitation of if-phase in ternary alloys...... 37 1.5.1 Precipitation in 2V-Permendur type alloys...... 37 1.5.2 Additions of nickel to 2V-Permendur...... 37 1.5.3 Additions of Ni and Nb to Fe50%Co...... 42 1.5.4 Effect of cold work on precipitation...... 42 1.5.5 Consequences of precipitation...... 44 1.5.6 Precipitation in Fe35%Co Ternaries...... 44

1.6.0 Texture...... 45 1.6.1 The Rolling Texture of BCC Alloys...... 45 1.6.2 The Texture of Binary Fe35%Co...... 45 1.6.3 The Texture of Binary Fe50%Co...... 46 1.6.4 The Texture of Fe50%Co2%V...... 46 1.6.5 Consequences of texture...... 47 1.6.5.1 Magnetocrystalline anisotropy...... 47 1.6.5.1.1 Values of Ki and K2...... 50 1.6.5.1.2 Magnetostriction...... 53 1.6.5.2 Microstructural effects due to texture...... 53 1.6.5.3 Variation of mechanical properties with direction...... 55 CHAPTER 2: Experimental Details.

2.1.0 Sample preparation...... 56 2.1.1 Arc Melting...... 56 2.1.2 Hot Rolling...... 56 2.1.3 Cold Rolling...... 56 2.1.4 Samples produced by industry...... 57 2.1.5 Heat Treatments...... 57 2.1.6 Compositions...... 58

2.2.0 Microscopy...... 59 2.2.1 Light Microscopy...... 59 2.2.2 Scanning Electron Microscopy, (S.E.M.)...... 59 2.2.3 Transmission Electron Microscopy, (T.E.M.)...... 59

2.3.0 X-Ray Diffraction...... 61 2.3.1.1 Theory of lattice parameter and order...... 61 2.3.1.2 Lattice Parameter Measurements...... 61 2.3.2.1 Theory of Superlattice Measurements...... 62 2.3.2.2 Superlattice Measurements...... 62 2.3.3 Texture Measurements...... 63

2.4.0 Differential Thermodynamic Techniques...... 66 2.4.1 Differential Scanning Calorimetry...... 66 2.4.1.1 Methods used in the D.S.C. study...... 66 2.4.1.2 D.S.C. Experiments...... 67 2.4.2 Differential Thermal Analysis (D.T.A.) and Thermo-Gravimetric Analysis (T.G.A.) Methods...... 67

2.5.0 Mechanical Testing...... 68 2.5.1 Impact Tests to investigate industrial problems with cold rolling...... 68 2.5.2.1 Tensile Tests...... 69 2.5.2.2 Test conditions...... 70 2.5.3 Hardness and microhardness tests...... 70

2.6.0 Magnetic Measurements...... 71 2.6.1 Saturation Magnetisation using the Sucksmith Balance...... 71 2.6.2 The Vibrating Sample Magnetometer, V.S.M...... 71 2.6.3 Permeammeter Measurements...... 75 2.6.3.1 Sample preparation and geometry...... 77

CHAPTER 3: Results.

3.1.0 Grain size and mechanical properties...... 78 3.1.1 Changes in grain size on annealing...... 78 3.1.2 Grain size and cold rollability...... 78 3.1.2.1 Rollability of different alloys with different grain sizes..78 3.1.2.2 Rollability of Fe35%Co0.2%Ta samples with different grain size...... 78 3.1.3 Impact tests on hot rolled sheet...... 84 3.1.3.1 Existence of a Ductile to Brittle Transition Temperature, (DBTT)...... 84 3.1.3.2 Dependence of DBTT on grain size...... 84 3.1.3.3 Impact tests at higher temperatures...... 84 3.1.3.4 S.E.M. of fracture surfaces of impact test specimens...... 89

/ 3.2.0 Hardness results...... 93 3.2.1 Dependence of hardness on annealing temperature...... 93 3.2.2.1 Effect of composition on hardening...... 93 3.2.2.2 Effect of prior cold work on hardening...... 97 3.2.3 Hardness change upon ordering...... 97 3.2.4 Variation in hardness on isothermal annealing...... 97 3.2.4.1 Annealing at 550°C...... 94 3.2.4.2 Annealing at 650°C...... 102

3.3.0 Texture results...... 105 3.3.1 Assessment of texture...... 105 3.3.2 Cold rolled textures...... 105 3.3.3 Recrystallisation textures of alloys with differing cold deformation...... 105 3.3.4 Texture and heat treatments...... 116 3.3.5 Two stage heat treatments...... 116 3.3.6 {100}<110> texture dependence on ternary additions...... 120

3.4.0 Results of magnetic measurements...... 124 3.4.1 Sucksmith balance results...... 124 3.4.2 Vibrating sample magnetometer measurements on textured material...... 124 3.4.2.1 General Results...... 124 3.4.2.2 Anisotropic Results...... 131 3.4.3 Permeammeter measurements on textured material...... 131 3.4.3.1 Comparison of B-H curves for the single stage and the two stage heat treatments...... 135 3.4.3.2 Comparison of the two B-H curves of samples cut in the roll direction...... 136 3.4.3.3 Comparison of the two B-H curves of samples cut at 45° to the roll direction...... 136 3.4.3.4 Best and Worst magnetic properties due to heat treatment and orientation...... 136 3.4.3.5 Error assessment...... 137

3.5.0 X-Ray Results...... 138 3.5.1 Lattice parameter measurements...... 138 3.5.2 Superlattice measurements...... 143

3.6.0 Thermal Analysis Results...... 146 3.6.1 Transition temperatures...... 146 3.6.1.1 Binary alloys...... 146 3.6.1.2 Additions of Nbf Ta and Ni to binary alloys...... 146 3.6.1.3 Addition of 3.5%Ni to Fe50%Col.8%V (Rotelloy 5)...... 150 3.6.1.4 Additions of Ce, Y and Misch metal to Fe35%Co...... 150 3.6.2 The 550 anomqly...... 150 3.6.2.1 The peak in binary alloys...... 150 3.6.2.2 Cancellation of peaks...... 153 3.6.2.3 Use of an annealed, ordered sample as a reference...... 153 3.6.2.4 Oxidation and the DSC trace...... 156

3.7.1.Tensile test results...... 159 CHAPTER 4: Discussion.

4.1.0 Impact tests, grain size and reliability...... 160 4.1.1 DBTT and reliability...... 160 4.1.2 Mode of fracture in impact test specimens...... 163 4.1.3 Grain size refinement by heat treatments...... 163 4.1.4 Grain size refinement by composition...... 163 4.1.5 Other effects of ternary on reliability...... 165

4.2.0 Magnetic anisotropy...... 166 4.2.1 Calculation of deformation of spherical precipitate...... 170 4.2.2.1 Allowing for precipitates to form in magnetised domains___ 175 4.2.2.2 Change in easy direction with ordering...... 176 4.2.3 High aspect ratio precipitates in a crystallographic direction...... 176 4.2.4 Elongated precipitates formed in the roll direction...... 179 4.2.5 The most likely explanation of experimental results...... 179 4.2.6 Summary...... 182

4.3.0 Texture discussion...... 183 4.3.1 Factors that determine texture...... 183 4.3.2 Explanations of recrystallised texture development...... 183 4.3.3.1 The same recrystallised texture as cold rolled texture.... 183 4.3.3.2 Recrystallised textures differing from the deformation texture...... 184 4.3.4 Effect of precipitate on texture...... 184 4.3.5 Likely explanation of textures observed in FeCo ternary alloys...... 185 4.3.6 Consequence of ternary additions...... 186

4.4.0 Hardening in FeCo base alloys...... 187 4.4.1 Order hardening...... 187 4.4.2 Order hardening in cold worked alloys...... 191 4.4.3 Precipitation hardening in cold worked alloys...... 191 4.4.4.0 Explanation of hardening in this work...... 194 4.4.4.1 Variation of hardness with anneal temperature...... 194 4.4.4.2 Variation in hardness with isothermal annealing...... 196

4.5.0 DSC Discussion...... 198 4.5.1 The derivation of the threshold temperature for atomic movement...... 198 4.5.2 Ordering activation energies from Td peak vs Heating rate graphs...... 202 4.5.3. Mathematical modelling of low temperature ordering...... 202 4.5.4 Fitting a theoretical curve to the experimental data...... 210 4.5.5 Why E derived by Kissinger method does not not agree with E derived from fitted equations...... 215 4.5.6 Reasons why theoretical fit to Fe35%Co base alloys is poor__ 215 4.5.7 Shape of the theoretical curves...... 216 4.5.8 Lattice parameter and diffusion energy reinforcement of lower ordering energy...... 216 4.5.9 Consequences of finding a full mathematical model of the system 218 CHAPTER 5: Conclusions.

5.1 Conclusions and implications...... 219 5.2 Factors that determine cold reliability...... 219 5.3 Rolling and recrystallised textures...... 220 5.4 Precipitation hardening in cold rolled alloys...... 220 5.5 Saturation magnetisation...... 220 5.6 Magnetic anisotropy...... 221 5.7 Methods of measuring LRO in Fe35%Co alloys...... 222

CHAPTER 6: Further work.

6.1 Suggestions for further work— ...... 223 6.1.1 Improvements in magnetic properties for DC applications.....223 6.1.2 Improvements in magnetic properties for AC applications.....223 6.1.3 Alternative methods of observing the '550'anomoly...... 224

CHAPTER 7

7.0 References...... 225

Appendix 1: Bragg Williams Theory of the Long Range Order Parameter...... 230

Acknowledgements 232 List of Figures.

Figure 1.1.1: The Co-Fe phase diagram...... 5 Figure 1.1.2: The mean atomic magnetic moment from saturation magnetisation measurements versus Co content for BCC iron cobalt alloys at 0 K ...... 6 Figure 1.1.3: The lattice parameter of FeCo binary alloys in the ordered and disordered condition...... 6 Figure 1.1.4: B-H curves for technically important and other FeCo base alloys...... 7 Figure 1.1.5: A comparison of isothermal ordering of FeCo alloys at 440°C, after quenching from 780°C...... 9 Figure 1.1.6: The change in equilibrium order with temperature..... 10 Figure 1.1.7: Isothermal ordering curves for Fe50%Co using the lattice parameter method...... 13 Figure 1.1.8: Isothermal domain growth for Fe50%Co2%V...... 15 Figure 1.1.9: The effect of cold ork on isothermal ordering...... 15 Figure 1.2.1: The mean magnetic moment of Fe atoms in FeCo alloys versus Co concentration of disordered and ordered alloys...... 20 Figure 1.2.2: Magnetic moments attributed to 3d electrons in FeCo alloys versus Co concentration...... 20 Figure 1.2.3: The average number of electrons per atom in 3d up, 3d down and 4s states versus Co concentration for disordered FeCo alloys...... 21 Figure 1.2.4: The average 57Fe hyperfine magnetic field for disordered and ordered FeCo...... 21 Figure 1.2.5: Effective magnetic field at the iron nuclei for FeCo alloys...... 23 Figure 1.2.6: Calculated local magnetic moments of d impurities in Fe...... 23 Figure 1.2.7: Calculated local magnetic moments of d impurities in ordered Fe50%Co...... 26 Figure 1.2.8: Variation of the mean saturation moment per atom in ternary FeCo alloys with concentration for various soute elements...... 26 Figure 1.3.1: Specific heat versus temperature upon constant heating rate, of various heat treated Fe50%Co alloys..28 Figure 1.3.2: FeCo alloy with a high degree of LRO approaching the equilibrium degree of order...... 28 Figure 1.3.3: A disordered sample of FeCo approaching the equilibrium degree of order...... 30 Figure 1.3.4: The 550°C anomoly observed by specific heat, for initially ordered FeCo alloys...... 30 Figure 1.4.1: Variation in hardness of 80% cold deformed FeCo alloys with temperature after 2hr annealing..... 32 Figure 1.4.2: Variation in the recrystallisation parameters of FeCo binary alloys...... 32 Figure 1.4.3: Grain growth kinetics during a recrystallisation anneal of Fe53at%Co & Fe34at%Co...... 34 Figure 1.4.4: Anisotropy of cold rolled strip annealed for 2 hours at indicated temperatures...... 34 Figure 1.5.1: FeCo-2%V ternary section...... 40 Figure 1.5.2: TTT curve for annealed Fe47%Col5%V3.5%Ni compared with that for FeCo2%V...... 41 Figure 1.5.3: Dependence of grain size on ^ precipitate...... 41 Figure 1.5.4: TTT curves for annealed and cold worked FeCo2%V...... 43 Figure 1.6.1: The magnetisation of single crystals of iron and nickel in different crystallographic directions...48 Figure 1.6.2: Directions of easy magnetisation for cubic crystals as a function of the anisotropy constants Ki, and K2..48 Figure 1.6.3: Perspective drawings of the magnetocrystalline energy surfaces for cubic crystals...... 49 Figure 1.6.4: The variation of the first and second anisotropy constants Ki and K2 with composition in binary FeCo alloys...... 51 Figure 1.6.5: The saturation magnetostriction in the [100] and [111] crystallographic directions for FeCo alloys...... 54

Figure 2.3.1: The variation in received intensity with angle from the sample normal for iron powder...... 64 Figure 2.3.2: Spiral contour plots of the raw data from the texture goniometer, illustrating how a fictitious jagged ridge can be produced...... 65 Figure 2.6.1: Schematic diagram of the sample chamber of the vibrating sample magnetometer...... 73 Figure 2.6.2: Block diagram of the permeammeter...... 76

Figure 3.1.1: Grain size of Fe50%Co0.35%Nb0.5%Mo after 2 hour heat treatments...... 79 Figure 3.1.2: Grain size of Fe35%Co0.2%Ta after 2 hour heat treatments...... 80 Figure 3.1.3: Grain size of hot rolled Fe50%Co base alloys...... 81 Figure 3.1.4: Grain size of hot rolled Fe35%Co base alloys...... 82 Figure 3.1.5: Effect of grain size on reliability of hot rolled and quenched Fe35%Co0.2%Ta...... 83 Figure 3.1.6: Impact test results for Fe50%Co2%V...... 85 Figure 3.1.7: Impact test results for Fe50%Co0.2%V0.4%Nb...... 86 Figure 3.1.8: Impact test results for Fe50%Co0.35%Nb0.5%Mo...... 86 Figure 3.1.9: Impact test results for Fe35%C00.2%Ta...... 87 Figure 3.1.10: Impact test results for Fe35%Co0.3%Ta...... 87 Figure 3.1.11: The grain size dependence of the DBTT...... 88 Figure 3.1.12: SEM micrographs of typical fracture surfaces...... 90 Figure 3.2.1: Variation of macrohardness with annealing...... 94 Figure 3.2.2: Variation of microhardness with annealing...... 95 Figure 3.2.3: Microhardness of Fe35%Co0.3%Ta with l-4%Ni...... 96 Figure 3.2.4: Effect of deformation on macrohardness...... 98 Figure 3.2.5: Variation in macrohardness with isothermal annealing at 550°C...... 99 Figure 3.2.6: Variation in microhardness with isothermal annealing at 550°C...... 100 Figure 3.2.7: Variation in macrohardness with isothermal annealing at 550°C...... 101 Figure 3.2.8: Variation in macrohardness with isothermal annealing at 650°C...... 103 Figure 3.2.9: Variation in microhardness with isothermal annealing at 650°C...... 104 Figure 3.3.1: Pole positions in the (200) and (110) pole figures for the predominant textures observed in BCC metals...... 106 Figure 3.3.2: Relief and contour pole figure for Fe50%Col.8%V3.5%Ni (Rotelloy 5) cold rolled 88%..... 107 Figure 3.3.3: Relief and contour pole figure for Fe50%Col.8%V3.5%Ni (Rotelloy 5) cold rolled 92%..... 108 Figure 3.3.4: Relief and contour pole figure for Fe50%Col.8%V3.5%Ni (Rotelloy 5) cold rolled 96%..... 109 Figure 3.3.5: Strength of central peak of (200) pole figures for cold rolled and annealed (760°C/2hrs) samples__ .110 Figure 3.3.6: Relief and contour pole figure for recrystallalised Fe50%Col. 8%V3.5%Ni...... Ill Figure 3.3.7: Relief and contour pole figure for recrystallalised Fe50%Co2%Vf (cold rolled 96% and 760°C/2hrs)...... 112 Figure 3.3.8: Relief and contour pole figure for recrystallalised Fe50%Col.8%V3.5%Ni, (Rotelloy 5, cold rolled 96% and 760°C/2hrs)...... 113 Figure 3.3.9: Relief and contour pole figure for recrystallalised Fe35%CoO.3%Tal%Ni, (cold rolled 95% and 760°C/2hrs)..114 Figure 3.3.10: Relief and contour pole figure for recrystallalised Fe35%Co0.3%Ta, (cold rolled 95% and 760°C/2hrs).... 115 Figure 3.3.11: Relief and contour pole figure for recrystallalised Fe35%Co0.3%Ta4%Ni, (cold rolled 95% and 575°C/2hrs).117 Figure 3.3.12: Relief and contour pole figure for recrystallalised Fe35%Co0.3%Ta, (cold rolled 95% and 670°C/2hrs).... 118 Figure 3.3.13: Relief and contour pole figure for recrystallalised Fe35%Co0.3%Taf (cold rolled 92% and 750°C/2hrs).... 119 Figure 3.3.14: Relief and contour pole figure for 86% cold rolled Fe50%Co2%V after 2 stage heat treatment of Table 3.3.1...... 121 Figure 3.3.15: Relief and contour pole figures for 96% cold rolled Fe50%Col.8%V3.5%Ni...... 122 Figure 3.3.16: The height of the central peak of the (200)pole figure for various samples after the heat treatment that gave the strongest {100}<110> texture...... 123 Figure 3.4.1: Saturation magnetisation of iron cobalt base alloys measured by the Sucksmith Balance technique...... 125 Figure 3.4.2: A comparison of the induction curves of cold rolled and annealed Rotelloy 5 ...... 126 Figure 3.4.3: A comparison of the induction curves of cold rolled and annealed Fe35%Co0.3%Ta4%Ni...... 127 Figure 3.4.4: The induction curves of recovered 95% cold rolled Fe35%Co0.3%Ta4%Ni...... 128 Figure 3.4.5: Magnetic properties of samples with different ternary additions...... 130 Figure 3.4.6: Magnetic properties of samples cut from textured Fe35%Co0.3%Tal%Ni sheet as determined by V.S.M...... 132 Figure 3.4.7: Magnetic properties of samples cut from textured Fe35%Co0.3%Ta4%Ni sheet as determined by V.S.M...... 133 Figure 3.4.8: Normal induction curves for textured Rotelloy 5 sheet measured in the roll direction and at 45° to the roll direction...... 134 Figure 3.4.9: Normal induction curves for isotropic ring specimen cut from 88% deformed sheet...... 134 Figure 3.5.1: The effect of a 465°C anneal on the lattice parameter of initially disordered iron cobalt base alloys..... 139 Figure 3.5.2: The change in lattice parameter upon isothermal ordering at 465°C of two initially disordered Fe35%Co base alloys...... 140 Figure 3.5.3: Rate of isothermal ordering of two initially disordered Fe50%Co base alloys introduced to a furnace set at 675°C...... 141 Figure 3.5.4: The rate of heating of a thermocouple introduced to a furnace at 675°C 142 Figure 3.5.5: Superlattice and fundamental x-ray diffraction lines for iron cobalt alloys ...... 144 Figure 3.6.1: D.T.A. trace of Fe50%Co0.1%Ta showing order/disorder and alpha/gamma phase transitions.... 147 Figure 3.6.2: D.S.C. trace of Fe49%Co0.2%V0.4%Nb showing '550 anomoly’, order/disorder and alpha/gamma phase transitions...... 147 Figure 3.6.3: D.T.A. trace of Fe35%Co0.2%Ta showing much smaller order/disorder peak...... 148 Figure 3.6.4: D.S.C. trace of Fe35%Co0.3%Nb0.3Misch showing double peak in heating cycle at s500°C and s550°C in addition to the order/disorder and alpha/gamma peaks.148 Figure 3.6.5: D.S.C. trace of binary Fe35%Co showing that peak at 500°C cannot be due to precipitation...... 151 Figure 3.6.6: D.S.C. trace of binary Fe50%Co showing exothermic peak at 525°C in first heating and endothermic peak at 600°C in second heating, characteristic of the '550 anomaly'...... 151 Figure 3.6.7: D.S.C. trace of Fe50%Co2%V repeated three times without changing the sample, showing the effect of oxidation and ordering...... 152 Figure 3.6.8: D.S.C. trace of Fe35%Co0.3%Ta repeated three times without changing the sample, showing the effect of oxidation and ordering...... 152 Figure 3.6.9: D.S.C. trace of identical reference and sample of Rotelloy 5 (Fe50%Col.8%V3.5%Ni) illustrating how complete cancellation of peaks can be achieved...... 154 Figure 3.6.10: D.S.C. trace of disordered Fe50%Co2%V run against an ordered reference, isolating the '550 peak'..... 155 Figure 3.6.11: D.S.C. trace of disordered Fe35%Co0.3%Tal%Ni run against an ordered reference, showing the '550 peak' which is not present if the run is repeated__ 155 Figure 3.6.12: Heating curves for Fe35%Co base alloys with different references used to isolate oxidation peak...... 157 Figure 3.6.13: Heating curves for Fe50%Co2%V with different references used to isolate oxidation peak...... 157 Figure 3.6.14: Combined Thermogravimetric and D.T.A. traces of Fe50%Co2%V showing how some of '550 peak' is due to oxidation above 450°C, indicated by weight gain..158

Figure 4.1.1: Schematic ductile to brittle transitions in iron cobalt alloys...... 162 Figure 4.2.1: SEM micrographs of Rotelloy 5 cold rolled 96% and recrystallised...... 169 Figure 4.2.2: Deformation of precipiates due to magnetostriction__ 172 Figure 4.2.3: Illustration of direction cosines...... 173 Figure 4.2.4: TEM micrograph showing how dislocation tangles obscure observation of precipitates...... 180 Figure 4.4.1: Comparison of the as quenched and elevated temperature flow stress for binary Fe50%Co...... 188 Figure 4.4.2: Comparison of the as quenched and elevated temperature flow stress for Fe50%Co2%V...... 189 Figure 4.4.3: Comparison of the as quenched and elevated temperature flow stress for Fe50%Co2%V measured at different strain rates...... 190 Figure 4.4.4: Effect of annealing temperature on the mechanical properties of 90% cold rolled FeCo-2%V alloy...... 192 Figure 4.4.5: Variation of the 0.2% yield stress and the ultimate tensile stress as a function of heat treatment temperature for 90% cold worked FeCo-2%V alloy...... 192 Figure 4.4.6: Mechanical properties of 91.7% cold worked FeCo2%-V versus quench temperature...... 193 Figure 4.4.7: Yield strength versus aging temperature of >90% cold rolled FeCo-2.8%V...... 193 Figure 4.4.8: Microhardness of 90% cold rolled FeCo-2%V alloy as a function of heat treatment temperature___ 195 Figure 4.4.9: Isothermal hardening of 90% cold worked FeCo-2%V.... 195 Figure 4.5.1: Experimental DSC curves of '550 anomoly' for Fe50%Co2%V at different heating rates...... 199 Figure 4.5.2: Experimental DSC curves of '550 anomoly' for Fe35%Co0.2%Ta at different heating rates...... 200 Figure 4.5.3: Extrapolation of ordering peaks at different ramp rates to give ordering threshold temperatures__ 201 Figure 4.5.4: Graph to obtain the ordering activation energies by the Kissinger method...... 203 Figure 4.5.5: Variation of the temperature derivative of electrical resistivity caused by the ordering or disordering, &(Xr (.) measured at a heating rate of 2°/min after quenching from various temperatures..206 Figure 4.5.6: Variations of (.) measured at the indicated heating rates after cooling from 800°C...... 207 Figure 4.5.7: Calculated curves of dX/dT vs temperature for Fe35%Co using Tahara's model...... 212 Figure 4.5.8: Calculated curves of dX/dT vs temperature for Fe50%Co using Tahara's model...... 212 Figure 4.5.9: Activation energies derived from the calculated DSC peaks using Kissinger's Method...... 214 Figure 4.5.10: Calculated specific heat curves for Fe35%Co...... 217 Figure 4.5.11: Calculated specific heat curves for Fe50%Co...... 217 List of tables.

Table 4.5.1: Threshold temperatures and ordering ranges...... 198 Table 4.5.2: Parameters for best fit curves...... 211 Table 4.5.3: Calculated ’550 peak' temperatures...... 213 Table 1.1.1: Ordering activation energies...... 16 Table 1.1.2: Diffusion activation energies...... 17 Table 1.1.3: Techniques for measuring L.R.0...... 18 Table 1.4.1: Recrystallisation threshold temperatures...... 35 Table 1.5.1: Precipitation of second phase in FeCo base alloys...... 38 Table 1.5.2: The classification of elements on their effectiveness in imparting ductility to FeCo alloys...... 39 Table 1.6.1: Directions of easiest, intermediate and hardest magnetisation in cubic crystals...... 50 Table 1.6.2: Values of Ki and K2 for binary FeCo alloys...... 52 Table 1.6.3: Effect of test direction on mechanical properties..... 55 Table 2.1.1: Materials and the main experiments they were used in...58 Table 2.1: Tensile samples...... 69 Table 2.6.1: Details of V.S.M samples...... 75 Table 2.6.2: Samples for permeameter measurements...... 77 Table 3.1.1: Impact test data for Fe35%Co0.3%Ta...... 89 Table 3.2.1: Effect of ordering on macrohardness...... 97 Table 3.3.1: Two stage heat treatments...... 120 Table 3.4.1: Summary of VSM results...... 128 Table 3.4.2: B-H characteristics of textured Rotelloy 5...... 135 Table 3.5.1: X-ray diffraction line intensities for fully ordered samples...... 145 Table 3.6.1: Transition temperatures for FeCo base alloys...... 149 Table 3.7.1: Tensile test results...... 159 Table 4.2.1: Magnetostriction values...... 170 Table 4.2.2: Magnetostriction values...... 174 Table 4.2.3: Fractional change of the axes of a sphere due to magnetostriction for Fe50%Co alloys...... 175 Table 4.2.4: Change in strain of a rod in Fe50%Co...... 178 Introduction

The iron cobalt magnetic system has been known to have high saturation magnetisation since 1912, when Weiss and Preuss (1912) reported the highest saturation magnetisation of any binary in Fe35%Co. In 1929, Elman patented the Fe50%Co alloy and named it Permendur. It was used for small magnetic pole pieces of loud speakers and sound recorders, and had the advantage of a lower permeability than the Fe35%Co composition with only a small sacrifice in saturation. Permendur and related alloys are still widely used in speaker and receiver diaphragms and this is reflected by the large number of papers produced by the Bell Telephone Laboratories, e.g. Bozorth (1951).

Permendur, as it stood, was extremely brittle, and in 1932, White and Wahl were granted a USA patent for FeCo2%V which they called 2V-Permendur. Under the right processing conditions, the vanadium ternary imparted ductility to the equiatomic binary and increased its resistivity by a factor of about 20. It was now possible to produce soft magnetic laminates with acceptable resistivity and consequently, 2V-Permendur is now used in aerospace a.c. applications where weight saving is more important than cost. Rotors and stators in aircraft generators are a good example, and iron cobalt alloys have been employed in aircraft as glamorous as Concorde.

Additions of more than 5% vanadium to various FeCo compositions lead to permanent magnetic materials with trade names of Vicalloy or Reraendur. Less expensive chromium is also used in alloys of this type. They offer high saturation magnetisation and can be machined.

Towards the end of the thirties the BCC to B2 ordering reaction at around 720°C was discovered. The ordered alloy is brittle and the disordered phase must be retained in a quench if the alloy is to be rolled or machined, and, to this day it is not clear exactly how vanadium assists in imparting ductility to the quenched alloy.

It was in the early fifties that the incredibly high magnetostrictive properties of iron cobalt alloys were investigated. The Fe70%Co alloy was reported to have the highest known magnetostriction of the day of +130 x 10-6. A consequence is that components under tensile stress,

- 1 - such as modern rotors, have enhanced magnetic properties, (Hoses et al. 1975). These high magnetostrictive values also led to FeCo alloys being assessed for use in transducers.

In 1957, Gould and Wenny reported the discovery of Supermendur, a magnetically annealed FeCo-2%V alloy with a much squarer B-H loop than the 2V-Permendur. Selected properties are shown below:

Permendur Supermendur Max permeability 5 000 70 000 Coercive force (Oe) 1.00 0.23 Residual induction (G) 10 000 21 400 Saturation magnetisation (G) 23 500 24 000

(Source Metals Handbook, 9th edition)

Supermendur finds uses in specialised transformers where weight saving is critical, such as radar pulse transformers. Unfortunately, in drum type motors and generators, designs are such that anisotropic materials are not advantageous. However, pancake type electric motors can be designed to take advantage of the superior magnetic performance offered by anisotropic materials, A dramatic illustration of this came in 1989, when "Lynch" motors, using grain orientated iron silicon, were used to propel the Duchess of Arran to a new water speed record for an electric powered craft.

The last property that iron cobalt alloys offer is an extremely high Curie temperature of about 930°C. This has led to both 2V- Permendur and Superraendur being used as temperature compensators. Another use of the high Curie temperature was made in the eighties when sintered FeCo was used for a 24 pin dot matrix printer head which was reported to print faster than any other head in the world.

Most of the alloys and applications mentioned have involved equiatomic based compositions. However, alloys based on 27%Co and 35%Co do exist; the 27%Co alloy is inherently less brittle and less subject to degradation by stresses and is appropriate for both a.c. and d.c. applications whereas the 35%Co alloy can be alloyed with chromium (Hiperco 35) or more recently, with niobium or tantalum and produces a

-2- less expensive material with the highest saturation, suitable for d.c. applications. Medical body scanners are an example of equipment where these alloys are finding applications. Little is known about the mechanical properties of these alloys and anisotropy has not been developed to try to bring the a.c. magnetic properties up to those of Supermendur.

This thesis aims to investigate and improve both the magnetic and mechanical properties of the Fe35%Co base alloys through the addition of Nb, Ta and Ni. A review of the atomic structure of the FeCo system is provided to explain why Fe35%Co has a higher saturation magnetisation than Fe50%Co, and how the magnetic moments of ternary additions align with the matrix.

The exact role of various additions is not clear and the thesis work aims to increased understanding. Additions may retard the ordering kinetics and hence enable the more ductile disordered phase to be retained after quenching. This work develops a new method of measuring ordering kinetics based on thermal analysis.

The ternary addition may lead to precipitation effects and hence affect the grain structure. The dependence of grain size and, consequently, the mechanical properties, on composition is also investigated.

Finally, precipitation, grain size, deformation, and annealing temperature often leads to a crystallographic texture. Other magnetic systems, for example Fe3%Si, make use of this anisotropy to enhance susceptibility, and lower coercivity in a given crystallographic direction. A thorough texture investigation of both the Fe35%Co and Fe50%Co base alloys is undertaken.

-3 - 1: Review of the literature

1.1.0 Phase Equilibria and the Effect of Ternary Additions.

1.1.1 The Iron Cobalt Phase Diagram.

The Co-Fe phase diagram is shown in Figure 1.1.1. The workable disordered a-phase is BCC and can be retained in a rapid laboratory quench. The ordered a'-phase, which exists between 30% and 70% Co at temperatures below s700°C, has a (i-brass structure and is inherently brittle. It is this phase that prevents binary iron cobalt alloys from being produced industrially. Above ~1000°C there is an FCC ft-phase, which on quenching, tra^j'ofiTiSto a martensitic BCC phase, ai.

Both the a and a' phases are strongly ferromagnetic up to a very high Curie temperature. The saturation magnetisation for different cobalt concentrations is shown in Figure 1.1.2, and is explained in Section 2. It is important to notice that the ordered structure has a higher saturation magnetisation than the disordered structure for a given composition. According to Shiga (1981), the change in saturation magnetisation can largely be accounted for by the concomitant change in lattice parameter. He reports an empirical relationship between the lattice constant, a(x), and the mean magnetic moment, <|m|>:

a(x) = aco(1-x) + aFeX + C<|m|>

where ac© , aFe and C are constants, and x is the atomic fraction of Fe. He reports excellent agreement between calculated values and experimental values, Figure 1.1.3. However, Asano et al., 1964, report lattice parameter values that do not relate to the mean magnetic moment for the cobalt rich alloys. He reports a lattice parameter expansion upon ordering in the Fe50%Co in agreement with Shiga, but a contraction upon ordering in the Fe35%Co, Figure 1.1.3.

In addition to saturation magnetisation, susceptibility and coercivity are important. Figure 1.1.4. shows that Fe50%Co base alloys are superior to the Fe35%Co base alloys in these areas. In particular, Fe50%Co2%V benefits from magnetic annealing to produce a very square hysteresis curve (Supermendur).

-4 - Temporal urn UC Temperature 1700------■ J 2 3 4 5 6 7 B 0 100 00 B0 70 60 50 40 30 20 JO 0 i J ------Figure 1.1.1: The Co-Fe phase diagram. diagram. phase Co-Fe The 1.1.1: Figure Nsiaa n Ihd 1984) Ishida and (Nishizawa egt ecn Iron Percent Weight tmi Pret Iron Percent ic Atom ------5 ------0 JO 20 30 40 50 60 70 Ai% Cobalt

Figure 1.1.2: The mean atomic magnetic moment from saturation magnetisation measurements versus Co content, for BCC FeCo alloys at 0 K. (Bardos 1969)

Figure 1.1.3: The lattice parameter of FeCo binary alloys in the ordered and disordered condition. (Shiga 1981, Asano et al. 1967)

- 6 - a)

FIELD, H (A/cml

r igure^ 1. j_4 : B H curves for technically important and other FeCo base alloys. (a: Pfeifer & Radeloff 1980, b: Chen 1904)

b)

- 7 - 1.1.2 Ternary additions to Binary iron cobalt.

Both Fe35%Co and Fe50%Co must be alloyed with a ternary^to enable cold rolling under normal industrial conditions, and until recently, it was only possible to make use of the high saturation magnetisation in the Fe50%Co base alloy. This was normally achieved by alloying 2-3%V to produce commercial alloys such as Permendur or Remendur. Chromium has also been used as a ternary addition. These alloys have been studied for over 40 years and detailed information can be found in the PhD. theses of Pitt 1980, Persiano 1986, and Orrock 1986. It is assumed that the ternary^helps in retaining disorder in the quench between hot and cold rolling, though the work of Eymery, Grosbras and Moine (1974), Glazyrina et al. (1983), and Clegg and Buckley (1973), suggest that vanadium accelerates isothermal ordering of previously quenched disordered samples, (Figure 1.1.5), contrary to the work of Orrock. Eymery et al. (1974), propose that the faster ordering rate is due to the alloy with vanadium additions retaining a higher vacancy concentration, though vacancy concentration is not thought to have a direct effect on reliability. Kawahara (1983a,1983c), proposed that elements that formed a CoaX precipitate produced rollable alloys. There are others, e.g Shiryayev et al. (1984), who show that interstitial C, H, 0 or N produce a more brittle alloy. Vanadium is known to have a strong affinity for these elements and may draw them from the surrounding matrix, thus facilitating dislocation movement. The fact remains that, for whatever reason, the vanadium alloys are ductile enough to cold roll after an industrial quench.

Vanadium additions do not enable the Fe35%Co alloy to be cold rolled. They require additions of <1% Ta or Nb, and as yet very little work has been done to explain why th e s e ^ elements work and others do not.

1.1.3 Ordering of Fe50%Co and Fe35%Co.

The order/disorder transition temperatures of Fe50%Co and Fe35%Co are 725° and 670°C respectively (Oyedele & Collins, 1977). The decrease in long range order with temperature is shown for the 35% and 50%Co alloys in Figure 1.1.6. The shape of the curve is of the form:

-8- Figure 1.1.5: A comparison of isothermal ordering of FeCo alloys at 440°C, after quenching from 780°C. (Evmery et al. 1974)

-9- INTCNSn-Y a ) b) iue ..: h cag i eulbimodr (I S2)order = with 1.1.6: inequilibriumFigure The change temperature for: a) Fe50%Co and b) Fe35%Co binary alloys. Fe35%Cotemperaturebinaryb) for: and a) Fe50%Co (Oyedele andCollins 1977) 10 -1 S = Tc - TV

where S is the Bragg-Williams long range order parameter, LRO, Tc is the order/disorder transition temperature, and p is the critical exponent.

Theoretically it is possible to produce a LRO value of unity in the equiatomic binary but in the 35%Co binary, a value of 0.7 is the maximum; see Appendix 1. However, in practice there is disagreement as to the full order achievable. Smith and Rawlings (1976), as part of their discussion, reported the experimental values of S, from neutron diffraction experiments, for alloys with 1.8 to 2.0%V at temperatures in the range 450 to 500°C and the values varied from 0.80 to 0.87, while Ye.I. Mal'tsev et al. (1975), report values of S for Fe50%Col-5%V of 0.89 to 1.07 at room temperature. Lyaschenko et al. (1964), found values of S of 1 and 0.7 for Fe50%Co and Fe35%Co at room temperature.

1.1.4 Mechanism of ordering in undeformed alloys.

The mechanism by which the alloys order is still disputed. In 1964 Lyashenko concluded from the absence of neutron superlattice broadening that the reaction was homogeneous. In the same year, Chen, in his review shxtcd that 'ordering in FeCo is a marginal case of, either 2nd degree, which can be converted into a 1st degree transformation by a small addition of certain solute atoms such as vanadium, or it is a true phase change, showing a small volume decrease and an apparently unresolvable a + a' two-phase region'.

At this stage it is of the utmost importance to clarify the two methods by which long range ordering mechanisms have been studied:

Method 1. Here, a sample is quenched into iced brine and, if the cooling is rapid enough, (it must be 3000°C/sec according to Clegg (1971), and 6000°C/sec according to Kadykova and Selissky (1960) and Bardos et al. (1967)) a thin sample will retain full disorder. This sample is then annealed at a temperature below the order/disorder temperature for a

- 1 1 - period of time and the degree of order imparted is measured. This is by far the most popular method of study, but it does not simulate the conditions of an industrial quench where some ordering may occur during the quench at a temperature close to the order/disorder temperature.

Method 2. Samples are quenched at different rates and the degree of order in each sample measured as the sample is cooled or, more easily the LRO of the quenched sample is measured. Because of the rapidity of the quench, this method is rarely used. However, Chen used dilatometry to monitor the quench in 1961, and Bardos (1969), quenched at different rates and measured the saturation magnetisation.

The more recent work of Buckley (1975), Eymery et al. (1974), and Tahara et al. (1978), all use methods similar to Method 1.

Ordering occurs by a vacancy diffusion controlled mechanism. It appears that the vacancy concentration of the quenched sample depends upon the quench temperature (Girifalco 1964), and on the alloy composition. Ternaries such as vanadium are thought to increase the vacancy concentration, (e.g. Eymery 1974).

Both Tahara and Eymery agree that the low temperature isothermal ordering (350-550°C) occurs in 2 stages, but they differ on what the two stages are.

According to Tahara, in the quenched binary alloy the vacancy concentration is above that of the equilibrium concentration and so the sample initially orders rapidly. Eventually the lower equilibrium vacancy concentration is reached and the rate of ordering decreases. Lattice parameter measurements were used to monitor ordering and typical results are shown, with those of Clegg and Buckley (1973), in Figure 1.1.7.

Eymery (1974), used superlattice lines and X-ray peak broadening to measure domains and LRO in Fe49%Co2%V after isothermal anneals between 400 and 500°C. On quenching from above Tc, the sample becomes homogeneously short range ordered, within this matrix domains nucleate and grow until they impinge, whereupon the maximum degree of long range

- 1 2 - S/^rriax. ) Log(TIME/mtn») b) iue ..: steml reig uvs o F5%ouig the Fe50%Co forusing curves 1.1.7: Isothermal ordering Figure lattice parameterlatticemethod. (a:1978)and Tahara b: Clegg1973,Buckley 13- 3 -1 order is achieved. Domains then coalesce at a faster rate, Figure 1.1.8. The final domain structure in Fe49%Co2%V, according to English (1966), is 'swiss-cheese like', with only two domains; one domain represented by the cheese and the other by the interlinked holes.

Eymery's publication is one of the few reports that compares the ordering rates of different alloys and concludes that the fastest ordering occurs in Fe40%Co, then Fe49%Co2%V and finally Fe50%Co, Figure 1.1.5. Eymery also quenched Fe49%Co2%V from below Tc, and annealed. The quenched alloy was of S=0.25 with large domain size of about 200 nm and upon annealing, homogeneously ordered with little change in the domain size.

Buckley 1975, reported some surprising results at low ordering temperatures. Initially he agrees with Eymery that for isothermal heat treatments between 500 and 600°C a homogeneous 'short range' ordering occurs followed by coalescence of antiphase domains. This is for Fe50%Co, Fe50%Co0.4%Cr and Fe50%Co2.5%V. For the two alloys, Fe50%Co and Fe50%Co0.4%Cr, at temperatures as low as 260°C ordered regions precipitate at the grain boundaries of disordered grains. These areas develop into the grains and hence the material becomes ordered.

1.1.5 Ordering of cold worked material.

In general cold worked material orders at a slower rate than the simply quenched material at temperatures above 440°Cf Figure 1.1.9, (Smith & Rawlings (1976), Eymery (1974)). With these ternary alloys, the precipitation of a #-phase complicates the issue and precipitation is discussed later on. In a heavily deformed Fe50%Co0.4%Cr alloy Buckley observed recrystallised ordered grains forming at temperatures between 250 and 475°C with an activation energy of 105kJ/mole (Section 4.2). At temperatures above 500°C ordering occurs by a bulk diffusion process with an activation energy of 160kJ/mol. (Buckley, 1975).

1.1.6 Ordering Energies.

According to Buckley (1975), the ordering activation energy depends on the ordering mechanism, and hence the ordering temperature, as well as the degree of order already present in the sample. The majority of

-14- Figure 1.1.8: Isothermal domain growth for FeCo-2%V quenched from 780°C; the small arrows correspond to maximum LRO. (Eymery et al. 1974)

1 10 100 1000 t ( h l

Figure 1.1.9: The effect of cold work on isothermal orderina of FeCo-2%V at 440°C. (Eymery et al. 1974)

-1 5 - researchers have studied ordering above 500°C, which occurs by a vacancy controlled bulk diffusion mechanism, and requires enough energy for the formation and migration of a vacancy. It is therefore appropriate to compare the diffusion energies of Fe and Co atoms in ordered and disordered FeCo. Table 1.1.1 gives a summary of ordering activation energies of previous researchers and Table 1.1.2 gives values of the diffusion activation energies.

Table 1.1.1: Ordering activation energies.

Ef Em Eact Deformation (kJ/mole) Ye^/No

Fe50%Co 91 139 230 N Tahara et al. 1979 Fe50%Co2%V 87 183 260 N Eymery et al. 1974 Fe50at%Co 214 N Yokoyama et al. 1971 Fe33.3at%Co 255 N Yokoyama et al. 1971 Fe50%Co2%V 158-284 N Clegg et al. 1973 Fe50%Co0.4%Cr 105-200 both Buckley 1979 Fe50%Col.8%V 223-270 both Smith et al. 1976

Ef: Vacancy formation energy. Em: Vacancy migration energy. Eact = Ef + Em: Ordering activation energy.

-1 6 - Table 1.1.2: Diffusion activation energies.

Degree diffusion of order Eact Fe50%Co ~0 230 (Fe) Fishman et al. 1970 ~o 251 (Co) ~1 555 (Fe) ~1 555 (Co) Fe29%Co ~0 188 (Co) Hirano & Cohen 1972 Fe30%Co 7 227 (Fe) Fe50%Co z0 248 (Co) Fe50%Co 7 244 (Fe)

1.1.7 Disordering Energy.

The work on the 550 anomaly (Section 3) shows that a higher energy is required to disorder an ordered sample than vice-versa, (Sato 1976). This implies that ordering activation energy is a function of the degree of L.R.O., and is illustrated by the shape of isothermal ordering curves, (e.g. Figure 1.1.7) . This point is further illustrated by the higher values of the diffusion energy in the ordered samples.

1.1.8 Methods of measuring long range order.

Several methods have been used to measure L.R.O. Some of these are applicable to both Fe35%Co and Fe50%Co alloys, whereas others, such as lattice expansion, are only applicable to Fe50%Co alloys, because the effect is too small in Fe35%Co alloys. A "?" indicates that the technique has not been tried. Table 1.1.3 gives a summary of these techniques.

-17- Table 1.1.3: Techniques for Measuring L.R.O.

Suitable for: Technique 35%Co 50%Co Reference Y/N Y/N

Neutron superlattice Y Y Smith 1976 Mal'tsev 1975 Lattice expansion N Y Orrock 1986 Clegg & Buckley 1973 Dilatometry N Y Chen 1961 Zel'Dovich 1971 Co X-ray superlattice Y Y Persiano 1986

Specific Heat ? Y Clegg & Buckley 1973 Sato 1976 Resistance Y Y Tahara 1983 Persiano 1986 Yokoyama & Takezewa 1971 N.M.R. ? Y Muraoka 1976

Mossbauer Spectroscopy Y Y Persiano 1986 deMayo 1981 Saturation Magnetisation ? Y Clegg & Buckley 1973

-18- 1.2.0 The Atomic Structure of Binary Iron Cobalt Alloys

Most of this review is taken from the comprehensive and recent work of Hamdeh et Al. (1989). Clearly, in this fast moving area, atomic descriptions are quickly updated and refined, making the recent papers more reliable. In addition the atomic explanations proposed by the Californian group are in agreement with most previous investigations.

It has been established by Bozorth (1951), Pfeifer (1979), Shiga (1981), and Bardos (1969), that the mean magnetic moment of FeCo alloys varies with composition as shown in Figure 1.1.2, the maximum mean magnetic moment being at approximately 30-35at%Co.

The ferromagnetic properties in iron and cobalt are mainly due to the unbalanced spins of the 3d orbitals, the 3d up shell being fuller than the 3d down shell.

1.2.1 Variation of magnetic moments with iron to cobalt ratio.

Theoretical and neutron diffraction studies have shown that the increase in the mean magnetic moment with additions of cobalt is due to the increase in the moment on the Fe atoms; that of the Co atom stays approximately constant at 1.85pe. The Fe magnetic moment increases from 2.22jib for pure iron to 3.0;ib at 70 at %Cc; Figures 1.2.1 & 1.2.2. This increase at the Fe atom must be due mainly to changes in the 3d shell. Hamdeh et al. used Electro Energy Loss Spectroscopy (EELS), to measure the population change of the 3d shell at the Fe atom, and found that the number of 3d up electrons at the Fe atom increased to a maximum of 5 at 35at%Co while the number of 3d down electrons steadily decreased. This result, along with the population of the 4s shell, for disordered alloys is shown in Figure 1.2.3.

Marshall (1958) had previously suggested that local atomic magnetisation was proportional to the hyperfine field in ferromagnets, and this postulate does hold for many cases, e.g. the temperature dependence of magnetisation in iron, Nagle et al. (1960). However, Kasper et al. (1983), drew attention to the fact that in iron cobalt alloys, the local moment on the iron atom increases with cobalt additions to 3 pb at 70at%Co, (Figures 1.2.1 & 1.2.2) whereas the

-19- £ «*-O .*tr•ft 3c

CCD Eo

0 LL)

Figure 1.2.1: The mean magnetic moment of Fe atoms in FeCo alloys vs Co concentration of disordered and ordered alloys, assuming pco = 1.85u b . The solid lines are least squares fit to the experimental data of the disordered phase. (Hamdeh et al. 1989)

Figure 1.2.2: Magnetic moments attributed to 3d electrons in FeCo alloys vs Co concentration. (Collins and Forsyth 1963)

-20- 5.1

fl3dT

2.5

0.88 20 40 60 80 Co Concentration (at.°/o) Figure 1.2.3: The average number of electrons per atom in 3d up, 3d dovm and 4s states vs Co concentration for disordered FeCo alloys. {Hamdeh et al. 1989)

0

L_© o "5c 0 2 ra ©c L_ ©Q.

Figure 1.2.4: The average 57Fe Hyperfine Magnetic Field for disordered and ordered FeCo alloys at 77K vs Co concentration. (Hamdeh et al. 1989)

- 2 1 - hyperfine field at the iron nucleus varies with cobalt concentration in the same way as the bulk saturation magnetisation and decreases for alloys of more than 35% cobalt (Figures 1.2.4 5c 1.2.5).

1.2.2 Variation of the mean magnetic moment upon ordering.

Iron cobalt can adopt an ordered structure between 30 and 70at%Co. For the equiatomic composition there is a clear increase (of about 4%) in the saturation magnetisation for the ordered alloy, Shiga (1981) or Bardos (1969). (Figure 1.1.2). However for the compositions at the edge of the ordering range, the change in saturation becomes marginal and, with respect to this thesis, at 35%Co, ordering does not provide a significant improvement in the saturation magnetisation of the bulk material.

1.2.3 Variation of the iron hyperfine field upon ordering.

The Mossbauer studies of DeMayo (1981) , Montano (1977), and Persiano (1986), show that the hyperfine field at the iron nuclei decreases, (net increases like the saturation magnetisation), upon ordering, (Figures 1.2.4 & 1.2.5).

To understand these apparently conflicting results, the factors that determine the hyperfine field must be understood. Firstly, the electron orbitals, most influenbial in producing the field around the iron nucleus, are the spherically symmetrical s-orbitals. In iron there are the localised Is, 2s, & 3s shells and the valence, or non-localised 4s shell. There are other contributions which we shall neglect here. It is thought that the magnetic moment of the iron atom is proportional to the component of the hyperfine field produced by the Is, 2s 5c 3s core electrons; the core electron polarisation, CEP, (Akai 1984).

The final magnitude of the hyperfine field is determined by the 4s electrons' contribution. These electrons, like some of the 3d . electrons, are involved m bonding and hencear^ itinerant. Further more, the energy levels of the 3d and 4s electrons overlap and hybridisation occurs. Therefore, the contribution to the hyperfine field at the iron nucleus from the 4s electrons is strongly dependent on what is happening to the energy levels of the 3d electrons.

-22- -370 O DISORDERED • ORDERED o — EARLIER RESULTS (REF. 5) 2*:

COMPOSITION(°/0 COBALT)

Figure 1.2.5: Effective magnetic field (i.e. Hyperfine Magnetic Field) at the iron nuclei for disordered and ordered FeCo alloys vs Co concentration. (deMayo 1970)

Figure 1.2.6: Calculated local magnetic moments of d impurities in Fe (closed circles). The experimental values are given by triangles. (Akai et al. 1984)

-23- According to Kasper and Salahub (1983) , upon ordering, some of the anti-bonding, majority-spin molecular orbitals of the disordered system become filled, (leading to a lattice expansion). The spin dependence of the new energy levels of the 3d shell are such that hybridisation is less favourable. This leads to a change in the spin balance in the 4s shell and a corresponding decrease in the total hyperfine field at the iron nucleus.

Hamdeh et al. (1989), suggest that, upon ordering, the number of 3d electrons in the minority spin down state decreases at the same time as the number of minority spin down 4s electrons increases, (leading to a lattice expansion). The resulting spin polarisation of the 4s shell is responsible for the observed reduction in the magnitude of the hyperfine field. (Note that bonding states can be spin up or spin down, and the same applies to anti-bonding states).

A successful quantative explanation of the hyperfine field has not been derived because of the complex nature of the s-d hybridisation and the sensitivity to the iron atom’s environment.

1.2.4 Variation of the cobalt hyperfine field upon ordering.

Varying the number of neighbouring iron atoms by increasing the iron concentration has little effect on the cobalt magnetic moment, so it is concluded that ordering has little effect too. However, Muraoka et al. (1976), used N.M.R. to illustrate that the cobalt hyperfine field increases upon ordering. They also found that up to 3rd nearest neighbour interactions affected the shape of the N.H.R. trace. This long range influence also applies to the hyperfine field at the iron sites {Hamdeh et al. 1989 and Kasper & Salahub 1983).

1.2.5 Effect of Ternary Additions on Electronic Structure.

Only the atomic structure of Fe50%Co base alloys and the pure iron or cobalt base alloys have been studied with small percentage additions and this has only come about in recent years.

The purpose of the ternary^is to improve the mechanical strength or improve the resistivity of the bmaryr'without introducing deleterious

-2 4 - effects to the magnetic properties. The magnetic moments of atoms in BCC alloys act individually, (Crangle 1977), so ideally, any addition that can carry a magnetic moment should, align parallel to the surrounding atoms.

The experimental results of Chen (1962), and Akai (1984), are shown with the theoretical results of Akai, and Anisimov (1989), in Figures 1.2.6, 1.2.7 & 1.2.8. In addition, Cable and Hicks (1970), reported that manganese aligns anti-parallel when added to cobalt.

It can be seen that the only addition that increases the saturation magnetisation is manganese when it is substituted for iron in Fe50at%Co. In both pure iron and pure cobalt, the manganese impurities produce a negative contribution to the saturation. In addition to these 3d impurities, Akai investigated 4d additions to pure iron, all of which had a similar, but much reduced effect on the saturation (Figure 1.2.6).

-2 5 - Figure 1.2.7: Calculated local magnetic moments of d impurities in ordered Fe50%Co when impurity is substituted for Co atoms and when substituted for Fe atoms. The dotted line shows the experimental values. (Anisimov and Postnikov 1989)

Figure 1.2.8: Variation of the mean saturation moment per atom in ternary FeCo alloys with concentration for various solute elements. (Chen 1962)

-26- 1.3.0 The 550 peak in Fe Co binary alloys.

1.3.1 History

In 1943, Kaya and Sato discovered an inexplicable peak at 550°C in the specific heat versus temperature curve of binary, equiatomic iron cobalt. This was in addition to the order/disorder peak at 730°C (Figure 1.3.1).

The first report of the ‘550 effect' outside Japan was in 1949, by a Russian, Grum-Grzhimailo. The anomaly was also observed, using lattice parameter measurements, by Goman'kov et al. in 1962, again in 50:50 binary iron cobalt.

Reporting in 1971 the Japanese researchers, Yokoyama, Takezawa and Higashida used resistivity measurements to establish that the effect was not present if the alloy was heated slowly enough to maintain the equilibrium state. They did, however, report a very slight anomqly at 650°C, though to date, no other workers have confirmed this.

1.3.2 Formation of the peak.

The events that produce the anomaly are best understood by taking the example of specific heat, which is known to reflect changes in L.R.O. If an ordered sample of FeCo is heated at a rate of 10°/min, at about 400°C the equilibrium degree of order starts to decrease but there is not sufficient energy available for the sample to start to disorder and follow the equilibrium value. The sample moves further from equilibrium as the temperature continues to rise, until at about 500°C the energy is greater than the activation energy for disordering and the sample moves quickly towards equilibrium (Figure 1.3.2).

The other possibility is to start with a disordered sample. Again as the temperature is increased, the sample reaches a point at which it can move towards equilibrium'. However because the different components that make up the alloys state, such as short range order, long range order and vacancy concentration, have different relaxation times, the sample overshoots the equilibrium position before approaching from the other side, (Figure 1.3.3).

-27- Figure 1.3.1: Specific heat vs temperature upon constant heating rate (2°C/min) of various heat-treated Fe50%Co alloys. (Kaya and Sato 1943)

Temperature/C

Figure 1.3.2: FeCo alloy with a high degree of long range order approaching the equilibrium degree of order.

-28- This anomaly has been observed in Fe50%Co using resistivity, lattice parameter, neutron diffraction and specific heat techniques, and there are many mathematical models to describe the effect, Tahara et al. (1983), Sato and Kikuchi (1976) and Chamberod et al. (1972). In 1966, Asano used specific heat to observe the '550' peak in non equiatomic Fe-Co alloys as fully ordered samples adopted the equilibrium state of order, (Figure 1.3.4). In 1971 Matsuda found good correlation between his theory and Asano’s results.

1.3.3 Effect of precipitation in ternary alloys on the '550 Anomqly'.

Zel'dovich et al. (1971) performed many experiments on Co37.4%Fel0.5%V (Vicalloy 1) and conclude that ordering of the BCC. a phase occurred between 300 and 500°C and the precipitation of a gamma phase started at 420°C. This overlap of the ordering and precipitation processes is likely to occur in the alloys used in this work and the ordering specific heat peak may be affected by a precipitate peak. However, the compositions in this work involve a much lower ternary addition, typically <2%, so the effect of precipitation will be much less than that reported by Zel'dovich.

1.3.4 Application of the *550 effect' to ternary alloys.

A tentative use of the 550 peak was made by Glazyrina et al. in 1983, to study the effect of ternary additions to the equiatomic alloy. They confirm that the 550 peak is kinetic in nature, and go on to say that this peakakrfhuiJ' shifts in the same manner as the a/a' peak with addition of a ternary^to Fe50%Co. For example the addition of 2% vanadium lowers the a/a'peak by approximately 15°C and also lowers the 550 anomoly by approximately 5°C.

-2 9 - Temperaiure/'C

Figure 1.3.3: A disordered sample of FeCo approaching the equilibrium degree of order.

Figure 1.3.4: The 550°C anomgly observed by specific heat, for initially ordered FeCo alloys. (Asano et al. 1967) 30- 1.4.0 Recovery and recrystallisation in cold worked alloys.

1.4.1 Iron cobalt binary alloys.

There have been two comprehensive studies of the recrystallisation parameters of cold rolled iron cobalt binary alloys. The first was by Borodkina, Detalf and Selisskii in 1959, who used hardness measurements and the resolution of an X-ray doublet to determine the onset of recrystallisation. The second was by Gol'denburg and Selisskiy, in 1963, who used optical microscopy to observe the first signs of new grains appearing.

In the first study, cold rolled binary alloys with compositions from 21 to 76%Co were annealed at various temperatures for different periods of time. Here, only the results for the 35 and 50%Co alloys are reported in detail. The following temperatures and times are associated with changes in the shape of the (210) X-ray doublet for 80% deformed samples. The start of recovery is indicated by a decrease in the doublet's width, the end of recovery by a complete resolution of the doublet and the start of recrystallisation by the appearance of the two peaks of the doublet.

Fe35%Co Fe50%Co Beginning of recovery 500°C/30mins 350°C/lhrs End of recovery 700°C/15mins 700°C/30mins Start of recrystallisation 650°C/2hrs 700°C/15mins

In the 35%Co alloy the hardness increases from 250 to 350 Hv at 450°C, beyond which the hardness remained constant to 550°C and decreased above 600°C as recrystallisation proceeded, (Figure. 1.4.1). The Fe50%Co alloy was harder than the 35%Co alloy and the equiatomic alloy also reached a maximum at the low temperature of 300°C and general softening did not occur until 500°C, (Figure. 1.4.1).

According to Borodkina et al., this hardening is due to the formation of ordered areas at low temperatures and can proceed at very low temperatures in the equiatomic alloy because of the great reduction in the free energy on ordering. The authors go on to conclude that order establishment and recovery are interdependent and at low temperatures

-3 1 - Temperature, °C

Figure 1.4.1: Variation in hardness of 80% cold deformed FeCo alloys with temperature after 2hr annealing. (Borodkina et al. 1959)

Q, kcal/g.at. too

/ 1 1— 3 i i-'-'i-a 1 1 > n 1 ! ! 1 1 1 I I i 1 lg/J -8 • M !_ Jl i l l - ! _ l l 1 -

7 T 14 -10 -35 i IV r T 1 - Figure 1.4.2: Variation in the - 1 -18 recrystallisation parameters of FeCo binary alloys. 0 is the activation SO 1 InT. ~ r energy, A is the pre-exponential WS T r factor, T is the time elapsed v 1 1 7 2 7 1 until the beginning of 3 | i recrystallisation at 727°C, and TP is 2 the recrystallisation threshold / Y i !/ v A temperature. 07 (Gol'denburg and Selisskiy 1963) -/ r 1 1 J c -z 1 | 1 ? 1 1 i * ^ TP> 1 ■ / v. ■v V i °i \ 1 1 1 I / \ w 1 1 1 / if 1, 1 i f r f ? 'Ad 1/7 * I p'T 1 7, j* 1 ! 1 1 I iO 20 30 00 SO 80_ i 70 V 80 SO 700 °/o at. Co -32- can occur simultaneously.

However recrystallisation does not occur in either alloy until temperatures of 650 and 700° are reached in the 35 and 50% alloys respectively. Smoluchowski (1949), confirms that recrystallisation had taken place after annealing 96% cold rolled Fe35%Co at 700°C.

Gol'denburg studied the full composition range from pure iron to pure cobalt, annealing 83% cold rolled samples at 25°C intervals until new grains appeared.

He found that the variation in recrystallisation parameters with cobalt concentration was not monotonic. There were three peaks in the threshold temperature for recrystallisation corresponding to the ordered structures of FesCo, FeCo and FeC03. This and other recrystallisation parameters can be seen in Figure 1.4.2. Of particular relevance to this thesis is the comparison between Fe34%Co and Fe53%Co. He reported that the recrystallisation activation energies for Fe34%Co and Fe53%Co were 53 kcal/g.at (222 kJ/mole) and 98 kcal/g.at (410 kJ/mole) respectively. The corresponding threshold temperatures for recrystallisation were 610°C and 630°C.

He further suggested that the activation energy for recrystallisation was greater in the ordered structure than the disordered structure. To illustrate this he monitered the grainof Fe53%Co, (known to fully order) and Fe34%Co, (which cannot fully order) after isothermal annealing at temperatures from 700°C to 850°C, Figure 1.4.3.

Geisler et al. (1953), found a distinct change in the texture from the cold rolled condition after annealing 98% deformed Fe50%Co for 2 hours at 454°C, (Figure. 1.4.4). He terms this a recrystallisation texture and similar samples are said to have a fine grain size, though no microscopic evidence of this is provided. 454°C is certainly well below the recrystallisation temperature reported by other workers.

From the limited evidence it would appear that the Fe35%Co binary recrystallises at a lower temperature than the Fe50%Co binary. However, the greater the degree of cold deformation, the lower is the recrystallisation temperature, Table 1.4.1.

-33- b

/ 2 3 4 5 6 7 Q Annealing time, hr

Figure 1.4.3: Grain growth kinetics during a recrystallisation anneal of (a) Fe53at%Co, (b) Fe34at%Co. (Activation energies of recrystallisation are 411 and 220 kJ/mol. respectively). (Gol'denburg and Selisskiy 1963)

90* 60* 30* R.D. 30* 60* 90? Angle to Rolling Direction Angle to Rolling Direction

Figure 1.4.4: Anisotropy of cold rolled strip annealed for 2 hours at indicated temperatures. (Geisler et al. 1953)

-3 4 - Table 1.4.1: Recrystallisation threshold temperatures.

Deformation Recrystallisation temp. Fe50%Co Fe35%Co

80% 700°C 650°C 85% 630°C 610°C 98% 454°C

1.4.2 Recrystallisation in Iron Cobalt alloys with ternary additions.

Fiedler, in 1974, was optimizing magnetic and mechanical properties of FeCo2%V (2.0%V, 49.1%Co, 48.7%Fe) through controlling microstructure. Cold rolled samples (unknown deformation) annealed at 695°C for 2 hours were not recrystallised, whereas samples annealed at 710°C for 2 hours were fully recrystallised. Hence the recrystallisation threshold temperature is approximately 710°C, the same as the order/disorder temperature. However, it should be noticed that a gamma precipitate exists below this temperature (Section 1.5) which may inhibit recrystallisation.

Couto and Ferriera (1989), used T.E.M. to show that 90% cold rolled Fe50%Co2%V annealed at 600°C for 18 hours recrystallised but ^-precipitates were present. Chen (1961), used optical microscopy to illustrate recrystallisation of Fe50%Co2%V between 500 and 700°C though he gives no further experimental details.

Ashby et cal. (1978), reported that increasing the % deformation from 25 to 50%, lowers the recovery temperature.

Davies and Stoloff (1966),and Fielder and Davies (1970) both report recrystallisation in cold worked Fe50%Co2%V at temperatures below the order/disorder temperature. For Davies and Stoloff this was for 90% cold worked material annealed for 1 hour at 675°C.

In summary, the papers suggest that ternary vanadium alloys have normal recovery/recrystallisation characteristics, e.g. greater deformation

-35- leads to lower temperatures for recovery, recrystallisation temperatures in the range 600 to 710°C, and smaller grain sizes, (Burke et al. 1952). However, there is one report of anomalous recrystallisation behaviour of a ternary Fe50%Co alloy.

The paper of Buckley, 1979, reported the ordering and recrystallisation of Fe50%Co0.4%Cr from 260° to 570°C. He claimed that Fe50%Co0.4%Cr, cold worked by 50%, recrystallised to very fine (~lpm) ordered grains, at temperatures between 260 and 475°C. Normal recrystallisation occurred at temperatures above 650°C. To explain this, to say the least, unusually low recrystallisation temperature, he gave 3 criteria. i) the alloy should be ductile enough to allow substantial cold work. ii) the critical ordering temperature must be relatively high so that a high driving force for ordering exists over the temperature range described in iii. iii) within this temperature range the mobility of the order/disorder interface should be sufficient to allow the formation of new ordered grains whereas the- rate of ordering by the alternative bulk mechanism should be relatively slow.

Much of this work is complicated by the effects of precipitation and ordering which occurs over the same temperature range. For instance, Buckley mentions that Fe50%Co2%V does not recrystallise at low temperatures because the vanadium affects the mobility of the order/disorder interface.

The exact role of ^-precipitates in recrystallisation is not understood and is disputed. Both Davies and Stoloff (1966) and Pitt and Rawlings (1981) noted that the presence of ^-precipitates in Fe50%Co2%V can impede grain boundary migration, and hence lead to a smaller grain size, whereas Branson et al. (1980)/iritiaf-'that the small amount of X in Fe50%Co2%V is not enough to affect the recrystallised grain size. This is further discussed in Section 1.5.

-36- 1.5.0 Precipitation of ^-phase in ternary alloys.

It is well established by numerous workers (Pitt 1980, Kawahara 1983, Orrock 1986, and Couto et al. 1989), that a wide variety of precipitates can form during solidification and on annealing quenched Fe50%Co ternary alloys between about 400 and 600°C. The summary table of Persiano (1986) is shown in Table 1.5.1, along with the results of Kawahara (1983) Table 1.5.2. This precipitate is often » fine (

1.5.1 Precipitation in 2V-Permendur type alloys.

Taking 2V-Permendur as the most studied example, the phase diagram of Martin and Geisler in Figure 1.5.1, shows a ‘nose* of

Mal'tsev (1975), reports that the impurity vanadium atoms are surrounded by cobalt atoms, setting off a K-phase nucleation process. The concensus of opinion is that the lathe-like precipitate is of the form C 0 3 V (Kawahara, 1983), (though Fe is also reported in the precipitate (Fiedler et al., 1970)), which habits the {1101 planes of the ordered matrix, with the long axis in the <111> direction and has a

FCC structure with possibly an ordered LI2 type structure (Ashby and Rawlings 1977, Couto et al. 1989). The precipitate is paramagnetic and according to Ashby et al. (1977), in undeformed material nucleates around vanadium atoms that have segregated to anti-phase boundaries.

1.5.2 Additions of nickel to 2V-Permendur.

Pitt (1980) found that small additions of nickel assisted in the precipitation process (Figure 1.5.2). Pitt and Rawlings' (1981) comprehensive study of nickel additions to Fe50%Co2%V, reported rod shaped #-phase with an ordered LI2 structure aligned in the <111> direction of the ordered matrix. The dimensions and shape of the

-37- Table 1.5.1: Precipitation of second phase in FeCo base alloys (Summary of Persiano 1986)

REFERENCE MATERIAL INFORMATION

very small (

V rich particles; lath shape; FeCo2V Lis structure; in previously annealed material tend to precipitate on the grain under ormed boundaries and APBs after 10*s; in cold-worked material Ashby et al-iS78 or tend to precipitate in the grains and sub-grains after 10s s. A transport controlled co i d-uorkec mecnanisn is suggestec; TTT curve has "nose" at 550°C.

Ni ennances F: formation but ? ii t - i960 nc precipitate was oDservec F e C o V N i after siow cooling from 750“C

and if Ni content is 1 . 5 % cr iess.

(V=0.7 - L . S % ) TTT curve for underormea alloy has the "nose" at lower

Pitt h C N i = 1.3-7.AX) times and higher temperatures Rau1ings-1981 when compared to FeCoV. Small

grains are due to V 7 present.

If Co3X compounds are present Severa i the alloy is ductile. The reason is attributed to LCD ternary cones that, being deficient in Co, do not permit the Kauahara-1983aib FeCoX formation of B2 (ordered) volumes. Some Co3X effective a 1 1oys elements are: C,V,C r,N i,N b,M o Ta and U; ineffective are: B, (X — 0.01-«atx) A 1,Be,Cu.Au,Mn, T i,Si , and Zr.

r2 particles with diameter FeCoZVCu less than 0.5 pm and lattice FeCoZVU parameter 3.B R: cause grain Or rock-1985 FeCoN i refinement in FeCoVCu, FeCoVU FeCoS i h. FeCoNi; no m i cros tructura 1 (X=l,3.5 wtX) change was observed in FeCoSi.

-38- Table 1.5.2: The classification of elements on their effectiveness in imparting ductility to FeCo alloys. (Kawahara 1983)

A possible classification of compounds occurring between iron/cobalt and additive elements Type of Effective elements Ineffective elements compounds C V Cr Ni Nb Mo Ta W Be B Al Si Ti Mn Cu Zr Ag Au Diffusion-produced Co,C Co,V Co,Ci (Co,Ni) CojNb Co,Mo Co,Ta Co,W Co,Ti (Co,Mn) type CoV, CoCt CoNi, Co,Ta Fe,C FeV FeCr FeNi, FeNb Fe,W FeBe„ Fe,Al Fe,Si Fe-Mn FeZr, FeAl Fe,Si, Solidification- CoV CoNb CoMo Co,Ta Co,W4 CoBe Co,B CoAl Co,Si Co,Ti COnZr, produced type Co,Nb Co,Mo, CoTa Co,B Co,AI, Co,Si Co,Ti Co,Zr CoTa, Co.AI,, CoSi CoTi Co,Zr COjAl, CoSi, CoTi, CoZr CoZr, Fe,C Fe,Nb, FeMo Fc,Ta Fe.W, FeBe, Fe,B FeAl, FcSi Fe,Ti Fe,Zr Fe,Nb Fe-Mo FeBe, FeB Fe,AI, FeSi, FcTi Fe.Zr FeAl. Sigma phase. Temperature iue151 FCV trayscin (Martinand 1952)Geislersection. ternary 1.5.1:Figure FeCoV' r* Us>an4i. (jabb o M re* if i-oif Co->-pjui^ehe V, l*>si ■i '.) lc*uu> .'2.') &■ i -40- Figure 1.5.2: TTT curve for annealed Fe47%Col5%V3.5%Ni compared with that for FeCo2%V. (Pitt & Rawlings, 1981)

o quaternaries • quaternaries (forged) _ „ ...... H Fe-Co-V-Mn JO -/-4 /oNi £ ternary disordcrcd/aged ▲ ternary 2h at 750°C _

k2-6%Ni H ^ 1-8°/oNi 8 ~ 12 16 20 24 28 32 GRAIN SIZE, [im

Figure 1.5.3: Dependence of grain size on ^ precipitate. (Pitt & Rawlings, 1981)

-41- precipitate depended on annealing temperature and varied between a minimum width of 0.02 and a maximum length of 0.5 pm. The volume fraction of tf-phase was nearly 7% for Fe50%Co2%V3.5-7.4%Ni slowly cooled from 750°C, but no precipitate was present if the nickel content was reduced to <1.8%. The precipitate was rich in both vanadium and nickel, and it was suggested that the Ni substituted for the Fe or Co in precipitates of the type (Fe-Co)3V. The presence of tf-phase reduced the grain size of the annealed material and this effect is shown in Figure 1.5.3. Mal'tsev suggests that unlike vanadium, the nickel has a preferred affinity to iron atoms, and two nickel atoms in the first coordination sphere of an iron atom, cu-e. sufficient to nucleate the K-phase.

1.5.3 Additions of Ni and Nb to Fe50%Co and Fe35%Co alloys.

In this work, small percentages of Ni, Nb and Ta are added to Fe35%Co. Although no information is available regarding Fe35%Co alloys, there is literature reporting the results of Nb and Ni additions to Fe50%Co. Working with l-3%Nb additions, Persiano, (1986), reported the precipitation of clusters of spherical particles of approximate composition, 49at%Co, 35at%Fe and 15at%Nb after quenching and ageing for 1 hour at 630°C. Hence the precipitates are not purely CoaNb and do not only form at grain boundaries. In addition, the Nb-rich precipitate decreased the grain size of both the quenched and furnace cooled condition for the Fe50%Col-3%Nb alloys investigated.

In Orrock's work, additions of l%Ni or more produced a grain boundary precipitate corresponding to (FerCo)aNi, which again led to a reduction in grain size of the recrystallised material and of the quenched material for Ni additions of 5%.

1.5.4 Effect of cold work on precipitation.

Ashby and Rawlings 1977, studied the effect of 25% and 50% cold deformation on low temperature precipitation in an Fe-Co-2V alloy. They report firstly, that cold work accelerates the formation of & -phase (Figure 1.5.4), and secondly that precipitates form on dislocations and sub-grains whereas in the unworked samples, precipitation is homogeneous together with formation on anti-phase boundaries, (APB’s).

-42- Figure 1.5.4: TTT curves for annealed and cold worked FeCo2%V. (Ashby, Flower and Rawlings, 1977)

-43- 1.5.5 Consequences of precipitation. i) The material is magnetically harder due to pinning of magnetic domain walls (Carey and Isaac 1966) and a final heat treatment at 760°C is normally employed to redissolve the ^-phase, followed by a cooling rate such that precipitation is minimised and LRO is maximised. ii) The saturation magnetisation is reduced because iron and cobalt atoms that align ferromagnetically in the matrix align paramagnetically in the precipitate particles. This leads to a greater reduction in the magnetisation than would be expected from a simple magnetic dilution of the matrix by a ternary. iii) The material is mechanically harder after tempering between 400 and 600°C (Couto and Ferreira 1989). This is due to precipitates hindering dislocation movement. iv) Ductility can be imparted to the material directly if non-coherent precipitates act as obstacles to dislocations. promoting cross-slip and thus preventing grain boundary dislocation pile-ups. This mechanism aeoena: >n particle size ana spacing ana nas been illustrates with a thcria dispersion in tungsten (Hahn and Rosenfield, 1966). Alternatively, precipitates may lead to a small grain size by hindering primary grain growth or by increasing the nucleation rate of new grains (Pitt and Rawlings, 1981). This in turn leads to greater ductility by reducing the pile-up length. v) The electrical resistivity will be affected. Either the matrix, will become more dilute in vanadium leading to a reduced resistivity, or the increase in grain boundary area and/or the presence of a precipitate with a greater resistance, will increase resistance.

1.5.6 Precipitation in Fe35%Co Ternaries.

All of the results quoted are derived from Fe50%Co base alloys. Up to now, no work has been published on precipitation in Fe35%Cc bare ternaries.

-44- 1.6.0 Texture.

1.6.1 The Rolling Texture of BCC Alloys.

BCC metals have been widely studied (Hatherly), and it is accepted that the rolling textures for BCC alloys are essentially the same, being largely unaffected by alloy additions or the temperature of rolling. The most commonly observed textures in the as cold rolled condition are: i) UOOKllO) 12111<110> 1111}<110> 1111}<211> and in the recrystallised condition are: ii) 11111<110> lllli <211> 1110}<100>

Dillamore & Roberts (1965), have summarised previous workers results for BCC alloys. As far as cold rolled alleys are concerned, all have a predominant U00}<011> texture with the 12111<011> component always present and sometimes the other 2 textures listed above.

In the recrystallised samples the annealing temperature is influential in determining the relative strength of the three recrystallised texture components suggested above. These were general results for BCC alloys, and there was no specific reference to iron cobalt in the review by Dillamore & Roberts.

1.6.2 The Texture of Binary Fe35%Co.

Smoluchowski (1949), working on the recrystallisation of Fe35%Co in a magnetic field, reported the cold rolling texture as predominantly {100}<110>, i.e. the normal B.C.C. cold rolled texture. After a normal anneal at 700°C, the recrystallised texture was (1001<110), {100I15°<110>, and {111 115°<110>, (1949). He later reports the

-4 5 - recrystallised texture as i1001<110>, 1100l±15°<110> and i554i <225>, (Sawyer and Smoluchowski 1957). This is atypical for B.C.C. alloys.

1.6.3 The Texture of Binary Fe50%Co.

Geisler et al. carried out a more detailed analysis of texture in 1952. They experimented with equiatomic FeCo and found the rolling texture was variable and consisted of a spread of planes in the rolling plane but a common direction of <110> in the rolling direction. Recrystallisation was undertaken at a variety of temperatures, and the resulting textures were also variable. A <110> direction was present after all anneals below 650°C but, although the most detailed of the investigations reviewed,it was concluded that "uncontrollable variations would make questionable any attempted correlation with annealing variables."

1.6.4 The Texture of Fe50%Co2%V.

Pinnel et al. (1976), used some involved heat treatments to produce Fe50%Co2.8%V with recrystallised textures of either predominantly 1211i <110> and lllli <211> or predominantly 10011<110>.

ICawahara (1983), quenched Fe50%Co2%V from the ^-phase field (1100°C), and cold rolled 90% to give a texture of mainly i1001<110> and {2111<110> with some {111}<211>. The sample was reheated to 1050°C and quenched to give a recrystallised texture of {2111<110> with some {1001<110> and {111}<211>.

In 1986, Orrock investigated the textures of various Fe50%Co2%V alloys. The recrystallised texture, after an 850°C anneal, in heavily deformed samples, was predominantly f111l<211> while {1001<110> was the texture of the cold rolled condition. For lesser degrees of deformation, both the rolling and the recrystallised textures are mainly {211}<110>.

-4 6 - 1.6.5 Consequences of texture.

1.6.5.1 Magnetocrystalline anisotropy.

The potential energy of a single crystal sample of a ferromagnet depends on the direction with respect to the crystal axis, in which it is magnetized. The energy is lowest in an easy direction and hence a lower field needs to be applied in this direction to fully magnetize it. Figure 1.6.1 shows that the magnetization of a single crystal of BCC iron is easy in the <100> direction and hard in the <111> direction, whereas for FCC nickel the reverse is true. In cubic crystals, the easy and hard directions lie along principle crystallographic planes because, for 3d ferromagnets, the spins couple with the lattice via the spin-orbit and the orbit-lattice interactions.

The crystal anisotropy energy, E, of a cubic crystal may be written as:

E = Ko + Ki (di 2d2 2 + 0,2 2CX3 2 + O] 2Cr< 2 ) + K2 (di 2U2 2 0:i 2) + ....

where o1 ,0 2 ,os are the direction cosines of the magnetization direction with respect to the cubic axes of the crystal, and Ko. Ki

& K2 are constants. Ki and K2 are round using torque magnetometers, and Ki is normally considerably greater than K i. When known they can be used to determine the easy direction of magnetisation from a graph such as Figure 1.6.2. A more complete picture is shown in Figure 1.6.3 where the potential energy is shown for all directions for cubic crystals. The value of the magnetocrystalline energy .•is given by the radial distance of the surface from the origin. To find the complete order for magnetisation of the 3 main crystallographic directions in cubic crystals, a table is used, Table 1.6.1.

-47- Figure 1.6.1: The magnetisation of single crystals of iron and nickel in different crystallographic directions. (e.g.Cullity 1977)

Figure 1.6.2: Directions of easy magnetisation for cubic crystals as a function of the anisotropy constants Ki, and K2 . (Kalvius and Tebble 1979)

-48- magnetocrystallme drawings of the Figure 1.6.3: Perspective crystals. energy surfaces for cubic a) Ki +ve, K2 = 0 b) Ki -ve, K2 = 0 c) K2 = - 9Ki = -ve (Kalvius and Tebble 1979)

-49- Table 1.6.1: Directions of easiest, intermediate and hardest magnetisation in cubic crystals. (Kalvius and Tebble 1979)

Ki + + + - - -

+ oO - 9 K i /4 -9 K i - 0 0 9|Ki |/4 9|K, | Ki to to to to to to - 9 K i/ 4 -9 K i - 0 0 9 |Ki |/4 9|K, | + 0 0

Easiest <100> <100> <111> <111> <110> <110>

Intermediate <110> <111> <100> <110> <111> <100> Hardest <111> <110> <110> <100> <100> <111>

1.6.5.1.1 Values of Ki and K2 .

In 1960, Hall investigated the first anisotropy constant, Ki , of iron cobalt using single crystals, and his results, along with Chamberod et al. and KcKeehan, are shown in Figure l.G.4. The only values available for Ke, are those quoted by KcKeehan in 1937 and are from the work of Shih (1934). It is not clear if they are for ordered or disordered material. McKeehan's and Hall's results are shown in Table 1.6.2, along with the predicted easy, medium and difficult magnetisation directions.

-50- Figure 1.6.4: The variation of the first and second anisotropy constants Ki and K2 with composition in binary FeCo alloys.

Ki ordered Ki disordered McKeehan (1937) + Hall (1960) A Hall x Ki (unknown order) * Chamberod (1972) v Chamberod . K2 (unknown order)

-51- Table 1.6.2: Values of Ki & K 2 for binary alloys.

Fe:Co (McKeehan) ratio Ki K2 at 20°C easy medium hard (Jnr3)

100:0 42.1 15 70:30 10.21 16.2 65:35 7.4* 2.6* 100 110 111 60:40 4.54 -11.1 50:50 -6.80 -38.7 111 110 100 30:70 -43.30 5.3

(Hall) (McKeehan) 65:35 ORD 12 2.6* 100 110 111 DIS 15 2.6* 100 110 111 50:50 ORD +0 -38.7 100 110 111 ORD -0 -38.7 111 110 100 DIS -17 -38.7 111 110 100

* denotes extrapolated value. For all of McKeehan's results the degree of order is unknown.

If the more recent value for Ki of =0 Jnr3 for ordered Fe50%Co is taken, the energy required to magnetise a single crystal in different directions is only affected by K2 . When the direction cosines are calculated, K2 only affects the magnetisation in the <111) direction by^frfa. i.e.

E i 00 = Ko

Ei 10 = Ko

Ei 1 1 = Ko +/^YK2

The <111> does not feature strongly in the plane of the cold rolled and annealed sheet, and it is concluded from Hall's and McKeehan's results that there is little magnetic advantage in producing a textured material of eguiatomic composition. However, for the 35%Co binary, Ki

-52- is not close to zero and theoretically, a strongly textured sheet cut with the <100) preferentially aligned should show greater susceptibility than a sample from a non-textured sheet.

Chen (1964) also points out that the very small magnetocrystalline anisotropy in the Fe50%Co alloy means that a texture offers no magnetic advantage whereas in the Fe35%Co alloy the magnetocrystalline anisotropy is such that a textured material would theoretically offer improved susceptibility leading to lower losses in a.c. applications. The often quoted grain orientated material Fe3%Si, which is used in transformer cores, exemplifies the advantages that can be gained from a marked texture.

1.6.5.1.2. Magnetostriction.

The values of magnetostriction for binary alloys are shown in Figure 1.6.5. The values are lower in the Fe35%Co alloy but, in general, all the values are high and consequently internal stresses are associated: with the movement of 90° domain boundaries. Fortunately the number of these 90° boundaries can be reduced by a magnetic anneal as in the commercial alloy ‘Supermendur’.

In non-orientated polycrystallme material, adjacent grains of differing crystallographic orientation will produce internal stresses when the material is magnetised. These strains are reduced in both the 35 and the 50% cobalt alloys if a strong texture is present. In addition, a disimilar expansion of a crvstallographically coherent paramagnetic precipitate may produce stresses depending on the habit planes.

1.6.5.2. Microstructural effects due to texture.

The, discussion in 1.6.5.1. concerns single crystals from which parallels can be drawn for heavily textured materials. However, iron cobalt alloys are polycrystalline and often have a fine paramagnetic 2nd phase present which for Fe50%Co2%V type alloys, is semi coherent and, in the ordered matrix, is known to habit the 11101 planes. It is not known if the formation of the precipitate, in the anneal after cold rolling, assists in producing the commonly found <110> roll direction

-53- 2 0 0 x 10 ------O WATER QUENCHED * DISORDERED ------x 20°C/HR, ORDERED C

Figure 1.6.5: The saturation magnetostriction in the [100] and [111] crystallographic directions for FeCo allovs (Hall 1960) * *

-54- texture, or if the textured matrix causes any second phase to align in the roll direction, but in either case, the second phase will be orientated with the matrix. This will modify the domain motion and have an effect on magnetic anisotropy.

1.6.5.3. Variation of mechanical properties with direction.

Josso (1974) found that the mechanical properties of a textured cold rolled or annealed Fe50%Co2%V alloy, were dependent on the test direction relative to the roll direction. His results are shown in Table 1.6.3.

Table 1.6.3: Effect of test direction on mechanical properties

Tensile properties Anneal: 2hrs at 670°C Young's Modulus Tensile Elongation Cold Worked Annealed Direction Strength 92% 850°C/4hrs (xi6r G‘I07A W 1’)

Longitudinal 144 19 22 400 17 200 Diagonal 105 n 17 900 15 100 Transverse 136 12.5 24 600 21 300

He found that the predominant texture consisted of 2 components; !112}<110> + I100}<110>, and this resulted in poorer mechanical properties at 45° to the rolling direction.

-55- 2: Experimental details 2.1.0 Sample preparation.

Samples produced at Imperial College were made from electrolytic iron, from Gliddon, Ohio, electrolytic cobalt and high purity ternary elements from BDH Ltd.

2.1.1 Arc Melting.

The elements were cut into small pieces, and washed in acetone. The pieces were placed in the hearth of the arc melter and melted in the normal way. The atmo^liere was evacuated 3 times before alloying in a partial pressure of argon in the presence of a titanium getter which was melted before and after every melting. The button produced was remelted 3 times before melting in a finger ingot mould. It was normally neccessary to remelt the ingot 2 or 3 times to produce a smooth sided ingot. Hence a homogeneous ingot was produced. The finished ingots weighed between 70 and, more normally, 150 grammes, and had cross-sections of l-2cm2.

When producing yttrium or cerium ternaries, because of oxidation problems, a binary iron cobalt button was produced first into which a hole was drilled. The requisite mass of ternary was then inserted into the hole and the button was put back in the arc melter and alloyed as above.

2.1.2 Hot Rolling.

The ingot was manually filed to remove protrusions from the surface and any oxide was removed at the same time. The ingot was then heated in a furnace to 850°C and rolled from about 1cm to 2.5mm. It was normally only possible to make one pass before reheating, though occasional nimble handling made for two passes. When 2.5mm thick, the sample was reheated to 850°C and rapidly quenched to cold water to retain the disordered a-phase. Samples referred to as hot rolled were taken at this stage.

2.1.3 Cold Rolling.

The hot rolled sheet was cold rolled to approximately .7mm depending on

-56- the hardness of the material, before changing the rollers and further deforming to the minimum roller setting. The thinnest samples were between 0.10mm and 0.14mm again depending on the alloy's composition. Samples were taken at different stages of the rolling and hence a full range of thicknesses were available for the ensuing investigations.

2.1.4 Samples produced by industry.

Some of the alloys produced by Telcon Metals Ltd. were available as hot rolled, 2.5mm sheet, directly from industry. This was cold rolled at Imperial College in the manner already described. In a few cases cold rolled sheet was available from industry, and for the Rotelloy 5, with a 2 stage heat treatment in a hydrogen atmosphere having already been given.

2.1.5 Heat Treatments.

Hot rolled samples used in Impact Tests.

These bars were encapsulated in silica under a partial atmoshere of argon and suspended in a vertical furnace at the required temperature. At the end of the annealing time, the wire was cut, allowing the samples to fall direct to an iced brine bath, where the capsule was broken as quickly as possible with a blow from a blunt instrument. In this way disordered specimens with grain size determined by annealing time were produced.

Hot rolled samples used for X-ray work.

Here the objective was to produce samples differing only by the degree of long range order. Samples were initially disordered by annealing at 850°C for 1 hour in the manner described above. Individual samples were then suspended in a vertical furnace in a stream of argon or nitrogen gas for the shorter time periods or encapsulated in silica tubes under argon partial pressure for the longer anneal times, at a temperature below the order/disorder temperature, and they were then quenched //?& iced brine as outlined above.

-57 Cold rolled samples used for texture and hardness experiments.

All of these samples were encapsulated in an argon atmosphere. After the heat treatment, these samples were not quenched but left to cool to room temperature in their silica capsules, which took about 3 minutes. This method of cooling was thought unlikely to affect any texture, but would impart a high degree of long range order to the samples.

Cold rolled samples used for DSC references.

The texture samples that had been annealed for two hours at 760°C and cooled in their capsules were sufficiently ordered to be used as references for the DSC 550 anomaly experiments.

2.1.6. Compositions

A wide variety of compositions were investigated. A list of alloys and the experiments they were used in is shown in Table 2.1.1.

Table 2.1.1: Materials and the experiments in which they were used.

Tensile tests Impact tests

Fe35%CoO.2%Ta Fe35%Co0.2%Co 0.3%Ta 0.3%Ta 0.25%Nb Fe50%Co2%V Fe50%Co0.3%Nb 0.2%V0.4%Nb 0.2%Ta 0.5%Mo0.35%%Nb

Magnetic measurements Hardness measurements

Fe35%Co0.2%Ta Fe35%Co0.3%Ta 0.3%Tal%Ni 0.3%Tal%Ni 0.3%Ta4%Ni 0.3%Ta4%Ni Fe50%Col.8%V3.5%Ni Fe50%Co2%V 2%V 0.1%Ta Texture measurements 0.25%Ta 0.2%Nb Fe35%Co0.3%Ta 0.3%Tal%Ni Grain sizes 0.3%Ta4%Ni Fe50%Co2%V Fe35%Co0.2%Ta 1.8%V3.5%Ni 0.3%Ta 0.25%Nb X-ray measurements

0.3%Nb Fe35%Co0.2%Ta 0.4%Nb 0.3%Ta 0.35%Nb0.5%Mo 1.2Y 0.4%NbO.2%V Fe50%Co2%V 0.2%V0.4%Nb Thermal analysis 0.3%Nb 0.5%Mo0.35%Nb Alloys investigated by thermal analysis are given in Table 3.6.1, page 149.

-58- 2.2.0 Microscopy.

2.2.1 Light Microscopy.

As cast, hot rolled, cold rolled and annealed samples were observed by light microscopy. In all cases the sample was mounted in bakelite, and ground with consecutively finer silicon carbide papers, before a final polish with 3 and 1pm particle size diamond impregnated polishing wheels.

The samples were then etched with 2-4% nital. The etching time for pure hot rolled a-phase was found to be 2 to 5 times longer than the etching time for cold rolled samples with #-phase present; i.e. several minutes as opposed to z30 seconds. To measure grain sizes, the linear intercept method was used and for elongated crystals, samples were observed in both the linear and transverse directions.

2.2.2 Scanning Electron Microscopy, (S.E.M.).

The fracture surfaces of the impact tested bars were observed in the S.E.M. (Jeol JSM T220A) which afforded a depth of field not acheivable by light microscopy. No specific sample preparation was required, and bars were cut to size and held in position with mounting screws. For observation of precipitates, polished samples were first etched in nital. If the sample required mounting in bakelite due to size and geometry, the normal gold coating process was used to prevent charging.

2.2.3 Transmission Electron Microscopy, (T.E.M.).

The high magnification available with a Philips EM301 T.E.M. was used to locate fine precipitates in the samples cold rolled to 0.1mm and annealed. Because of the high magnetic saturation of the alloys, it was not possible to produce the small specimens required using the standard Struers sample thinning technique, so thin films were prepared using the window technique. Two electro-polishing solutions were used:

-59- In one, a solution of 20% perchloric acid in ethanol at -20 to -40°C with a current of 200mA at 40V was used. It was not always easy to find the plateau region of the electropolishing curve, but when conditions were right it provided a good polish.

The other solution was 10% perchloric acid in glacial acetic acid at room temperature at 200mA and 20V. The electrical settings and temperature were not too sensitive and an even thinning and shiny finish were more easily attainable. Unfortunately, carefull washing of the sample was required to prevent the polishing solution from corroding the sample over the following hours.

Samples of z2mm2 were mounted between copper grids and observed.

A hot stage was available on the E.M.I. high voltage T.E.M. and this was employed to watch the low temperature annealing process. Samples prepared as already described, were heated to 700°C over 1 hour.

-60- 2.3.0 X-Ray Diffraction.

2.3.1.1 Theory of lattice parameter and order.

It has been shown by Clegg and Buckley (1973) and by Orrock (1986), that the 0.2% lattice expansion that occurs in Fe50%Co base alloys upon ordering can be used to determine the Bragg-Williams long range order parameter, S. The method assumes a linear relationship between S and 5a0 , the change in the lattice parameter.

2.3.1.2 Lattice Parameter Measurements.

Hot rolled samples of both 50%Co and 35%Co alloys which had been quenched and isothermally annealed to impart differing degrees of order, were polished to 1pm and lightly electropolished in a manner similar to that in 2.2.3, to provide a stress-free surface.

A. Philips PW1710 x-ray diffractometer was used to measure the d-spacings, using an internal program that automatically scanned through the desired 28 range and labelled the intensity peaks. Cu Ka radiation was used and all samples were spun to reduce anisotropy effects. The lattice parameter, a, from the 5 X-ray peaks: 200, 220, 211, 310, & 222 were plotted against sin28 for each sample. The best fit straight line was computed, and the intercept at sin28 = 1 gave a best value for the lattice parameter, ao.

Cold rolled and annealed samples were scanned to look for extra peaks attributable to second phase particles. These samples were polished to lpm and also spun.

-61- 2.3.2.1 Theory of Superlattice Measurements.

Superlattice lines appear in X-ray scans when the 2 types of atoms that make up the ordered structure posses? a different atomic scattering factor, f. Hence, when there are alternating, evenly spaced planes of atoms of type A and type B of different scattering factors, fA and fb , incomplete destructive interference occurs, and the small signal that emerges is of intensity:

I a S* (fA - fB)2 where S is the Bragg-Williams long range order parameter.

This method of measuring degree of order is applicable to any diffraction technique and provides the best results when the difference between the scattering factors is greatest. For FeCo alloys, neutrons are most satisfactory because, being a nuclear scattering effect, a small change in the number of nucleons manifests itself as a large change in the atomic scattering factors, e.g. Bacon (1962).

the. Some preliminary neutron diffraction work was attempted at Silwood Park Reactor, but the advantage of greater superlattice peak size was out-weighed by the low neutron flux and limited time in which to set up a diffractometer of high precision, when compared to the highly tuned commercial X-ray facilities available.

For Cu Ka X-rays superlattice lines are barely visible for even fully ordered FeCo alloys. However, by choosing Co Ka radiation the atomic scattering factor of the cobalt atoms can be reduced and hence the intensity of the superlattice peaks increased. According to Persiano the ratio of the (100) superlattice peak to the (200) fundamental can be as high as 1:10.6 for furnace cooled Fe50%Col%Nb.

2.3.2.2 Superlattice Measurements.

The same Philips PW1710 as used for the lattice parameter measurements but with a cobalt tube was used for the superlattice measurements. The samples were produced in the same manner as for the lattice parameter measurements, and often the same sample was used. The diffractometer

-62- was stepped through the (100) superlattice peak, the detector pausing every 50 and counting for a set time, 5t. Typical scans were through 1 degree, pausing every 0.02 to 0.1 degree to count for between 200 and 800 seconds. The (200) fundamental peak was scanned in a similar manner as a reference.

2.3.3 Texture Measurements.

A Philips X-ray texture goniometer with Co Ka radiation was available to provide (110) and (200) pole figures. The machine worked in reflection mode and provided a 5° spiral out to 80°, when plotted as a stereographic projection. The data were corrected to allow for the intensity fall-off for angles greater than 65°, and the corrected intensity as a function of the angle of reflection, is shown for a powdered iron sample in Figure 2.3.1. Software was availaHe to convert the spiral to a grid and this information was plotted as a 3-D intensity relief map, in addition to the pole figure. Software was written to plot the raw data as a 3-D spiral surface in order to investigate the artificial peaks produced when converting the spiral data to a grid format. Figure 2.3.2. shows a relief map of the (110) poles and this illustrates how the small artificial peaks are produced.

A variety of cold rolled samples with different degrees of cold rolling and heat treatments at different temperatures and for different times were investigated. It is known that shapes with high aspect ratios produce erroneous results so samples were square or hexagonal and were polished to a 3pm finish.

63 Intensity (units) Intensity meter F.S.D.

200

100 ...... r = 7 7 *

0 0 10 20' 30 40 50 60 70 80 90 Anale from Normal

Figure 2.3.1: The variation in received intensity with angle from the sample normal, for iron powder.

-64- Figure 2.3.2: Spiral contour plots of the raw data from the texture goniometer, illustrating how a fictitious jagged ridge can be produced. (110) pole figure viewed: a) along roll direction b) perpendicular to roll direction c) at 45° to roll direction.

-65- 2.4.0 Differential Thermodynamic Techniques.

Many books are available describing D.S.C. and D.T.A. equipment; e.g. Thermal Analysis, by Wesley Wm. Wenddlandt, John Wiley and Sons, 1985.

2.4.1 Differential Scanning Calorimetry.

The differential scanning calorimeter used was a Stanton Redcroft DSC 1500. Samples could be ramped up or down at rates of up to 60°C/min. An inert atmosphere of nitrogen or argon was used at a flow rate of 60 mm/sec and the ratio of flow between the sample chamber and the surrounding furnace was 0.2. Stanton Redcroft software was employed for data analysis and graphical output.

2.4.1.1 Methods used in the D.S.C. study.

Method 1:

To determine the a/a'(order/disorder) and the aJft (alpha/gamma) transition temperatures, an alumina crucible containing the sample and one containing a platinum reference were run against each other. In all experiments, the samples were cleaned with alcohol and one surface was polished to ensure good thermal contact. The reference was one or two discs of platinum, such that the mass was close to that of the sample.

Method 2:

To determine differences between two samples with different thermal histories, great care was taken to match sample geometries and masses. In general for samples of weight 15.0mg, the difference would be <0.2mg. This method was used to observe the '550 anomqly' and the initial baseline of the run and the degree of cancellation of the a/2f peaks gave an indication of the successfullness of the matching. Runs were frequently repeated until satisfactory cancellation was achieved. To increase sensitivity, samples were put directly onto the thin alumina washers and not in alumina crucibles.

-66- 2.4.1.2 D.S.C. Experiments.

1) In early experiments on hot rolled alloys, using Method 1, the sample and reference were matched to within a few milligrams. The purpose of the runs was to determine the a/tf and the order/disorder transition temperature for both Fe35%Co and Fe50%Co base alloys with different ternary additions. The samples were ramped to 1100°C and back to below the order/disorder temperature, at a rate of 20°C. The average of the ’up* and the ’down’ peak for both transitions was used as the transition temperature.

2) The same as experiment 1 only the effect of ramp rate on average peak temperature was determined.

3) Method 2 was used on totally identical samples to check cancellation of peaks.

4) To observe the peak due to oxidation, Method 2 was used on identical annealed and slowly cooled samples, one. cfodiiJi had had the oxide layer removed.

5) The samples used in the '550 anomaly’ work were heavily cold rolled, ensuring that they were disordered. The reference was annealed at 760°C for 2 hours and cooled in its silica tube to room temperature over 3 minutes. These samples were cut from the sheet used in the texture investigations. Both samples were polished on 3 pm grit polishing wheels. Initially samples were ramped to 1100°C, back down to 350°C and then up to 850°C to produce a baseline. It became apparent that close matching of samples could be assessed from the shape of the curve and so only the first ramping stage was used in the later runs. This experiment was repeated using different ramp rates.

2.4.2. Differential Thermal Analysis (D.T.A.) and Thermo-Gravimetric Analysis (T.G.A.) Methods

A Stanton Redcroft STA 781 differential thermal analysis and thermal gravimetric analysis machine was available. Transition temperatures were determined in the same way as for D.S.C. and oxidation was investigated using T.G.A. In these experiments a ramp rate of 10°C/min was used.

-67 2.5.0 Mechanical Testing.

2.5.1 Impact Tests to investigate industrial problems with cold rolling.

Some alloys of Fe35%CoX coda not^rolled at room temperature but warm rolled at about 100°C. It was thought that the lack of rollability was associated with a ductile to brittle transition similar to that in steel.

To investigate this hypothesis, bars of dimensions 43.3 x 9.0 x 2.5 mm with a 2mm deep, 30° V-notch were tested with a Charpy Impact Tester. The samples were heat treated at 850°C to produce single phase material of differing grain size. These samples were then impact tested at temperatures between -20 and 300°C. Sample temperature was determined by warming both the sample and a thermocouple probe in the same furnace, then quickly transferring the sample to the grips of the impact tester with warmed tongs and touching the thermocouple against the sample until the temperature was dropping at a steady rate. The time from furnace to test was close to a minute. As a check on accuracy, a bar with a thermocouple permanently secured to it was compared with the 'touch on' thermocouple method and good agreement was achieved. It is difficult to estimate the error in sample temperature because of two effects. i) localised cooling around the fracture area from the sample grips, suggests that the fracture area is cooler than the probe might suggest. ii) the interior of the specimen is warmer than the surface reading from the probe.

Both of these sources of error are systematic and because of their opposing nature it is concluded that sample temperatures are accurate to at least ±5°C. Flaws in individual specimens are a much greater source of inaccuracy.

-68- 2.5.2.1 Tensile Tests.

Hot rolled 2.5mm sheets were rolled parallel and and others cross rolled down to 0.5mm, as shown in Table 2.5.1.

Table 2.1: Tensile samples-

Alloy. Direction of cold roll.

Fe35%Co 0.2%Ta cross rolled 0.3%Ta cross rolled 0.25%Nb parallel rolled

Fe50%Co 0.3%Nb parallel rolled 0.3%Nb cross rolled 0.2%Ta parallel rolled

Waisted tensile specimens of gauge length 20mm, width 7.5mm, and thickness 0.5mm were stamped in the cold roll direction from the cold rolled sheet.

Heat treatments were designed to produce equiaxed isotropic grains in the ordered and disordered state. Two sets of twelve samples were heated at 850°C for 2 hours under an argon atmosphere in silicon capsules before one set was quenched and the other furnace cooled. Prior to testing on the 'Nene' tensile tester, the edges of the samples were manually smoothed using 1pm silicon carbide paper.

-69- 2.5.2.2 Test conditions.

A 'Nene 3000' universal testing machine with computer control and data collection was used for tensile tests to fracture. The machine parameters were:

Load cell...... 30kN Test speed...... lmm/min and all but 2 tests were repeated.

2.5.3 Hardness and microhardness tests.

Both micro and macrohardness tests were employed with a diamond indenter and a load of 10 kg for macro-hardness and 100 to 200g for micro-hardness. In both cases samples were polished to 3pm and between 5 and 10 readings were taken, depending on the scatter.

-70- 2.6.0 Magnetic Measurements.

2.6.1 Saturation Magnetisation using the Sucksaith Balance.

The Sucksmith Balance is a piece of equipment that measures saturation magnetisation using a force method. The sample is suspended on a rod between the shaped pole pieces of a magnet. The magnet provides a field big enough to saturate the sample in the horizontal plane and with a linear gradient in the vertical plane. The sample is attracted to the strongest part of the field which is arranged to be downwards, and the force with which the sample is pulled down, is proportional to the magnetisation. The top end of the sample rod is conected to a balance which gives a reading proportional to the magnetisation, (B+H).

In these experiments the magnet was set at a strength known to saturate the sample, and this was kept constant for all measurements. Before each iron cobalt alloy was measured, an iron reference of similar weight was used to confirm the correct calibration. Samples were z0.5x3x4mm with a central hole for connection to the sample rod.

Samples were cut from hot rolled sheet, which was known to be recrystallised and only mildly textured.

Errors.

The main source of error was differences in sample size. Although every effort was made to make samples as near to identical as possible. In addition, sample positioning was not always identical and problems were encountered with the rod pulling to one of the pole pieces. Because of these geometry, and later to be realised*texture effects, the errors in this method were about 5%.

2.6.2 The Vibrating Sample Magnetometer, V.S.M.

General information on the V.S.M. can be obtained from Cullity (1972), and information specific to the machine used, from the PhD of Duffield (1985).

-71- In this method, the sample is vibrated vertically in a steady magnetic field. In close proximity to the sample are oppositely wound pick-up coils in which an alternating current is induced, the magnitude of which is proportional to the magnetisation, M, of the sample. The strength of the applied magnetic field, H, can be varied by controlling the current flow through the magnet, and in this manner an M-H hysteresis loop can be obtained.

A long thin sample is mounted as shown in Figure 2.6.1. It can be seen that the magnetic moment is measured in the vertical direction, and hence by cutting samples from the textured cold rolled sheet, in different directions, and aligning them with the long axis lying vertically, it is possible to determine the magnetic properties in different crystallographic directions.

When measuring the magnetisation of strong ferromagnets, it is important to consider demagnetisation effects; i.e. the sample distorts the field in a manner that draws 'lines of flux' into the sample, reducing the field immediately adjacent. Hence the value of H is less outside the sample than it would be if the sample was not present at all. Demagnetisation effects are strongly dependent on sample size and geometry, and so samples were matched as closely as possible in this study, enabling comparisons between samples to be made. In addition a 'look-up' table was available to correct magnetisations due to differences in sample length. Demagnetisation most strongly affects susceptibility measurements and hence the results quoted here for coercivity, He, are not absolute.

Demagnetisation is not a problem as far as saturation magnetisation is concerned as the exact field that causes saturation does not need to be known.

Measurements normalised to mass, were taken at temperatures between -10° and 20°C rather than close to absolute zero, so as to simulate more closely the working temperatures of the alloys.

-72- Steel Tube to Vibrator Sample Thermometer Position

Anti-cryostat Araldite Rod Superconducting \ Solenoid \

P ick-u p C oils

Figure 2.6.1: Schematic diagram of the sample chamber of the vibrating sample magnetometer.

-73- Sample preparation

35%Co alloys with nickel additions were available in .13mm cold rolled sheet. They were encapsulated in argon in silica tubes and annealed for 2 hours at temperatures of up to 760°C. The samples were removed from the furnace and allowed to cool over about 3 minutes to room temperature. A degree of LRO was expected in both heat treatments whereas the cold rolled samples were disordered and internally stressed.

In addition Rotelloy 5 was available in the cold rolled and the 'soaked' (see Table 2.6.1) and annealed for two hours at 760°C in hydrogen conditions direct from industry. These samples were .1mm thick and the annealed samples were known to be extremely textured. Details of the samples are shown in Table 2.6.1

Sample geometry

Because of the problems of demagnetisation already mentioned, it was neccessary to ensure that samples of the same composition were of the same geometry and mass. Bars with dimensions of 5.5 x 0.13 x 0.6 mm for the Fe35%Co base alloys and 6.5 x 0.1 x 0.6mm for the Rotelloy 5 samples were chosen. To produce these straight bars, without curling or twisting, sheets of material were stuck to an aluminium substrate with 'Evostick', and guillotined to the required dimensions, then un-stuck with acetone.

-74- Table 2.6.1: Details of V.S.M* Samples.

Sample Temp, of Bar Bar Bar 2hr anneal length Thickness Mass (°C) (mm) (mm) (mg)

RD 45° RD 45°

Fe35%CoO.2%Ta 670 6.66 6.68 0.1 2.20 2.50

Fe35%CoO.3%Tal%Ni CR 6.54 6.62 0.13 2.90 3.05 670 6.38 6.62 3.40 3.10 760 6.20 6.48 3.40 3.00

Fe35%Co0.3%Ta4%Ni CR 6.58 6.54 0.13 4.05 3.65 575 6.45 6.71 3.40 3.60 670 6.36 6.60 3.25 3.15 760 6.61 6.60 3.75 3.60

Rotellov 5 CR 6.50 6.40 0.10 2.60 2.40 760* 6.45 6.38 2.40 2.20 6.62 6.38 2.40 2.30

* denotes sample was 'soaked' at 450°/1.2hrsf raised to 760°C/2hrs at 2°/min, and cooled at 1.7°/min to room temperature.

2.6.3. Permeammeter Measurements.

A 'Fahy Simplex1 type permeammmeter similar to that described by Cullity was used. Figure 2.6.2. shows a simplified diagram of the equipment.

-75- Meter to measure H

Figure 2.6.2: Block diagram of the permeammeter.

-76- 2.6.3.1 Sample preparation and geometry.

A sheet of Rotelloy 5, a commercial alloy supplied by Telcon Metals Ltd., in the cold rolled by 96% condition, was cut into strips 250x2.5x0.1mm with the long axis in either the roll direction or at 45° to the roll direction. The samples were then heat treated in hydrogen as shown in Table 2.6.2.

Table 2.6.2: Samples for permeameter measurements.

Sample Heat treatment

Rotelloy 5 45 (S) 150°C/min to 450°C/72 mins Rotelloy 5 RD (S) 2°C/min to 760°C/120 mins 1.7°C/min to 25°C

Rotelloy 5 45 As above only no anneal time at 450°C Rotelloy 5 RD

4 or 5 strips were lightly bound in the centre with mashing tape, and over this 20 turns were wound and connected via twisted wires to a fluxmeter. The cross section of the strips was calculated from the density, mass and sample lengths, and the distance between the clamps was measured and was about 10 cm. The sample was magnetised by a field measured by the ajacent Hall probe, and the poles reversed by a manual switch. The flux change was recorded for different magnetisations to saturation. The remanence was calculated after removing the field, and the coercive force was calculated by applying a reverse field to reduce the flux density to zero. A simple calculation of the form:

B = 0 /(2 x Ns x A) was used to calculate the flux density, B. 0 is the flux change, Ns is the number of turns around the strips, and A is the cross sectional area. No air corrections were made and hence the magnetisation values are slightly larger than expected but the results were quite accurate enough for the purpose of comparison between samples.

-77- 3: Results 3.1.0 Grain size and mechanical properties.

3.1.1 Changes in grain size on annealing.

Figures 3.1.1 and 3.1.2 show the change in grain size determined by the linear intercept method, of Fe50%Co0.35%Nb0.5%Mo and Fe35%Co0.2%Ta. The nucleation and growth of the FCC phase is indicated by the dip in the grain size between 900 and 1000°C. It should be pointed out that the grain growth in the 35%Co base alloy is much greater than that in the 50%Co alloy. The maximum grain sizes are approximately 150pm and 40pm for the 35%Co and 50%Co base alloy respectively.

3.1.2 Grain size and cold rollability.

3.1.2.1 Rollability of different alloys with different grain sizes.

Figures 3.1.3 and 3.1.4 shows the hot rolled grain size of 50%Co and 35%Co base alloys and indicate . whether the alloys were successfully cold or warm rolled after quenching from 850°C. A warm roll is performed at about 100°C. The Fe50%Co base disordered material with a grain size of <25pm cold rolls. Hot rolled sheet with a grain size of 28um warm rolls, whilst sheet with a grain size greater than 30pm cracks upon warm rolling. For the Fe35%Co alloys less information is available and the conclusion is less clear cut. However, the Fe35%Co0.3%Ta sheet has the largest grain size and does not cold roll, but does warm roll. It must be concluded that grain size and ternary additions both have an effect on rollability.

3.1.2.2 Rollability of Fe35%Co0.2%Ta samples with different grain sizes

Figure 3.1.5.(a) shows samples from impact test experiments that were cold rolled. It can be seen that the most successful cold roll was for the material with a grain size of 22pm and the least successful for the sample with a grain size of 40.3pm. All of these samples were put through the rollers until the minimum setting was reached. The thickness of each sample was then measured and this information is shown in Figure 3.1.5. (b). Hence, hot rolled and quenched Fe35%Co0.2%Ta cold rolls more readily and with less cracking if the grain size is kept to 22pm or less.

-78- dtameter (microns) 50

40

30

20

10 alpha mar tenslte

0 i i I 850 900 950 1000

Temp, of heat treatment (°C) Figure 3.1.1: Grain size of Fe50%Co0.35%Nb0.5%Mo after 2 hour heat treatment Grain diameter (microns) 2 0 0 ------

150 -

100 -

50 -

alpha martensite

700 800 0 0 0 1000 1100 1200 Temp, of heat treatment (°G) Figure 3.1.2: Grain size of Fe35%Go0.2%Ta after 2 hour heat treatments Qtuhdlij haf jxiUtl

80- Ternary addition (% wt)

0.2Nb

0.3Nb

0.4Nb

0.35Nb 0.5MO

0.4Nb 0.2V

i------1------1------1—:— i------r 0 5 10 15 20 25 30 35

Grain size (microns)

KSXI cold rolled E23 warm rolled EB3 didn't roll

Figure 3.1.3: Grain size of hot rolled Fe50&Go base alloys

-81- Grain size (microns)

ES3 cold rolled E5S3 warm rolled

Figure 3.1.4: Grain size of hot rolled Fe35%Co base alloys

-82- °)

Fe357.CoO-27.Ta

d=22-0}jm

KfWf “^ T

d=25'6jjm

d=35-3jum

d=A-0-3jjm 1cm 1— •

Max deformation (%) 100 r— ------

90 -

Cotcl Figure 3.1.5: Effect of grain size on/reliability of hot rolled and quenched Fe35%CoO.2%Ta. a) illustrates that samples with a greater grain size crack more under cold rolling, and b) illustrates how the final sheet thickness of the same samples depends on the grain size.

- 83- 3.1.3 Impact tests on hot rolled sheet.

3.1.3.1 Existence of a Ductile to Brittle Transition Temperature, (DBTT).

Figure 3.1.6 shows that a DBTT occurs in disordered Fe50%Co2%V at 75°C for a grain size of 21pm. Also shown is the result for ordered Fe50%Co2%V which shows that the ordered material remains brittle up to 300°C.

3.1.3.2 Dependence of DBTT on grain size.

Figures 3.1.7 to 3.1.10 show the DBTT for the two Fe50%Co base and the two Fe35%Co base alloys, namely Fe50%Co0.35%Nb0.5%Mo, Fe50%Co0.2%V0.4%Nb, Fe35%Co0.2%Ta and Fe35%Co0.3%Ta, with varying grain sizes. For a given composition the transition temperature increases as the grain size increases, e.g. in Fe50%Co0.2%V0.4%Nb, a change in grain size from 21.0 to 34.6pm, shifts the DBTT by 150°C. However, ternary addition also determines the transition temperature. For alloys with a grain size of about 21pm, Fe50%Co0.35%Nb0.5%Mo (20pm) is brittle at 50°C, Fe50%Co0.2%V0.4%Nb (21pm) is ductile at 50°C, and Fe35%Co0.2%Ta (22pm) is brittle at 50°C. The variation of DBTT with grain size is summarised in Figure 3.1.11.

3.1.3.3 Impact tests at higher temperatures.

A Fe35%Co0.3%Ta sample with the largest grain size of 67.6pm was tested, and it was found that to reach the estimated required impact test temperature of greater than 280°C to exceed the DBTT, a furnace temperature of >420°C was neccessary, Table 3.1.1.

-84- Energy Absorbed (J)

( Chtuf.j ) Figure 3.1.6: Impact test results^for Fe50%Co2%V (21pm).

-85- Figure 3.1.7: Impact test results for Fe50%Co0.2%V0.4%Nb.

Figure 3.1.8: Impact test results for Fe50%Co0.35%Nb0.5%Mo.

-86- Energy Absorbed (J).

Energy absorbed (J)

-87- DBTT (Kelvin).

Grain diameter (microns).

Figure 3.1.VU The grain size dependence of the DBTT

- 8 8 - Table 3.1.1: Impact test data for Fe35%Co0.3%Ta (67.6pm).

Furnace Temp. Test Temp. Time from furnace Energy (°C) (°C) to test (secs) (J)

425 280 35 4.6 420 305 24 5.0 • 508 340 20 2.0, 540 350 21 1.8' 503 >350 ~5 3.1

The samples were in the furnace for up to 1 hour, and in this case, the samples were found to become more brittle as the test temperature was increased, Figure 3.1.10.

3.1.3.4 S.E.M. of fracture surfaces of impact test specimens.

S.E.H. pictures of the fracture surfaces are shown in Figure 3.1.12. For all the samples,(xaipt the Fe35%Co0.3%Ta (67.6pm) sample, the brittle fracture occurs predominantly by transgranular cleavage (a,b) whereas the micrographs of the ductile fracture surface (c,d) show a typical deformed and dimpled surface. Pictures e and f show that intergranular cleavage occurs in the Fe35%Co0.3%Ta (67.6pm) samples that broke in a brittle manner.

-89- Figure 3.1.12: S.E.M. micrographs of typical fracture surfaces.

(a): Fe50%Co0.2%V0.4%Nb (29.2pm) impact tested at 75°C, showing intragranular cleavage typical of brittle fracture below the DBTT.

(b): Fe35%CoO.3%Ta (44.2pm) impact tested at 125°C. An area in the centre of the specimen showing intragranular cleavage.

-90- (c): Fe35%CoO.3%Ta (44.2um) impact tested at 150°C showing a typical ductile fracture surface.

(d): A close-up of the fracture surface of Fe35%CoO.3%Ta (44.2pm) impact tested at 150°C.

-91- (e): The centra of the fracture surface of Fe35%CoO.3%Ta (67.6pm) after impact testing at 280°C, showing intergranular fracture.

(f): An area of the Fe35%CoO.3%Ta (67.7pm) specimen close to the notch, showing both inter and intragranular fracture. (Tested at 280°C).

-92- 3.2.0 Hardness results.

3.2.1 Dependence of hardness on annealing temperature.

Figures 3.2.1 and 3.2.2 show the macro and microhardness values for 92%cold rolled Fe50%Co2%V and Fe35%Co0.2%Ta after they were annealed for 2 hours at the temperatures indicated. The macrohardness values are lower than the microhardness values and is simply due to systematic errors from the thinness of the samples. The microhardness values should be taken as the true values.

The Fe50%Co base alloy is consistently 5 to 20% harder than the Fe35%Co base alloy. Both alloys show a hardening peak, which is at 520°C for the 50% alloy and at a slightly higher 550°C for the 35% alloy. The magnitude of this hardening is 40 to 45% for the 50%Co and 25 to 35% for the 35%Co alloy.

Hardening has already started at 350°C in the 50%Co alloy whereas the 35%Co alloy does not start to harden until above 400°C. There is a suggestion of softening at 400°C in the 35% alloy reminiscent of the low temperature recrystallisation reported by Buckley (Section 1), but this may be due to experimental error.

The 50%Co alloy returns to a hardness value below that of the cold rolled alloy at 620° and at 710°C according to macro and raicrohardness results respectively. For the 35% alloy these values are 580°C and 690°C. On averaging, this gives peak widths of 315°C for the Fe50%Co base alloy and 235°C for the Fe35%Co base alloy.

Of the anneals investigated in this work the softest condition was produced after annealing at 750°C for 2hrs and corresponds to VHN’s of

291 and 261 for the Fe50%Co and Fe35%Co base alloys^ 5. 3.2.2 *) .

3.2.2.1 Effect of composition on hardening.

In addition, the same experiment was performed on 95% cold rolled Fe35%Co0.3%Ta with additions of 1 to 4% Ni. The same trends are observed with the greater Ni content resulting in a harder material. Figure 3.2.3 shows the results and the widths of the hardness peaks are

-93- Anneal temperature l°C) Figure 3.2.1: Variation of macrohardness with annealing

-94- Anneal temperature (°C1

Figure 3 .2.2: Variation in microhardness with annealing

-95- Anneal temperature (°C) Figure 3.2.3** Microhardness of Fe35%Co0.3%Ta with l~4%lsli. All samples 95% cold rolled and annealed for 2 hours

-96- approximately 250 and 300°C for the l%Ni and the 4%Ni respectively, slightly wider than the Fe35%Co0.3%Ta peak.

3.2.2.2 Effect of prior cold work on hardening.

Figure 3.2.4 shows selected hardness values for material undergoing the same 2 hour anneals but after differing cold deformation. The same trend in hardness variation with anneal temperature to that for the 92% i$ ohs&rvsul.

3.2.3 Hardness change upon ordering.

The 92% cold rolled and annealed at a-j-tes 800°C/2hrs samples of Fe35%Co0.3%Ta and Fe50%Co2%V that were highly ordered,^ were heated for a further 5 minutes at 800°C and quenched to produce disordered recrystallised specimens. Hardness values were taken to see if ordering could account for the observed hardness peak. Table 3.2.1 shows these macrohardness values and those of the recrystallised disordered sample for the alloys investigated in Section 3.2.1. The change in hardness due to ordering is considerably less than the low temperature hardness peak.

Table 3.2.1: Effect of ordering on macrohardness.

Condition Fe35%Co0.3%Ta Fe50%Co2%V VHM (lici/niwx) y tttv

Recrystallised disordered 174- 260 Recrystallised ordered '2.03 24-0 Peak hardness 373 466

3.2.4 Variation in hardness on isothermal annealing.

3.2.4.1 Annealing at 550°C.

The macro and microhardness values of the same 92% cold rolled Fe50%Co2%V and Fe35%Co0.3%Ta alloys after annealing at 550°C for up to 6 hours are shown in Figures 3.2.5 and 3.2.6. Samples with 96% cold deformation were macrohardness tested^for periods of <1 hour and show no anomolous behaviour, Figure 3.2.7.

-97- Anneal Temperature (°C1 Fig 3.2,4* Effect of deformation on macrohardness

-98- VHN (kg/mm1] 500

100 Cold rolled 92% FeS0%Co2%V 0 Fe3S%Co0.3%Ta o 1- 0 1 2 3 4 5 6 7 Anneal time (hours! riqure 3.2.5s Variation in macrohardness with isothermal annealing at 550°C

-99- 700 VHN (kg/mm3!

200

100 Cold rolled 92% ■ Fe50%Co2%V 0 Fe35%Co0.3%Tc

0 0 1 2 3 4 5 6 7 Anneal time (hours! Figure 3.2.6! Variation in microhardness with isothermal annealing at 550°C

-100- (kg/mnfl

400

300

< 200

100 Cold rolled 96% • Fe50%Co2%V 0 Fe35%Co0.3%Ta

0 0.2 0.4 0.6 0.8 1 1.2 Anneal time (hours) 'igure 3.2.7: Variation in macrohardness with isothermal annealing at 550°C Both alloys reached their maximum hardness after 2 hours. This hardness value was then maintained for heat treatments up to 6 hours. According to the microhardness results, the change in hardness for the Fe35%Co0.3%Ta alloy is 50% (316 to 466) and 40% (390 to 550) in the Fe50%Co2%V alloy, though the macrohardness suggests a change.

3.2.4.2 Annealing at 650°C.

The hardness values versus anneal time are shown in Figures 3.2.8 and 3.2.9. During this period the microhardness values increased by 12% and 33% for the 50%Co base and the 35%Co base alloys respectively. There is some disagreement here between micro and macrohardness results for the 35%Co base alloy. The macrohardness values show a much lower hardness change. In addition, there is a distinct peak after 0.5 hours in the Fe50%Co2%V alloy which is not present in the 35%Co alloy.

- 1 0 2 - -103- VHN Ikg/mm'2') 600

100 Cold rolled 92% Fe50%Co2%V 0 Fe35%Co0.3%Ta

Q ______1______1______1______I______1______1______I______0 0.5 1 1.5 2 2.5 3 3.5 4 Anneal time (hours) Figure 3.2.9: Variation in microhardness with isothermal annealing at 650CC

-104- 3.3.0 Texture results.

3.3.1 Assessment of texture.

Figure 3.3.1 shows the (200) and (110) pole positions for the five main textures observed in BCC metals (Hatherley and Hutchinson). Although other components were occasionally present the texture was only described by these five components. Initially (110) and (200) pole figures were produced, but once the predominant texture components were identified, (200) pole figures were found to be sufficient.

3.3.2 Cold rolled textures.

In general, all the cold rolled textures were of weak 11001<110>. The central peak in the 200 pole figures was approximately 3x random and the 4 non central-peaks were slightly more pronounced as the deformation increased from 88 to 96%. Typical 200 pole figures are shown in Figures 3.3.2 to 3.3.4.

3.3.3 Recrystallisation textures of alloys with differing cold deformation.

Fe35%Co0.3%Ta, Fe35%Co0.3%Tal%Ni, Fe50%Co2%V and Fe50%Col.8%V3.5%Ni were all fully recrystallised after 2 hours at 760°C. The predominant recrystallised texture was {100}<110>. In general, for a given composition, it was found that greater deformation resulted in a stronger recrystallised texture, (Figure 3.3.5). Figure 3.3.6 shows the recrystallised texture of Fe50%Col.8%V3.5%Ni with 86% deformation. However the strength and type of texture was also dependent on the sample’s composition.The pole figures of two recrystallised equiatomic alloys (Fe50%Co2%V and Fe50%Col.8%V3.5%Ni), are compared with those of two Fe35%Co base alloys(Fe35%Co0.2%Ta and Fe35%Co0.3%Tal%Ni) in Figures 3.3.7 to 3.3.10. It is clear that the 50% alloys have produced a significantly stronger {100}<110> texture with the exclusion of all the other components. Central peak values in the 200 pole figures correspond to 12 and 8x random for Fe50%Col.8%V3.5%Ni and Fe50%Co2%V respectively. The pole figure for the Fe35%Co0.3%Tal%Ni sample had a strong central peak, and an almost cylindrically symmetrical ring at 80°, suggesting that the {100} planes were in the rolling plane and

-105- TD

Figure 3.3.1: Pole positions in the (200) and (110) pole figures for the predominant textures observed in BCC metals. a>

Figure 3.3.2: Relief and contour pole figure for Fe50%Col.8%V3. 5%Ni (Rotelloy 5) cold rolled 88%.

-107- D TP

Figure 3.3.3: Relief and contour pole figure for Fe50%Col.8%V3.5%Ni (Rotelloy 5) cold rolled 92%.

-108- Jk

Figure 3.3.4: Relief and contour pole figure for Fe50%Col.8%V3.5%Ni (Rotelloy 5) cold rolled 96%.

-109- 14 MAX. INTENSITY/R

12

10

8

r* 0

4

2

l______i______i 0 85 87 89 91 93 95 97 99 % DEFORMATION

Figure 3.3.5: Strength of central peak of (200) pole figures for cold rolled and annealed (750°C/2hrs) samples.

-1 1 0 - AO

Figure 3.3.6: Relief and contour pole figure for recrystall«5£<^ Fe50%Col.8%V3.5%Ni, (Rotelloy 5, cold rolled 88% and 760°C/2hrs).

-Ill- A* I

2W lilt- Z

Figure 3.3.7: Relief and contour pole figure for recrystalltsedl. Fe50%Co2%V, (cold rolled 96% and 760°C/2hrs).

-1 1 2 - Figure 3.3.8: Relief and contour pole figure for recrystallist?c/ Fe50%Col.8%V3.5%Ni, (Rotelloy 5, cold rolled 96% and 760°C/2hrs).

-113- nI

V) bJz Zh

Figure 3.3.9: Relief and contour pole figure for recrystall/S&J Fe35%Co0.3%Tal%Ni, (cold rolled 95% and 760°C/2hrs).

-114 Relief and contour pole figure for r e c r y s t a l W Figure 3 ^cold rolled 95% and 160»C/2hrs). Fe35%CoO

-115- there is no crystallographic direction that aligns preferentially with the rolling direction. Fe35%CoO.3%Ta has a texture very similar to that in the cold rolled material.

3.3.4 Texture and heat treatments.

The pole figures for two Fe35%Co alloys after annealing for 2 hours at temperatures below 800°C and in the main, below Tc , were produced. The compositions investigated were designed to produce differing degrees of precipitation in the range 400 to 600°C. They were:

Fe35%Co0.3%Ta (92%)...... Ta to allow cold deformation. Fe35%Co0.3%Tal-4%Ni (95%).. Ni to enhance precipitation.

Typical 200 pole figures for the 35%Co base alloys are shown in Figures 3.3.11 and 3.3.12. It was found that up to about 500°C there was little change in the predominant {1001<110) texture from that in the cold rolled condition. The {100}<110) texture of the 575°C/2hrs (Figure 3.3.11) was slightly stronger and the sample annealed above Tc , at 670°C had the strongest {100)<110> texture of all the Fe35%Co base samples (Figure 3.3.12). There was also a tendency for the samples with greater nickel content to have a greater {1001<110> texture at temperatures up to 670°C. Samples of Fe35%Co0.3%Ta (92%) alloy were annealed at 750 and 800°C for 2 hours and the texture became equally {1001<110) and {1111<112), Figure 3.3.13.

The samples annealed below Tc were not recrystallised, and etched in less than half the time of the recrystallised samples of the same composition. The 670°C/2hrs samples were found by optical microscopy to be partially recrystallised.

3.3.5 Two stage heat treatments.

The samples and heat treatments are shown in Table 3.3.1.

-116- iue331: eif n cnorpl fgr fr recrystall*^/ for figure pole contour and Relief 3.3.11: Figure e5C03T4N, cl old 5 ad 575°C/2hrs). and 95% (coldrolled Fe35%Co0.3%Ta4%Ni,

INTCNSITV/Cn— 32) -117- Figure 3.3.12: Relief and contour pole figure for recrystallis^ for figure pole contour and Relief 3.3.12: Figure Fe35%Co0.3%Ta, (cold rolled 95% and 670°C/2hrs). and 95% (coldrolled Fe35%Co0.3%Ta,

jNTEMStTvyCrt — 3-7) -118- iue 3Figure Fe35%CoO

IfMTENSITX/CR— 335 .3 Rle ad otu pl fgr fr recrystalli>e,c'/. for figure contour3.13: pole and Relief %a (cold 750°C/2hrs). and3%Ta,92%rolled -119- PJ> Table 3.3.1: Two stage heat treatments.

Sample Heat treatment

86% CR1d Fe50%Co2%V 775°C/lhr, fc to 525°C/2hrs, 63% CR’d Fe35%Co0.3%Ta fc to RT.

96% CR’d Fe50%Col.8%V3.5%Ni 150°C/min to 450°C/1.2hrs, (Rotelloy 5) 2°C/min to 760°C/2hrs, 1.7°C/min to RT.

fc: furnace cooled, CR’d: cold rolled, RT: room temperature.

77$°C i S> u t - i (Jul cUscrrzL

3.3.6 {1Q0K110> texture dependence on ternary additions.

For each composition the cold deformation and anneal temperature that provided the strongest {100 J <110> texture was compared. 94 to 96% deformation was required in all alloys and the best anneal temperature was 670°C for the 35%Co base alloys and 760°C for the Fe50%Co alloys. Figure 3.3.16 shows the height of the central peak in the samples with the strongest {100}<110) texture against ternary addition. The greater the ternary addition the greater is the {1001<110> component of the texture. However, smaller amounts of ternary were added to the 35%Co alloys and, although the trend was for greater {100}<110> texture with greater ternary addition, it has not been shown that 5-6% ternary produces the same strength of alignment in the Fe35%Co base alloys as in the Fe50%Co alloys.

- 1 2 0 - ftp

Figure 3.3.14: Relief and contour pole figure for 86% cold rolled Fe50%Co2%V after 2 stage heat treatment of Table 3.3.1.

-1 2 1 - °) 20 i r\A 1 VC £tn U2 f-z

Figure 3.3.15: Relief and contour pole figures for 96% cold rolled Fe50%Col.8%V3.5%Ni after: (a) : 2 stage heat treatment of Table 3.3.1. compared with (b) : single 760°C/2hrs treatment, showing no difference in crystallographic texture.

- 1 2 2 - Figure 3.3.16: The height of the central peak of the (200) pole figure for various samples after the heat treatment that gave the strongest 1100]<110> texture.

-123- 3.4.0 Results of magnetic measurements.

3.4.1 Sucksmith balance results.

The saturation magnetisation per unit mass, Bsat/m, is shown in Figure 3.4.1. for a selection of hot rolled iron cobalt alloys. There are considerable errors with the force method of measuring saturation in these alloys due to the tendency of the sample to be pulled to one of the pole pieces. In addition, it was found that the results were suscept/ble to sample geometry, such as thickness, and later to be discovered, orientation effects (Sections 3.4.2. & 3.4.3.). Nevertheless, the accuracy of the data is sufficient to ascertain general trends.

It can be seen that the Fe35%Co0.2%Ta alloy has a higher saturation than the Fe50%Co base alloys, and that the ternary addition tends to reduce the saturation as illustrated by the Fe50%Co2%V sample. In this respect, the smaller additions of Ta and Nb offer a distinct advantage.

3.4.2 Vibrating sample magnetometer measurements on textured material.

In all of the results problems were encountered with calibration and allowing for demagnetisation effects. Though every effort was made to determine absolute values of the magnetic parameters, the results must be thought of as semi-quantitative ; more quantitative results are presented in the permeammeter results. The maximum magnetisation did not correspond to saturation though the magnetisation at the maximum field was reproducible.

3.4.2.1 General Results.

Shape of M-H curve irrespective of direction sample was cut from sheet

Figure 3.4.2 shows the M-H loop of 96% cold rolled Rotelloy5 (Fe50%Col.8%V3,5%Ni) compared to that of recrystallised Rotelloy5. Figure 3.4.3 shows the equivalent curves for Fe35%Co0.3%Ta4%Ni cold rolled 95%. The cold rolled sample has a much lower susceptibility, a lower remanence and a higher coercive force than the recrystallised sample. The M-H loop of a recovered sample is shown in Figure 3.4.4 and

-124- Fe35SGo0.2*Ta

Fe49%Go2%V

F©60%Co0.1Ta

F©50%Go0.25Ta

F©608>Go0.2N b

Pur© Fe.

210 220 230 240 250 Saturation magnetisation (©mu/g)

L i relied Figure 3.4.1: Saturation magnetisation of^iron cobalt base alloys measured by the Sucksmith Balance Method.

-125- a ) 5D0 45° to roll direction 200

100

\3 E 0

-100

-200

-300 -f------BO -60 -40 -20 0 20 40 60 BO H(mT)

Figure 3.4.2: A comparison of the induction curves of: a) Rotelloy 5 cold rolled by 96% b) recrystallised Rotelloy 5 after the 2 stage annealing at 450°C/1.2hrs and 760°C/2hrs. (Table 3.3.1, page 120).

Samples were cut in the roll direction and at 45° to the roll direction in both cases, as shown.

-126- Figure 3.4.3: A comparison of the induction curves of: a) Fe35%CoO.3%Ta4%Ni cold rolled by 95% b) recrystallised Fe35%Co0.3%Ta4%Ni after annealing at 670°C/2hrs.

Samples were cut in the roll direction and at 45° to the roll direction in both cases, as shown.

-127- Figure 3.4.4: The induction curves of Fe35%Co0.3%Ta4%Ni cold rolled by 95% after a recovery anneal at 575°C/2hrs.

Samples were cut in the roll direction and at 45° to the roll direction in both cases, as shown.

Table 3.4.1: Summary of V.S.M. results.

Sample Anneal Temp. Remanence Magnetisation Coercivity (°C/2hrs) Mr M(700mT) He RD 45° RD 45° RD 45° Fe35%Co 0.2Ta 670 46.5 41 203.6 207.7 2.2 2.1

0.3TalNi — 28.5 25.5 199.8 202.6 2.6 2.05 670 31 30.5 187.4 195.3 1.95 1.85 760 23 23.5 197.9 187.8 1.65 1.55

0.3Ta4Ni — 16.5 19.5 167.2 168.5 2.5 2.15 575 24.5 25 157 167.5 2.35 2.15 670 29.5 26 176.7 189.1 2.15 1.85 760 19.5 22 169.3 172.9 1.6 1.65 Fe50%Co 2V3.5Ni - 17.5 21 159.2 166.4 3.8 3.3 (Rotelloy 5) 760* 43.5 19 158.2 175.6 2.8 1.15 760* 52.5 20.5 167.6 178.6 3.25 1.25

He = mT: All magnetisation values are in arbitrary units. * denotes 'soaking' at 450°C/1.2hrs prior to 760°C/2hrs. -128- shows an intermediate squareness. A summary of all the V.S.M results is presented in Table 3.4.1. This squaring of the M-H curve is typical of soft magnetic materials and illustrates the deleterious effect of cold work on iron cobalt alloys.

Comparison between Fe35%Co and Fe50%Co base alloys.

From the literature review presented in Section 1, a higher magnetisation and a lower coercive force would be expected in the Fe35%Co alloys when compared with the Fe50%Col.8%V3.5%Ni (Rotelloy 5) sample. It can be seen that on the whole the 35%Co base alloys show a higher magnetisation at 700A/m. However, it should be noticed that this difference may partly be due to dilution of the Rotelloy by the ternaries (Table 3.4.1). As for coercivity, the direction in which the samples are cut from the cold rolled sheet is the more significant than composition and prevents any comparison based on composition being made. For example, the recrystallised Rotelloy shows both the highest and the lowest values of He.

If the remanence values are compared for the Fe35%Co0.3%Tal-4%Ni samples after different heat treatments, it can be seen that the anneals at 575 and 670°C give the greatest reman^nce values. In the Rotelloy 5 sample the remanence is more strongly dependent on the direction the sample is cut from the sheet, after the 760°C/2hrs heat treatment.

Effect of Ternary/Quaternary on M(700A/m) and He.

Figure 3.4.5 shows He and M(700A/m) as a function of the wt% ternary/quaternary added. It is clear that the greater the addition the lower the magnetisation though the coercive force is more greatly affected by the direction in which the sample was cut from the sheet.

-129- Figure 3.4.5: Magnetic properties of textured samples with different ternary additions. Heat treatments were chosen to give maximum {100}<110> texture and were 670°C/2hrs for the Fe35%Co base alloys and 450°C/1.2hrs plus 760°C/2hrs for the Rotelloy 5.

-130- 3.4.2.2 Anisotropic Results.

Magnetisation at 700A/m.

In almost all cases the magnetisation at 700A/m was greater in the samples cut at 45° to the rolling direction, (Figures 3.4.6 and 3.4.7). The spurious result for the Fe35%Co0.3%Tal%Ni annealed at 760°C/2hrs, is thought to be due to experimental error, which is large in all cases, and due mainly to geometric effects in the samples and to slight differences in positioning samples in the VSM. i There is a tendency for the difference in magnetisation with sample direction to be greatest for the Rotelloy 760°C/2hrs and Fe35%Co0.3%Ta4%Ni 670°C/2hrs, (Figure 3.4.5).

Coercivity.

In all the alloys the coercive force is lower in the samples cut at 45° to the roll direction, apart from the Fe35%Co0.3%Ta4%Ni annealed at 760°C/2hrs which is again thought to be due to experimental error. This effect is greatest in the alloys with greater ternary additions (Figure 3.4.5).

Remanence.

In the 35%Co base alloys there is little difference in the remanence values of samples cut in the roll direction and samples cut at 45° to the roll direction. However, in the Rotelloy 5 the difference is considerable. For the 760°C/2hrs annealed Rotelloy 5, the remanence value for the 45° sample is less than one half the value in the roll direction.

3.4.3 Permeammeter measurements on textured material.

The results of the permeammeter measurements on 96% cold rolled Rotelloy 5, are shown in Figure 3.4.8, and Table 3.4.2 in which (s) denotes the 2 stage heat treatment followed by furnace cooling, and the other sample had a single stage heat treatment followed by furnace cooling. There was insufficient accuracy to distinguish between maximum

-131- Figure 3.4.6: Magnetic properties of samples cut from textured from cut samples 3.4.6:of Figureproperties Magnetic Fe35%Co0.3%Tal%Ni sheet, as determined by V.S.M.by Fe35%Co0.3%Tal%Ni sheet, as determined MAGNETISATION (700mT) 220 * ol i‘ 4° o RD to 45° ■*— Dir‘nRoll -132-

4 ORIEFRE (nT) FORCE COERCIVE Fe35%Co0.3%Ta4%Ni sheet, as determined by V.S.M.by Fe35%Co0.3%Ta4%Ni as sheet,determined Figure 3.4.7: Magnetic properties of samples cut from textured cut from samples 3.4.7:of Figure Magnetic properties MAGNETISATION (700nT)

-133- ORIEFRE (mT) FORCE COERCIVE Magnetizing Field. H. (A/m) the roll direction the 45° and roll direction. at to the Figure 3.4.8: Figure Normal 3.4.8: induction measured curves Rotelloy sheet textured for 5 in specimen cut specimen deformed cut from heat sheet 88% after and same stage treatment. two Figure 3.4.9: (Transparent overlay) Figure Normal 3.4.9: overlay) induction ring isotropic curve for (Transparent Flux Flux Density, B, (T)

-134- Flux Density, B, (T)

Figure 3.4.0: Normal induction curves for textured Hotelloy 5 sheet measured in the roll direction and at 45° to the roll direction.

rigure 3.4.9: (Transparent overlay) normal induction curve for isotropic ring specimen cut from 831> deformed sheet and after sane txo stage heat treatment. susceptibilities but it is apparent from the shapes of the curves in Figure 3.4.8 that the initial susceptability is poorest for the two samples cut in the roll direction.

Table 3.4.2: B-H characteristics of textured Rotelloy 5.

Sample Remanence, Br Coercive Force, He (T) (A/m)

Rotelloy5 RD (S) 1.53 375 Rotelloy5 45 (S) 1.48 315 Rotelloy5 RD 1.52 361 Rotelloy5 45 1.41 326

If the permeameter results for Rotelloy 5 (S) are compared to the VSM results for the same material, the VSM values for the remanence and the coercive force in the roll direction and at 45° to the roll direction, have a greater difference, though qualitatively, the results agree with those from the permeameter. This is thought to be due to inaccuracy in measuring the field H, due to demagnetisation, and due to incomplete saturation of the samples in the VSM.

3.4.3.1 Comparison of B-H curves for the single stage and the two stage heat treatment.

It may have been predicted that the longer 2 stage heat treatment would produce a material of larger grain size, and that the longer period in a hydrogen atmosphere would lower still further the deleterious sulphur and carbon . content and lead to a material magnetically superior to the single stage heat treatment.

If this were the case the 4 curves would separate into 2 pairs corresponding to the 2 heat treatments. This does not happen and, although these effects may be present, by far the greatest effect of the two stage thermal treatment is to produce magnetic anisotropy without affecting the crystallographic anisotropy.

-135- 3.4.3.2 Comparison of the two B-H curves of samples cut in the roll direction.

Both of these curves show lower initial susceptibilities, higher coercive forces and higher remanences than a reference sample of 88% deformed Rotelloy 5 measured by a ring technique (Cullity 1972 or more specifically Persiano 1986) (Figure 3.4.9).The sample undergoing the additional 450°C tempering shows the poorest values in Br , He and Uinit. The nature of a ring sample is to average out any directional effects and hence the anisotropy induced in the roll direction in the strips tested here, is deleterious to the magnetic performance in the roll direction.

3.4.3.3 Comparison of the two B-H curves of samples cut at 45° to the roll direction.

There is very little to choose between these two samples. The coercive force is slightly lower in the 2 step heat treatment sample, but the remanence is slightly higher. When compared with the isotropic ring specimen, the differences are again very small and when the slight variations in experimental procedure are accounted for, these differences fade into the associated errors. Hence it is apparent that the anisotropy induced by the heat treatments does not impair the magnetic performance in a direction 45° to the roll direction.

3.4.3.4 Best and Worst magnetic properties due to heat treatment and orientation.

To obtain an appreciation for the magnitude and hence significance of the figures presented in Table 3.4.2, it should be noted that the difference in coercive force between the Rotelloy5 RD (S) and Rotelloy5 45 (S) sample is 19% due solely to the direction in which the sample is cut from the sheet. The single stage heat treatment reduces this gap to 11% but it is still very large.

-136- 3.4.3.5 Error assessment.

The samples were produced with the objective of comparison between samples cut in the roll direction and at 45° to the roll direction. Hence the only source of random error is from differences in positioning the samples in the permeammeter and strains produced by clamping. Because of the large dimensions of the samples these sources are negligible.

Systematic errors are present. The most important, and the source of error most common in permeammeters, is the determination of the field, H. The Hall probe touched the secondary coils around the sample but demagnetisation would reduce the field experienced by the probe.

In addition the cross section of the secondary coils was taken as the cross section of the sample. The cross section of the coils is clearly greater than of the sample and the difference is filled with air. Hence the change in flux, and induced e.m.f., was more than that produced by the Rotelloy 5 alone. This is thought to cause the values of magnetisation to be =2% greater than those of the ring specimens.

-137- 3.5.0 X-Ray Results

3.5.1 Lattice parameter measurements.

The 35%Co alloys have a larger lattice spacing than the Fe50%Co alloys for a given degree of order. Figure 3.5.1 shows the room temperature lattice parameter of Fe35%Co base and Fe50%Co base alloys upon isothermal annealing of initially disordered samples at 465°C. The expansion of the lattice on ordering in the Fe50%Co alloys is clearly detected though there is no corresponding significant increase in the lattice parameter of the Fe35%Co base alloys. To confirm that there was no noticeable change in lattice spacing in the 35%Co base alloys, disordered samples of Fe35%Co0.3%Ta and Fe35%Col.2%Y were isothermally annealed at 465°C. The results are shown in Figure 3.5.2, and it is apparent that the scatter is greater than any lattice parameter change upon ordering.

It is clear that the lattice parameter method is not suitable for measuring LRO in 35%Co base alloys.

In a quench, the temperatures at which the sample^orders are close to the order/disorder transition temperature, Tc. Figure 3.5.3 shows two Fe50%Co base alloys 'isothermally' annealed in a furnace set at 675°C, about 45°C below Tc. At temperatures in this range the samples order within 120 secs (2 mins), and this is too short a time period for the samples to reach the furnace temperature. In other words, the samples start to order at temperatures that are lower than 675°C. This point is illustrated more quantitatively by Figure 3.5.4 which shows the temperature of a thermocouple introduced to the furnace in the same way as the samples. It can be concluded that most of the ordering occurs between 500°C and 600°C but this is the same for all samples and does not invalidate comparisons between different compositions.

There is considerable scatter in the data but .tfia- Fe50%Co2%V alloy orders more quickly than the Fe50%Co0.3%Nb alloy.

-138- 2-865

64 - Fe35%Co0.2%Ta

63 - Ordered Fe35%Co{ Fe35%Co0.3%Ta 15 62 -

56 n

2-855 0 5 17-6 96 NO. OP HOURS AT 465 *C

Figure 3.5.1: The effect of a 465°C anneal on the lattice parameter of initially disordered iron cobalt base alloys. The Fe50%Co2%V curve was produced from two separate sets of samples and illustrates reproducibility. The value for ordered Fe50%Co is from Clegg (1971).

-139- 1600 s 1400 1200 ih tvi uctj ’Ctlictzj • • Fe35%Co0.3%Ta $ Fe35%Col.2%Y

L besnctho -£-nrsT. ______5^y5 U ______J Time Time at 465°C (secs) $ 600 800 1000 ctriit- jr&m ct $>.l. i ______H'fCar cCfbcesiccs kikcerc'- tL n$alh jrtr ~L.. ^UJsra*. 200 400 _ i $■ 3. >. >. 3. Fh. lUote- : 465°C of 465°C of Fe35%Co alloys. disordered two base initially Figure 3.5.2: Figure ordering at parameter Change upon 3.5.2: isothermal in lattice

Lattice parameter (angstroms) 0 2.865 2.864<> 2.863 L

-140- Figure 3.5.3: Rate of isothermal ordering of two initially disordered Fe50%Co base alloys introduced to a furnace at 675°C. a) Fe50%Co2%V b) Fe50%Co0.3%Nb.

'Y* or 'N* denote whether superlattice peaks were detected using CoKa radiation.

-141- Figure 3.5.4: The rate of heating of a thermocouple introduced to a furnace at 675°C which was used to assess the temperature range over which the samples in Figure 3.5.3 ordered.

-142- 3 .5 .2 Superlattice Measurements.

Figure 3.5.5 (a) shows that it is possible to detect the 100 superlattice line in fully ordered Fe35%Co base alloys after isothermal annealing at 465°C for 5 days, and Figure 3.5.5 (b) shows the corresponding 200 fundamental lines. The X-ray line intensities are given in Table 3.5.1.

However, it was not possible to detect superlattice lines with any accuracy for partially ordered samples. This is a consequence of the peak intensity being proportional to the degree of long range order squared:

l a S2

Because the maximum degree of order attainable in the 35%Co alloys is 0.7, the superlattice peak size for fully ordered Fe35%Co is half that of the fully ordered equiatomic alloy, at 465°C. If a temperature closer to Tc is chosen, where the equilibrium degree of order is nearer to one half of the maximum degree of order, then the superlattice peak size is reduced by a factor of about 4. Combine this with the requirement that partial degrees of order need to be measured to determine ordering rates, and the size of the peak that needs to be measured reduces to a size indistinguishable from the background.

Whatever the superlattice peak size, it must then be compared with a fundamental line which is not diminished with temperature and for the fully ordered Fe35%Co alloys, is approximately 50 times bigger, Table 3.5.1. This peak must be known with a similar percentage accuracy, so as not to augment the errors.

Clearly X-ray superlattice line measurement is not an accurate method for measuring partial LRO in Fe35%Co base alloys.

-14 3 - Counts (xlQQO)

Figure 3.5.5: Superlattice and fundamental x-ray diffraction lines for iron cobalt alloys annealed at 465°C for 5 days.

-144- Table 3.5.1: X-Ray diffraction line intensities for fully ordered samples.

Specimen Area under X-ray peak (200) Fundamental (100) S'lattice Ratio

Fe49%CoO.2%V0.4%Nb 119911 4871 24.6 Fe49%CoO.3 5%NbO.5%Mo 41255 1963 21.0 FeCo2%V 23575 906 26.1 Fe35%CoO.3%Ta 146549 3026 48.4 3.6 .0 Thermal Analysis Results

A typical DTA and DSC plot for an equiatomic base alloy are shown in Figures 3.6.1 and 3.6.2. The DTA clearly shows peaks corresponding to the order/disorder and the BCC/FCC phase transitions on both the heating and cooling cycle. Though similar, the DSC shows an additional broad peak at about 500°C. This is the '550' anomaly described in Section 1.

Figures 3.6.3 and 3.6.4 show the same traces for an Fe35%Co base alloy. In these alloys the order/disorder peak is much reduced but still detectable, and the 550 anomaly is still present in the heating cycle of the DSC trace.

3.6.1 Transition temperatures.

3.6.1.1. Binary alloys.

The mean temperatures of the heating and cooling peaks for the order/disorder and the a/tf transitions for Fe35%Co base, and Fe50%Co base alloys as determined by DSC are shown in Table 3.6.1. As already illustrated in Figures 3.6.1 to 3.6.4, the hysteresis between the heating and cooling peaks is greater for the a/)f transition than for the ordering transition.

3.6.1.2 Additions of Nb, Ta and Ni to binary alloys-

On addition of less than 1 wt%Ta or Nb, there is little change in the order/disorder or the a/8 transition temperatures for the Fe35%Co based alloys, (Table 3.6.1). However, the table demonstrates that the addition of up to 4 wt%Ni to Fe35%Co0.3%Ta does significantly reduce the a/% transition temperature, and to a lesser extent, the order/disorder transition temperature.

-146- mJ/sec iue ..: ... rc f e9C02V.%bsoig '550showing anomqly', 3.6.2: Figure trace D.S.C.of Fe49%Co0.2%V0.4%Nb re/iodradapagmapae transitions.order/disorderandalpha/gammaphase eeec: apems o P. etn ae 20°C/min rate: Heating mg 29.7 weight: Sample Pt. of mass Sample Reference: rolled. cold condition: Initial eeec: lmn odr Haigrt: 10°C/minrate: Heating Reference:powder. Alumina Figure 3.6.1: D.T.A. trace of Fe50%Co0.1%Ta showing order/disordershowing andFigure 3.6.1: Fe50%Co0.1%Ta traceof D.T.A. nta cniin cl old Sml egt 27 Samplemgweight: Initial condition:rolled. cold alpha/gamma transitions.phase 147 mJ/sec eeec: apems o P. etn rt: 20°C/min rate: Heating Pt. of mass Sample Reference: showingpeakdouble Figure 3.6.4: trace of Fe35%Co0.3%Nb0.3Misch D.S.C. nta cniin cl old Sml egt 2. mg 29.9 weight: Sample rolled. cold condition: Initial smallershowingFigure traceFe35%Co0.2%Ta3.6.3:muchof D.T.A. andalpha/gammapeaks. the order/disorderadditiontoinin atand=550°C cycles500°Cheating eeec: lmn odr Haigrt: 10°C/minrate: Heating Sample21.9mgweight: Reference:powder. Alumina Initial condition:rolled. hot order/disorderpeak. 0 TMEAUEC 1100 TEMPERATURE(C) 500 -148- Table 3.6.1: Transition temperatures of FeCo alloys. (All figures ± 5°C)

Order/disorder (°C) a/JT (°C) Fe50%Co 730 980 0.1%Ta 728 978 0.25%Ta 727 976 0.5%Ta 727 976 1.0%Ta 717 957

0.l%Nb 720 968 0.2%Nb 723 971 0.3%Nb 726 968 0.4%Nb 730 972 0.5%Nb 725 970

Fe50%Col.8%V3.5%Ni 685 851 (Rotelloy 5) Fe50%Co2%V 707 913 (2V-Permendur)

Fe35%Co 638 967 0.2%Ta 643 974 0.3%Ta 646 972

0.25%Nb 642 967 0.3%Nb 643 969

0.04Ce 640 974

Fe35%CoO.3%Tal%Ni 630 944 2%Ni 634 926 3%Ni 626 903 4%Ni 634 879

-149- 3.6.1.3 Addition of 3.5%Ni to Fe50%Col.8%V (Rotelloy 5).

Additions of Ta and Nb to Fe50%Co slightly lower the cl/ $ transition temperature and produce a marginal reduction in the order/disorder temperature. These effects are small and any change over^ternary addition range is less than the inherent errors in the experiment.

However, on addition of 3.5%Ni, there was a lowering of the order/disorder peak by z2Q° to 685°C compared to 707°C for the FeCo2%V. More dramatic was the lowering and separation of the a/2f transition. The heating peak occurred at 932°C and the cooling peak at 770°C, giving a mean of 851°C, 60°C lower than the mean for FeCo2V. These results are similar to those of Persiano (1986) who performed DTA at a rate of 5°C/min on FeCo5.4%V, (i.e. the same alloy but with vanadium substituted for nickel) and found the order/disorder peak at 680°C and the a/^ transition at 803°C. The atomic weights are such that there are zlO% more vanadium atoms in equivalent weights of vanadium and nickel. Hence, it appears that V and Ni lower the order/disorder and a/X transition temperatures by a similar amount per atom.

3.6.1.4 Additions of Ce, Y and Misch metal to Fe35%Cc.

These alloys had very small percentages of ternary which had little effect on the transition temperatures.

3.6.2 The 550 anomqly.

3.6.2.1 The peak in binary alloys.

The traces for disordered binary Fe35%Co and Fe50%Co using Pt as a reference are shown in Figures 3.6.5 and 3.6.6. At 490°C and 570°C there are two overlapping peaks for the Fe35%Co sample whereas the Fe50%Co binary alloy has a single larger peak at 520°C. If the trace is repeated, as shown for Fe50%Co2%V and Fe35%CoO.3%Ta in Figure 3.6.7 & 3.6.8, then the second run does not follow the path of the first run, but the third run does follow the path of the second run. In spite of the inert atmosphere, the sample undergoes a degree of oxidation during the first run which contributes to the exothermic peak at 520°C. On the second and third runs, the thin oxide layer which is coating the

-150- mJ/sec iue ..: ... rc f iayF5%osoigeohri ek atexothermicshowingpeak Fe50%Co 3.6.6: Figure trace of binary D.S.C. eeec: t 15mg. = Pt Reference: hrceitco te '550characteristic theanomqly'. of nta cniin ht rolled. hot condition: Initial 2° nfrthaigadedtemcpa a 600°Cinsecond atheating,peak 525°C inandfirstendothermic heating eeec: apems o P. etn ae 20°C/minrate: Heating Reference: of SamplePt.mass thatat 500°C showing peakFigure 3.6.5:Fe35%Co tracebinaryof D.S.C. nta cniin ht old Sml egt 29.5mg Sampleweight: Initial condition: rolled. hot cannot toprecipitation. bedue -151- etn ae 20°C/minrate:Heating apewih: 23 mgSampleweight: Figure 3.6.7: D.S.C. trace of Fe50%Co2%V repeated three times without changing the sample, showing the effect of oxidation and ordering.

Initial condition: cold rolled. Sample weight: 15 mg Reference: Sample mass of Pt. Heating rate: 20°C/min

Figure 3.6.8: D.S.C. trace of Fe35%CoO.3%Ta repeated three times without changing the sample, showing the effect of oxidation and ordering.

Initial condition: cold rolled. Sample weight: 15 mg Reference: Sample mass of Pt. Heating rate: 20°C/min

-152- sample, combined with the fact that the sample is highly ordered for the second and third runs, causes a deviation from the baseline of run 1 below 450°C. Bearing in mind that the sample is highly ordered for the second and third runs it is clear that the endothermic peak at 600°C corresponds to the rapid disordering described in Section 1, and the exothermic peak due to ordering, of the first run is no longer present.

The two overlapping peaks at about 490 and 570°C observed in the Fe35%Co sample run against a platinum reference (Figure 3.6.5) are observed in many of the Fe35%Co base alloys. Note that because it exists in the binary, neither of these peaks can be due to precipitation.

3.6.2.2 Cancellation of peaks.

Figure 3.6.9 shows how complete cancellation of any peaks can be achieved by running identical samples against each other. The order/disorder, atransition and the oxidation peaks are not observed.

3.6.2.3 Use of an annealed, ordered sample as a reference.

In order to observe the 550 anomaly more closely, an annealed sample was run against a cold rolled sample. This corresponds to running an ordered sample against a disordered sample and produces peaks that are due solely to differences in order if there is complete cancellation of the a/a', a/# and the oxidation peaks. The degree of cancellation of the very large a/& peak gives an indication of how closely matched the sample and reference are. Figures 3.6.10 & 3.6.11 show typical traces using an annealed sample as the reference. Unfortunately the binary alloy was not available cold rolled sheet i.e. highly disordered, and consequently was not suitable for investigation by this method.

-153- mJ/sec eeec Sml Haigrt: 20°C/rainrate: Heating SampleReference= Figure 3.6.9: D.S.C. trace of identical reference and sample ofandsamplereference identical 3.6.9: of Figure trace D.S.C. oely5 (Fe50%Col.8%V3.5%Ni) cancellationcompletehow illustrating5 Rotelloy nta cniin cl old Sml egt 7.8 Sampleragweight: Initial condition:rolled. cold of canbeachieved. peaks -154- -155- ordered reference, the Note reference, isolating ordered how peak'. not peak is '550 Initial condition: cold rolled. rolled. condition: cold Initial weight: mg 21 Sample present if run the heating. is continued if present to a second Figure 3.6.10: D.S.C. D.S.C. of trace disordered Fe50%Co2%V Figure 3.6.10: an run against Reference: ordered sample. Reference: Heating rate: 20°C/min an ordered reference, ordered reference, showing an the peak' which not present is if '550 Initial condition: cold rolled. Sample weight: 3.7 mg Reference: ordered sample. Heating rate: 20°C/min the run is run repeated. is the Figure 3.6.11: D.S.C. D.S.C. trace of Figure 3.6.11: disordered Fe35%Co0.3%Tal%Ni against run D B S /p iU 3.6.2.4 Oxidation and the DSC trace.

An annealed sample from a prior DSC run with an oxidised surface was run against an identical sample with the oxide polished off. The results for a Fe35%Co base alloy and for a Fe50%Co base alloy are shown in Figures 3.6.12 and 3.6.13. Here they are superimposed on the sample run against a Pt reference and on a run where the reference was an annealed sample. The lower peak in the Fe35%Co binary vs Pt is due to oxidation, and in the Fe50%Co base alloys the oxidation peak is at a temperature above the '550 peak’. Thermo Gravimetric analysis was also used to determine the onset of oxidation and was found to agree with the DSC results, Figure 3.6.14.

-15 6 - 12 -

Shoulder due to oxidation.

------1 ------i------1------1------1------[ | | i 0 100 200 300 400 500 GtiO 700 800 900 D eg C Figure 3.6.12: Heating curves for Fe35%Co base alloys with different references used to isolate oxidation peak. a) polished disordered Fe35%Co vs Pt reference. b) polished ordered Fe35%Co0.3%Ta vs oxidised ordered Fe35%CoO.3%Ta. c) polished disordered Fe35%Co vs polished ordered Fe35%Co.

Figure 3.6.13: Heating curves for Fe50%Co2%V with different references used to isolate oxidation peak. a) polished disordered Fe50%Co2%V vs Pt reference. b) polished ordered Fe50%Co2%V vs oxidised ordered Fe50%Co2%V. c) polished disordered Fe50%Co2%V vs polished ordered Fe50%Co2%V.

-157- WEIGHT

Figure 3.6.14: Combined Thermogravimetric and D.T.A. traces of Fe50%Co2%V showing how some of '550 peak' is due to oxidation above 450°Cf indicated by weight gain.

Initial condition: 760°C/2hrs. Sample weight: 10.51 mg Reference: alumina powder. Heating rate: 10°C/min

-158- 3.7.1.Tensile test results.

The results of the tensile tests are shown in Table 3.7.1. A thorough investigation of the influence of grain size,composition and cold roll direction upon the tensile properties was not undertaken and the results are presented as a reference for future workers.

Table 3.7.1: Tensile test results.

Sample Condition Grain Max. Max Tensile Yield size Elong. stress point (pm) (%) (MN/m2) (MN/m2) ord dis ord dis ord dis ord dis

Fe35%Co0.2%Ta X-rolled 20.0 18.9 10.6 30.0 560 660 380 380 Fe35%Co0.3%Ta X-rolled 18.6 18.2 1.6 30.2 360 620 360 330 Fe35%Co0.25%Nb //-rolled 13.5 12.3 10.6 29.8 610 600 400 350

Fe50%Co0.3%Nb //-rolled 22.8 18.9 3.3 13.7 350 700 350 430 Fe50%Co0.3%Nb X-rolled 19.6 15.0 6.1 29.5 550 580 350 320 Fe50%Co0.2%Ta //-rolled 12.7 22.8 5.3 16.5 450 720 340 410

It is clear that long range ordering induces a considerable degree of brittleness to the Fe35%Co alloys.

Cross rolling produces a more ductile cold rolled sample than parallel rolling.

The disordered parallel Fe35%Co base alloy does show greater ductility than the equivalent disordered Fe50%Co base alloys. The Fe35%Co alloys are generally easier to file and cut when preparing samples and, as already seen in Section 3.2, are softer.

-159- 4: Discussion

4.1.0 Impact tests, grain size.and reliability.

4.1.1 DBTT and reliability.

The impact test experiments were initially undertaken to try to explain why some alloys did not cold roll, but did warm roll (at sl00°C). As can be seen from Section 3.1. alloys which required warming to sl00°C to roll, underwent a brittle to ductile transition between 20 and »100°C. This DBTT is dependent on many factors including strain rate, so it is not appropriate to say that a specimen with a DBTT at 50°C will roll at 55°C but not at 45°C. The lower strain rates of cold rolling suggest that the temperature at which cold rolling will become feasible is lower than the DBTT. However, it is apparent that there is a qualitative correlation between the DBTT and the minimum temperature at which an alloy will cold roll.

A material that cold rolls without the need for prior warming is desirable and so the relationship between grain size and the DBTT was determined. For this set of experiments, it was only possible to enlarge the grain size of the hot rolled sheet available.

Many polycrystalline metals obey the Hall-Petch relation, which states that: a = oi + kd-* 4.1.1. where a is the yield or flow stress, ai is the intercept stress, k is the Petch slope and d is the grain diameter. In binary Fe50%Co, oi increases linearly with strain rate (Marcinkowski and Fisher, 1965) and is a function of test temperature in many alloys (Jordon and Stoloff, 1969). The Petch slope, k, has been reported to depend on composition and temperature (Johnston et al., 1965).

-160- For a metal to fracture at the yield point, due to dislocation pile ups at the grain boundaries or at other obstacles,

crkd* £ {SpY (jyctlpo/xj ktroLop-i:') 4.1.2. where 3 is a geometric factor dependent on the macroscopic ratio between shear and tensile stress, p is the shear modulus and is the true elastic surface energy. According to Jordon and Stoloff's review, Armstrong suggested that (Spif was dependent on temperature and grain size. At the ductile to brittle transition temperature, equation 4.1.2. becomes an equality, but it is not possible to ascertain the temperature dependence of all the terms.

Figure 3.1.11. shows In(DBTT) vs In d with a reference gradient of %. Allowing for the unprecise nature of the DBTT, a d* dependence is observed for grains smaller than 30pm, corresponding to DBTT's below 150°C. By using equation 4.1.2, it can tentatively be suggested that, for high strain rates and small grain sizes:

DBTT a (SpY/ak 4.1.3.

For larger grain sizes, it is not clear if there is a ceiling temperature above which the disordered samples always break in a ductile manner. The result for Fe35%Co0.2%Ta shown in Figure 3.1.9. suggests that this is the case, and Johnston et al. (1965), found that at low strain rates ordered FeCo-2V fractured in a ductile manner above 450°C.

The experimental factors are responsible for the brittleness incurred at test temperatures > 280°C, by the Fe35%Co0.3%Ta sample with a grain size of 67.6pm. In this case the result of the DSC '550 anomaly’ must be consulted. To reach a test temperature of 280°C the sample must be kept in a furnace at s420°C; to test at 350°C, a furnace temperature of 540°C is required. The low temperature DSC peak, Td, at which ordering can occur is at 520°C for Fe35%Co0.2%Ta, and it has a peak width of tens of degrees. The high temperature brittleness is a result of the specimen ordering in the furnace prior to testing. This puts an upper DBTT limit on the samples. A schematic brittle to ductile to brittle curve is shown in Figure 4.1.1.

- 1 6 1 - Energy

Limit Due to Ordering

Rolling Temperature Range

Ordered Alloy

60 100 160 200 260 300 360 Test Temperature (°OJ

Figure 4.1.1: Schematic ductile to brittle transitions in iron cobalt alloys.

- 1 6 2 - If an FeCo base alloy were to b.e used in the disordered state, at temperatures greater than «400°C, it should be noted that the material will order and hence its properties will change. As these alloys are normally used in the fully ordered condition, this upper temperature limit on ductility is unlikely to cause any problems for the in-service components but should be borne in mind in the production and processing of the alloys.

4.1.2 Mode of fracture in impact test specimens.

Transgranular cleavage is observed in the disordered Fe35%Co base specimens at temperatures below the relevant DBTT, Figure 3.1.12.(a & b). This is similar to the transgranular cleavage reported by Marcinkowski 1970, for disordered Fe50%Co. He also reported that the brittle fracture of ordered Fe50%Co was by intergranular cleavage and this fact is in keeping with the explanation given, that Fe35%Co0.3%Ta (67.6pm) is ordering in the furnace prior to testing. The brittle fracture surface of the Fe35%Co0.3%Ta tested at 280°C is shown in Figure 3.1.12. (e & f). No reference to the mode of fracture of Fe35%Co alloys has been found prior to this.

4.1.3 Grain size refinement by heat treatments.

It may be of use if the grain size of an alloy that warm rolled could be reduced by heat treating to enable cold rolling to proceed. To this end the grain size of material heated through the a/X phase boundary was monitored. Figures 3.1.1 and 3.1.2 show that, although there is a dip in grain size as the phase boundary is crossed the initial grain size is still the smallest. Hence a finer grain size cannot be produced in either the 35%Co or 50%Co alloys by this method.

4.1.4 Grain size refinement by composition.

It has been noted throughout this study that Fe35%Co base alloys have a larger hot rolled grain size and that more difficulties are encountered with cold rolling them. The results of Section 3.1, Figures 3.1.3 and . dlc/Aj-nfe 3.1.4, show a clear tendency for greater wt% additions of ternary/to produce smaller grains in both 50%Co and 35%Co base alloys. An extreme case is 96% cold rolled Rotelloy 5 (Fe50%Col.8%V3.5%Ni)

- 1 6 3 - i which has a grain size of 7pm after recrystallising at 760°C/2hrs. Between 1 and 4% Ni was added to Fe35%Co0.3%Ta with a view to increasing precipitation and decreasing grain size. The grain size in the 95% cold rolled and annealed^at 760°C/2hrs was 16 and 13 pm for the 1% and 4%Ni alloys respectively.

There are several ways in which a nickel addition can reduce grain size. If it remains in solution it may accelerate the ordering kinetics. Hence, if ordering occurs before recrystallisation, atomic mobility in the ordered material is slower and grain growth will be retarded. However, there is no evidence for nickel affecting the ordering kinetics (Pitt and Rawlings 1981).

It is more likely that the grain size refinement is due to X phase precipitation. Second phase particles can modify the recrystallised grain structure in three ways. Firstly, particles can nucleate new grains. This is not observed by microscopy in this work or in the work of Pitt and Rawlings. Secondly, precipitates can k n a J k r4" grams through the deformed matrix. K i a y grain growth ceases when the average grain size, Dmax ~ 4r/3f, where r and f are the radius and volume fraction of the particles. This corresponds to the results in this study, where greater quantities of precipitate result in smaller grain size. Thirdly, the interface between recrystallised grains can be pinned by precipitate particles, and results in precipitates being observed at grain boundaries. This also fits the experimental results from this work.

It is concluded that greater ternary addition produces a finer grain size through limiting primary grain growth and impeding the movement of recrystallised grain boundaries, which in turn lowers the DBTT and enables the alloy to be rolled at a lower temperature.

-164- tK

It is known that interstitial impurities such as C, 0, H r & N, reduce ductility in BCC structures by pinning dislocations, e.g. C and N in iron, and 0 in tantalum. In addition, Shiryayev (1984), report that these interstitials limit structural domain size, and found that with very high purity Fe50%Co alloys (C+O+N+H = 0.005at%), ductility was still observed in ordered material at -196°C.

It is possible that the ternaries added to FeCo alloys in this study, form carbides, oxides etc. with the interstitials, and the compounds produced do not diffuse to, and hence do not pin, the dislocations. Alternatively, if precipitates of this nature, become large enough to be incoherent, ductility may be induced by limiting the size of dislocation pile-ups or by promoting cross-slip. This latter type of precipitate could be of the C03X type reported by many workers. Ductility has been successfully imparted to tungsten alloys by the addition of a thoria dispersion (King et al. 1965).

It is known that vanadium, niobium and tantalum form carbides and nitrides in steels at between 450 and 600°C, so compounds of the form

V4 C3 , Nb4C3, Ta4C3, VN, NbN and TaN exist (Honeycombe P60). This 'removal' of the free atomic interstials from the matrix would offer an explanation as to why alloys of identical grain size but different composition, have different DBTT's.

- 1 6 5 - 4.2.0 Magnetic anisotropy.

The result of the study of magnetic anisotropy was that the 96% cold rolled and 2 stage annealed (450°C/1.2hrs and 760°C/2hrs) Fe50%Col.8%V3.5%Ni (Rotelloy 5) sample had a very pronounced magnetic anisotropy, whereas the single stage annealed Rotelloy 5, and the 35%Co base alloys had marginal magnetic anisotropy.

The question that concerns us is whether the magnetic anisotropy is a consequence of the crystalline anisotropy, or if it is due to microstructural features, such as lath-like precipitates lying in the roll direction which in turn could lead to directional magnetic properties. The question as to whether a strong crystallographic texture can be produced without precipitation, or if anisotropic precipitation can occur without a strong texture is recognised, and discussed in Section 4.3.

In passing, Josso 1974, reported that the best magnetic properties for Fe50%Co2%V were achieved by cooling from 870°C at an intermediate rate of 250°C/hr, (as opposed to 25° or 10000°C/hr) and reasoned that this was because K1 was close to zero. In my opinion a cooling rate of 250°C/hr will produce an alloy with a value of long range order of >0.8, Section 1.1. The greater susceptibility and lower coercivity than the quenched alloy is due to fewer strains in the lattice, fewer vacancies and the fact that the ordered material per se, has better magnetic properties than the disordered material. The extremely slow cooling rate of 25°C/hr, enables precipitates to form between 600 and 450°C, which, as shall become clear, are detrimental to magnetic properties, (also Bozorth 1951).

McKeehan, Hall and Villemain agree that K1 is between 10 and -10 J/m3 depending on whether the sample is ordered or disordered. Here, the Rotelloy 5 samples were cooled at 100°C/hr and hence a degree of order of >0.5 is expected with a corresponding value of K1 between 0 and +10'3//n5 In general a value of K1 close to zero is beneficial to susceptability and coercivity in a soft magnetic alloy, c.f. Permalloys. The 35%Co alloys were cooled in air over 2 to 3 minutes, but in this case K1 is between 10 and 15 J/m3 , depending on order.

-166- If the existence of any second phase is not accounted for and it is assumed that the sample has the same values of K1 and K2 as the relevant binary, then if the magnetic anisotropy is greatest in the samples with the strongest crystalline anisotropy, a crystallographic interpretation of magnetic anisotropy would be appropriate. However, the results from the current work lead to an alternative interpretation; these results are: i) The 50%Co base alloys have a value of K1 close to zero, measured independently by 3 different research groups, Mckeehan (1937), Hall (1960) and Villemain (1972) and hence there should be no easy direction of magnetisation. ii) Two Rotelloy 5 samples, with indistinguishable crystallographic texture were produced; one alloy had the additional temper at 450°C/1.2hrs. The magnetic anisotropy in the 2 stage heat treatment, was much greater than that in the single 760°C/2hrs heat treatment, but the texture was identical. iii) The greater value of K1 for the 35%Co alloys should compensate for the weaker textures present in these alloys, and give values of magnetic anisotropy similar to those of the Fe50%Co base alloys, but this is not observed.

The following results indicate that an alternative interpretation of the magnetic anisotropy may be associated with precipitation: i) The magnetic anisotropy is greater for alloys with greater ternary additions, and hence the potential to form more precipitates, e.g. Rotelloy 5 has the greatest magnetic anisotropy and additions of 1.8%V + 3.5%Ni, whereas little anisotropy is seen in the magnetic properties of Fe35%Co0.2%Ta. (jcibU 3,^.1^ ii) The magnetic properties in a given direction in the alloys exhibiting magnetic anisotropy are equal or worse than those of the same alloy with no crystalline anisotropy. They are not better than random in either of the main crystallographic directions, <100> or <110>. Precipitation can only detract from soft magnetic properties by, for example, the precipitates disrupting magnetic domain movement.

- 1 6 7 - iii) The magnetostriction values are higher in the 50%Co base alloys than the Fe35%Co base alloys. iv) Precipitation on prefen-ed crystallographic planes has been established in Fe50%Co2%V (Ashby et al. 1977) and 'Remendur', Fe50%Co3%V (Mahajan et al. 1974, Pinnel et al. 1976). v) The SEM pictures of Rotelloy 5 after the two stage and single stage heat treatments are indistinguishable with both samples showing considerable globular precipitation and a small grain size, (Figure 4.2.1, a & b ) .

On this evidence, a purely crystallographic interpretation must be ruled out in favour of an explanation in terms of precipitation. The precipitate morphologies that could lead to magnetic anisotropy are:

1) Spherical precipitates. If the values for magnetostriction in the <110> and the <100> direction are significantly different, then the strains created at the precipitate-matrix interface will be different in the 2 directions. Persiano and Orrock showed precipitates of Nb and Ta were spherical, and globular precipitates are found at grain boundaries in recrystallised FeCo2%V and Fe50%Co-V-Ni (Pitt & Rawlings 1981).

2) Elongated precipitates forming on existing crystallographic planes. Previous work has shown that rod-like C03X precipitates can habit the (1101 planes, and have the long axis in the <111> direction, (Ashby, Flower & and Rawlings 1977). If the magnetisation is parallel to the precipitates' long axis, more of the matrix is affected by magnetostriction strains at the precipitate-matrix interface than if the magnetisation is in the short axis direction.

3) Elongated precipitates other than those formed on existing crystallographic planes. These can be formed by the rolling of a precipitate present during hot or cold rolling, or by the growth of precipitates along the grain boundaries of elongated grains.

- 1 6 8 - a)

Figure 4.2.1: SEM micrographs of Rotelloy 5 cold rolled 96% and recrystallised. Etchant: 2% Nital. a) 760°C/2hrs and furnace cooled, x 1000,* '10pm. b) 450°C/1.2hrs & 760°C/2hrs and furnace cooled, x 5000, 5pm. -169- 4) Inclusion, Strain Theories and N6el spike domains. The energy per unit area of a domain wall is reduced if it is occupied by a non-magnetic inclusion. Strains around the precipitates can pin domain walls and sometimes nucleate reverse domains. These effects are normally small when compared to magnetostrictive effects, and do not lend themselves to an explanation of anisotropy.

Explanations 1,2 and 3 are the most likely explanations of magnetic anisotropy because the magnetostrictive strains are very large in iron cobalt alloys. In fact, in 1951, the highest reported value of magnetostriction for any material was 130 x 10-6 in ’hard rolled Fe70%Co tape’, (Bozorth 1951). In the following calculations, the more recent values of Hall (1960) are used.

4.2.1 Calculation of deformation of spherical precipitate.

If the saturation magnetostriction values, Mijk in the <100> and <111> principle crystallographic directions are taken from Hall, Muo can be calculated from the simplified "two-constant" formula for cubic crystals (e.g.Cullity 1972):

Mijk = MlOO + 3 (Ml 1 1 - Ml 00 ) (ell 2ct2 2 + Cl22(l3 2 + (13 2ai 2 ) ...... (1) where an are the direction cosines between the magnetising field and the principal crystallographic directions. The values of Mijk for ordered FeCo are shown in Table 4.2.1.

Table 4.2.1: Magnetostriction values after Hall (1960)

Composition Magnetostriction x 10~6 Fe:Co <100> <110> <111>

50:50 140 65 40 65:35 120 45 20

All values ± 10

-170- If we now consider a spherical hole in a demagnetised single crystal of Fe50%Co, it is possible to calculate the radii of the induced ellipsoid under conditions of saturation magnetisation, in different directions, Figure 4.2.2.(a). It should be noted that the shape change of the hole in the matrix is the same as the shape change as a sphere of matrix material.

The coercivity is due to the precipitate particles resisting the shape change of the matrix, and thereby introducing localised stresses in the matrix. Hence, a magnetisation direction that produces the greatest distortion of the precipitate particle, has the greatest coercivity. We shall assume that the precipitate has no magnetostriction itself and that it was originally formed when the matrix was fully demagnetised, Figure 4.2.2.(a).

The greater the shape change, the greater the strain energy in the lattice around a real precipitate, and the higher the coercivity. The full "Two constant” formula for cubic crystals is applicable when one is interested in magnetostriction values at an angle to the magnetising field:

M = 3/2 MlOO (fll 2{Sl *+02 2f&2 2+Cl3 2|$3 2”l/3)

+ 3M l l l (d l 02 P i 02 +0203 02 03 +03 Ol &3 P i ) ...... (2) where are the cosines of the angle between the principle crystallographic directions and the direction of magnetostriction, and On are as before. Figure 4.2.2.

Table 4.2.2. shows the magnetostriction in the principle crystallographic directions after magnetic saturation in the <100>, <110> and <111> directions.

- 1 7 1 - Figure 4.2.2: Deformation of precipitates due to magnetostriction.

-172- 2

y (Oioj *rootl0n u 8Snet. aire°t!o^,0n hiO]

Dlre°tlon c siphs 0ln°* !*QtQ j ",. " 00* otc. oo* 87 6 7 fiSilt 6 <-*.3. ■ list ?ation ° * * itQct ion cos lt*Gs

"i 73- Table 4.2.2: Magnetostriction values.

Field Magnetostriction 6 Magnetostriction Direction Direction x 10-6

Pi P2 03

<100> lx<100> 1 0 0 0 140 a =1 2x<010> 0 1 0 90 -70 ii © p 4x<110> 1//2 1//2 0 45 35 II o p 2x<011> 0 1//2 1/42 90 -70 4x 1//3 1//3 1/43 55 0

<110> 2x<100> 1 0 0 45 35 a =l//2 lx<001> 0 0 1 90 -70 ii p to lx<110> 1//2 1//2 0 0 65 II p o lx 1/42 -1//2 0 90 5 4x<011> 0 1/42 1/42 60 -35 2x 1//3 1//3 1/43 35 20 2x<-lll> -1//3 1//3 1/43 90 -20

<111> 3x<100> 1 0 0 55 0 a =l//3 3x<110> 1/42 1/42 0 35 20 a =l//3 3x 1/42 -1/42 0 90 -20 a =l//3 lx 1//3 1/43 1/43 0 40 3x 1//3 -1/43 1//3 70 -13 1/46 1/46 -2/46 90 -103

NB: 2x<100> means there are 2 directions with equivalent magnetostriction to the <100> direction. 0 is the angle between the magnetisation and the magnetostriction direction.

The fractional changes in length of the axes of a sphere to form an ellipsoid by magnetostriction are shown in Table 4.2.3.

-174- Table 4.2.3: Fractional change (x 10-6) of the axes of a sphere due to magnetostriction for Fe50%Co alloys.

[100] field [110] field [111] field 140 x -70 x -70 65 x 5 x -70 40 x -20 x -103

From these results it appears that a spherical precipitate is considerably more disruptive when the matrix is magnetised in the <100> as opposed to the <110> direction,conflicting with the experimental results.

4.2.2.1. Allowing for precipitates to form in magnetised domains.

It is possible to extend our assumptions and say that the easy direction of magnetisation is the <100> and hence the precipitates formed in a matrix which consisted of domains already magnetised in the <100> direction of magnetisation.

If a field is applied so that all the precipitates are in a matrix domain of [100] orientation, 2/3rds of the spherical particles (the ones magnetised in [010],[001],[0-10] & [00-1]) will form ellipsoids in the [100] direction. These ellipsoids will be surrounded by strains up to twice the strains experienced by ellipsoids formed from a theoretical demagnetised matrix, Figure 4.2.2.(b). i.e. The total stored strain energy must be less than the strains produced by going to the demagnetised state plus the strains due to magnetising the demagnetised state in the new direction. However, in reality, the ellipsoids do not go via a demagnetised state, and intermediate ellipsoids will be created as the domains rotate to the new [100] direction, so there will be some sort of preferred path involving minimal deformation. The mathematics are complicated and this concept is not likely to affect the strains in different directions by a factor of more than 2 or 3. This will not be significant if extended precipitates are present with aspect ratios much greater than 2 or 3.

-175- 4.2.2.2. Change in easy direction with ordering.

While discussing precipitation in a domain structure the results of Chamberod et al. should be noted. Below s650°C, (e.g. 450°C), Kiis positive and the easy direction is the <100> direction leading to a corresponding domain structure with domains at 90° to each other. Above s650°C (e.g. 760°C), Ki is negative, and the <111> direction is the easy direction, leading to a more complex domain structure, with angles between the magnetisation vectors in adjacent domains of 70°32', 109°28’ and 180°. This may offer a possible explanation of anisotropy in terms of precipitation at high energy points in the magnetic domain structure. The argument is as follows:

At 450°C, the matrix orders, and precipitates are formed in a <100> domain environment, which is then heated to 760°C, a <111> domain environment. The precipitates coarsen and some new precipitates may form at points of high energy in the <111> domain structure. On furnace cooling to room temperature the domain structure again reverts to <100> and the majority of precipitates will be positioned for easy magnetisation in the <100> directions.

A single stage heat treatment at 760°C leads to precipitates forming at prefered sites in the <111> domain structure. When the sample is cooled to room temperature, the precipitate positions are not so favourable for magnetisation in the <100> or <110> direction.

This argument does not hold fully, because the two stage heat treatment would produce the best magnetic properties in the <100> direction; better than the properties of the single stage heat treatment. This is not the case. On its own this change in Ki is not the reason for the magnetic anisotropy but it may be a factor that contributes to another explanation. Further work using magnetic annealing to control domain structures is required.

4.2.3 High aspect ratio precipitates in a crystallographic direction.

If precipitates have di/nwfijyiSflf C * u-fn, where c>a, there will be ~c/a times more interface in the c direction. Hence, although the stored energy per unit volume is the same for any volume at the interface

-176- (s &E£2), the stored strain energy will be dominated by the magnetostriction in the c direction, and the problem is reduced to one dimension. It is known that rod like precipitates with the long axis in the <111> direction exist in Fe50%Co2%V.

Because of the proximity of K1 to zero, and the uncertainty in K2, there is a possibility that the easy magnetisation direction is <111> and not <100>. This is unlikely to be the case, especially as shape anisotropy factors associated with thin sheets, would favour an easy direction in the rolling plane, such as <100>. However, to encompass this possibility, a rod precipitate in a matrix with the <111> easy direction is also considered.

If <100> is the easy direction, the demagnetised state will consist of <100> domains and similarly <111> domains for a <111> easy direction. Let us now consider the change in rod length when the sheet is magnetised in the [100] direction, and when it is magnetised in the [110] direction for both of the starting conditions. The results for a precipitate evenly distributed between the <111> directions on 1110} planes are shown in Table 4.2.4. using the values from 4.2.2.

The rods remain strain free along the c axis in a matrix with a <100> easy magnetisation direction, when magnetised in the [100] direction, but undergo a length change of 20c x 10*6 when magnetised in the flit/, [110] direction. This agrees with^results.

If the matrix has a <111> easy magnetisation direction, the length change is 20c x 10~w6 for a magnetisation in the [100] direction and 25c x 10^6 in the [110] direction. A greater field will have to be supplied to magnetise the sample in the [110] direction which agrees with^results, but the change in precipitate length is much smaller than for a <100> easy direction. Hence a <100> easy direction seems more likely.

-177- Table 4.2.4: Change in strain of a rod in Fe50%Co

Precipitate Easy Applied Initial Final Difference direction direction field strain strain in strain xl0-6 xl0-« xlO-e

<111> <100> [100] 4x 0 4x 0 0 (rod type matrix precipitation) [110] 4x 0 2x 20 20 2x-20 20 mean=20

<111> [100] 3x-13 0 3x13 40 0 40 mean=~20

[110] 3x-13 3x 20 3x33 40 20 20

3x-13 3x-20 3x 7 40 -20 60 mean=25

[110] <100> [100] 2x 35 2x 35 2x 0 (Grain boundary -70 35 105 precipitation) mean=35

[110] 2x 35 2x 65 2x30 -70 65 135 mean=65

<111> [100] 2x 20 2x 35 2x15 2x-20 2x 35 2x55 mean=35

[110] 2x 20 2x 65 2x45 2x-20 2x 65 2x85 mean=65

-178- 4.2.4. Elongated precipitates formed in the rollydirection.

It is in accordance with previous workers that precipitates form on grain boundaries, and in a heavily rolled sample, lines where 3 grain boundaries meet are found in the rollijdirection. Table 4.2.4. shows the change in strain, for a [110] precipitate when magnetised in the [110] and the [100] directions from initial <100> domains and from initial <111> domains. A change in precipitate length of 30c x 10-^ is observed for both <100> and <111> easy directions.

4.2.5. The most likely explanation of experimental results.

Clearly an orientated precipitate is producing magnetic anisotropy. The orientations that could lead to the observed results are <111> or <110>. To select between the two, the following criteria are considered: i) Have precipitates with this orientation been observed? ii) Is there a suitable mechanism to induce precipitation in this orientation? This mechanism must operate at 450° but not at 760°C.

In this work, globular precipitates have been observed by SEM and TEM, at grain boundaries in Rotelloy 5 for both heat treatments, but no significantly elongated precipitates. Figure 4.2.1. This result is in accord with Pitt and Rawling’s work on Fe50%Co-2%V-Ni. There is a tendency for the precipitates to be in rows in the roll direction, but ilie M M * ' more precipitates in the single 760°C/2hrs heat treatment. Magnetic measurements in the transverse direction will clarify the [110] precipitate explanation of anisotropy.

Because the magnetic anisotropy is a recent discovery, the <111> precipitates have not been looked for in the Rotelloy 5 samples. A comparison between the single stage and two stage heat treatments needs to be undertaken. However, general precipitation in cold rolled Fe50%Co2%V and Fe35%Co base alloys at temperatures below Tc has been looked for using TEM. Figure 4.2.4. shows how the dense dislocation tangles obscure any fine precipitates that may be present.

-179 Figure 4.2.4: TEM micrograph showing how dislocation tangles obscure observation of precipitation, (xiOk)

-180- For observation of <111> precipitation we must turn to the work of Mahajan, Pinnel, Olsen and Bennett at the Bell Telephone Laboratories, working on FeCo3%V (Remendur) and of Ashby et al. working at Imperial College on FeCo2V (Permendur). According to Ashby et al., at annealing temperatures below Tc, precipitates are observed on grain boundaries, dislocations and subgrain boundaries in 55% cold rolled FeCo2%V sheet. For annealed sheet, an additional low temperature precipitation occurs in FeCo2%V. Rod like precipitates form in the <111> directions. An excellent illustration of this is given by Ashby et al.1977, who attribute its formation to a low lattice misfit of 1.4% for [lll]a to [0-11]t. A similar result was reported by Mahajan et al. 1974, who performed the following heat treatment on heavily cold rolled FeCo3%V; 1100°C/24hrs and air cooled, 900°C/9hrs and quenched, 600°C/2hrs and air cooled. The same research group found <111> precipitates <0.1pm long in FeCo3%V after >90% cold rolled material underwent 1100°C/lhr and air cooled, 900°C/^hr and quenched, 600°C/2hrs and air cooled (Pinnel et al.1976).

Assuming that <111> precipitates are responsible for the magnetic anisotropy, we can speculate on why they are present in the Rotelloy 5 with the additional anneal at 450°C/1.2hrs, and not in the single anneal at 760°C/2hrs.

According to Ashby et al. <111> precipitates are stable because the small misfit between precipitate and matrix, can lower the total surface free energy. At 450°C the matrix can order and APB’s will be produced. If some of these boundaries contain the <111> direction, there will be a further lowering of the free energy, if precipitation occurs on the APB’s, and hence a greater driving force for precipitation that is not present at 760°C.

Alternatively it may simply be a question of atomic mobility. At 450°C the atoms are not as mobile and although it is energetically more favourable to form precipitates at the grain boundaries, the vanadium atoms may be too far away and lower the free energy by a lesser amount by forming precipitates on matrix planes and in <111> directions.

-181- A third possibility is that dislocations are still present at 450°C that would rapidly disappear at 760°C due to recovery and recrystallisation. Screw dislocations in the <111> direction on the 1110} planes may form nucleation sites.

4.2.6. Summary

Both single and two stage anneal^ {100}<110> Rotelloy 5 samples have similar amounts of spherical precipitate, mainly at grain boundaries which results in identical hardness values and identical reductions in saturation magnetisation from the Fe50%Co value. The two stage Rotelloy 5 is postulated to have an additional fine precipitate in the <111> direction which leads to magnetic anisotropy due to magnetostriction. This effect is greater if the easy magnetisation direction in Rotelloy 5 is the <100> and not the <111>. As both samples have identical crystallographic texture, these <111> precipitates are not responsible for the strong texture.

This effect is not seen in Fe35%Co base alloys, because the texture is not as strong, the magnetostriction is less, and it is not known if the precipitates observed in tempered cold rolled samples, and responsible for grain refinement, are coherent with the Fe35%Co lattice.

In 1953 it was reported by Geisler et al. that after annealing cold rolled Fe50%Co at 454°C for 2 hours, the maximum permeability was at 90° to, and in the rolling direction. 45° to the rolling direction had the lowest permeability, completely contrary to the observed results here. This is further evidence that the ternary additions, and hence the precipitates, are responsible for the magnetic anisotropy observed in the present work.

-182- 4.3.0 Texture discussion.

4.3.1 Factors that determine texture.

The results are in line with many systems in that the textures observed in the present study are dependent on cold deformation, anneal temperature and composition. To be more specific, to produce a strong {100}<110> texture in Fe50%Co base alloys, >92% cold deformation, an anneal at approximately 50°C above Tc for 2 hours, and ternary additions of 2%V or greater are required. In Fe35%Co the same degree of texture was not achieved, but the conditions are thought to be similar but with a lower temperature anneal at 670°C/2hrs, corresponding to ~30°C above Tc. 760°C is too high in the case of 35%Co alloys. Under these conditions, the recrystallised texture is an enhancement of the {100}<110> cold rolled texture.

4.3.2 Explanations of recrystallised texture development.

Much of this explanation of cold rolled and recrystallised textures comes from the recent work of Hutchinson (1988) with bicrystal specimens of a-iron. In all cases recrystallisation was by growth of new grains from grain boundaries. Here it is assumed that the mechanism of texture development in disordered FeCo a-phase is analogous to that in a-iron. As many BCC alloys with a wide range of compositions show similarities in cold rolled and recrystallised textures, the close similarity of FeCo and Fe suggests that this is a very reasonable assumption to make.

4.3.3.1 The same recrystallised texture as cold rolled texture

In a heavily deformed sheet the grain boundaries are predominantly in layers parallel to the sheet surface. New grains nucleate at these boundaries and grow most quickly vertically into the adjacent grain with the greatest stored strain energy. Grains with the {111} planes in the rolling plane have more stored energy than the grains with {100} planes in the rolling plane. The {110} planes have an intermediate amount of stored energy.

-183- Alternatively, no new grain may be nucleated, but at the boundary, the grain with the lower stored strain energy may grow into the adjacent grain. This can be thought of as a migration of the grain boundary, and has been observed in iron by Hutchinson. It is of interest to note that the {100l<110> orientation is known to have a low stored energy of deformation and this often leads to a reluctance for these grains to recrystallise.

For either of these mechanisms to occur the orientation of both grains at the interface must be favourable. These mechanisms tend to enhance the rolled texture.

4.3.3.2 Recrystallised textures differing from the deformation texture.

It is more usual for new grains to be nucleated by subgrains rotated by the deformation process in conjunction with constraints associated with the neighbouring grain. This leads to new grains being nucleated close to grain boundaries. The textures formed in this manner are often a rotation of the deformed texture.

The growth rate of these new grains also determines the final recrystallised texture, and hence this mechanism may be dependent on the recrystallisation temperature. In Fe35%Co0.3%Ta, a change in predominant texture from the cold rolled texture was observed at 750° and 800°C.

4.3.4 Effect of precipitate on texture.

Precipitates are observed at grain boundaries, Figure 4.2.1. in the recrystallised structure.They are globular, and they tend to have an alignment in the roll direction. This suggests they formed on grain boundaries in the deformed structure and have not moved with the growth of new grains.

-184- Pitt, (Pitt and Rawlings 1983) working on similar alloys, concluded that the main role of precipitates is to limit primary grain growth, and not to act as nucleating sites for new grains. The recrystallised matrix grains are not elongated in the roll direction, which might be expected if a limited number of new grains were growing between planar, restNCtive boundaries of precipitation. There must be a sufficiently high nucleation rate of recrystallised grains, for them to impinge on each other and restrict growth in the roll direction to the same degree as precipitates restrict growth in the other directions.

This line of analysis leads to a small grain size, but how does this effect texture? The grains with highest stored energy, the (1111 planes parallel to the rolling plane, will have new grains nucleate and grow most rapidly within them. The nucleation sites may be rotated areas close to the grain boundaries (c.f. 4.3.3.2.) and these new grains will have a different orientation to the deformed host grain, i.e. no longer (1111.

4.3.5. Likely explanation of textures observed in FeCo ternary alloys.

High deformations >92% were needed to enhance the recrystallised texture. The degree of deformation of a grain, to a certain extent, determines its stored energy, and higher deformations ensure that all of the grains are significantly deformed.

There is now the possibility that grain boundaries will migrate into the grains with higher stored energy, as the {1001 orientated grains have a low stored energy; they will grow and this will increase the {100} in agreement with the results. Clearly, this will only occur if the grain boundary is relatively free of precipitation.

The variation in texture with annealing temperature in the Fe35%Co0.3%Ta suggests different grains have different stored energy and a higher temperature is required to induce recrystallisation in grains with different orientations. This reasoning applies to all the alloys investigated, so there must be an additional effect due to the ternary addition. This is associated with smaller grain sizes and implies that precipitates are present before and form during the early stages of recrystallisation.

-185- Smaller grains and the observation of precipitates at grain boundaries, imply that primary grain growth and grain boundary migration is impeded by JJ-phase precipitates. The strong textures observed with higher ternary content are not entirely due to this grain boundary migration mechanism, but due to enhanced nucleation in areas close to the grain boundaries. This would lead to new grains growing in grains of high stored energy.

The strongest (100}<110> texture was in the samples with the smallest grains. Grain size has been reported to affect textures and a reduced grain size will provide a greater driving force for new grain growth and this effect must not be neglected.

4.3.6. Consequence of ternary additions.

c*-ckl*h "iTM A ternary^can reduce the grain size in the hot rolled condition which allows the sheet to be rolled. The smaller the grain size the better the ductility and rolling behaviour. The ternary^also produces a strong texture in heavily rolled sheet but it is not clear how the texture depends on precipitate, and how it depends on grain size. What is clear is that the smaller the grains, the greater is the {100}<110> texture. This crystallographic texture should not affect the magnetic properties in the Fe50%Co base alloys as K1 is close to 0, but may be a problem in the Fe35%Co base alloys if a stronger texture is achieved by the addition of more ternary.

There is a magnetic anisotropy produced due to shape or strain anisotropy of the precipitate particles, which can be reduced by careful heat treatments, (Section 4.2.).

-186- 4.4.0 Hardening in FeCo base alloys.

There is no literature concerning the hardening of Fe35%Co base alloys, 'b ’L'O. so all of the mechanisms discussed below^ are derived from studying Fe50%Co base alloys.

4.4.1 Order hardening.

This has been observed in binary FeCo and Fe50%Co2%V by Chessin (1963), Stoloff et al. (1964) and Moine et al. (1971). The samples which show order hardening are recrystallised, and hence soft and ductile when disordered and harder and brittle when fully ordered at room temperature. Later, cold rolled samples are tested and are obviously initially hard, due to cold work.

The peak in flow stress (in all figures flow stress is at low strains between 0.002 and 0.2%) is between 50 and 70 MPa, and is observed when tensile testing 'at temperature' and when testing quenched samples. It appears at s20°C below Tc for the 'at temperature' tests, and <10°C below Tc if the samples are annealed and quenched. The peaks for FeCo and FeCo2%V are shown in Figures 4.4.1. and 4.4.2.

This peak is attributed to dislocations acting individually as unit dislocations, or in pairs, as superlattice dislocations, depending on the degree of order. The dislocation spacing being inversdy proportional to S2, thus the spacing of the superlattice dislocation at low degrees of order is large, and the dislocations have to be viewed as unit dislocations. This order hardening is observed in other ordering systems such as FeaAl.

Moine et al. (1971) advance the theory, and attribute the hardening peak in FeCo2%V to a change in the mechanism for dislocation generation. Their result is similar to that of Figure 4.4.2. and is shown in Figure 4.4.3. It is seen that neither Stoloff or Moine observe a peak at temperatures below 600°C.

187- 55 I I 1 I 1 I O - FLOW STRESS IN COMPRESSION AT £p = 0.002 MEASURED AT ROOM TEMPERATURE AFTER QUENCHING FROM THE TEMPERATURES INDICATED. 50 500 • - FLOW STRESS IN TENSION AT £p =0.002 MEASURED AT TEMPERATURE INDICATED. ----- COMPONENT OF S T R E SS ASSOCIATED WITH SHORT 4 5 ~ RANGE ORDER IN THE QUENCHED SPECIMENS.

40 400

35

CM e 30 -3 0 0 to UJ MPa >-O 25

to to UJ cc 20 200 to UJZD cc -e - 15

10 100

- 0 0 59 o THE0RETICALLY '||| PREDICTED LONG RANGE ORDER PARAMETER S 0.97 0.9 0.81 0.7 0.48

600 700 800 900 1000 1100 1200 1300 EQUILIBRATION TEMPERATURE, °K

Fig. 4.4.1: Comparison of the as quenched and elevated temperature flow stress measured at 0.002 plastic strain, for binary Fe50%Co. (After Marcinkowski and H. Chessin, 1964).

-188- MPa

Fig. 4.4.2: Comparison of the as quenched and elevated temperature flow stress measured at 0.1% plastic strain, for Fe50%Co2%V (After Stoloff and Davies, 1964).

-189- quenching temperature (°C)

Fig. 4.4.3: Comparison of the as quenched and elevated temperature flow stress for Fe50%Co2%V. (a) samples strained at a crosshead speed of 27 mm/s; (b) samples strained at a crosshead speed of 0.3 mm/s. (After Moine, Eymery & Grosbras, 1971). -190- In all of the results so far precipitates were not seen by light or transmission elecron microscopy. The experimental method did not afford more than a few minutes at temperatures between 400 and 600°C, so precipitates were unlikely to form to any great degree.

4.4.2 Order hardening in cold worked alloys.

Thornburg (1969), Josso (1974) and Couto et al. (1989) annealed cold rolled FeCo2%V samples at temperatures between 550 and 800°C and found no evidence of order hardening, in yield stress results. The work hardening when cold rolling, increases the yield stress by -1200 MPa and the softening associated with recovery and recrystallisation totally swamps the sl5x smaller order hardening peak, Figures 4.4.4., 4.4.5. & 4.4.6.

However, there is an increase in the yield stress from the initial cold rolled value which occurs at temperatures below 600°C and this has been attributed to precipitation hardening.

4.4.3 Precipitation hardening in cold worked alloys.

Pinnel et al. (1976), used a variety of heat treatments on an 48.05%Fe48.8%Co2.80%V alloy, culminating in furnace cooling from 500-600°C. All these samples were said to be fully ordered and hence differences in mechanical properties were attributed to dislocation densities, grain size, texture and precipitation. Order hardening was completely ruled out.

Of particular interest, for samples annealed for 2 hours and then furnace cooled, a peak in the 0.2% yield stress was observed for annealing temperatures between 600 and 750°C. (600°C was the lowest anneal temperature), Figure 4.4.7. They attributed this peak to precipitation, and indeed precipitation was observed by microscopy However, this explanation, as the author admits, is not completely satisfactory, as the precipitate starts to redissolve at a temperature at which the yield stress is still increasing.

-191- inches 2 2 Percent Elongation, Percent Elongation,

Figure 4.4.4: Effect of annealing temperature on the mechanical properties of 90% cold rolled FeCo-2%V alloy. All samples were annealed for 2 hrs and cooled at approximately 900°C/hr. (Thornburg 1969)

TEMPERATURE (K)

Figure 4.4.5: Variation of the 0.2% yield stress (YS) and the ultimate tensile stress (UTS) as a function of heat treatment temperature for 90% cold worked FeCo-2%V alloy. All samples were annealed for 2 hrs. (Couto and Ferriera 1989)

-192- R MPa

Figure 4.4.6: Mechanical properties of 91.7% cold worked FeCo2%-V versus quench temperature. A% elongation; R tensile strength; ^DPN hardness. (Josso 1974)

Legend 2500-i 2 50 A A series; cold-rolled strip A B series; cold-rolled strip □ CondD series 300 O Data from Ref. 9 2000 - >' 20 0 • Data from Ref. 9 E — qO. s O A series; strand onneoled wire J? * A series; cold-drawn wire JC 1500• £ 150- CT CT» 5 200 QJC - f - i n c o •o X3 N 4 * CO t o it: ^ IOO o o 500- 5? C\J ■ h 50- C\J O—< o o o

100 200 300 400 500 600 700 800 900 Aging temperature, °C

Figure 4.4.7: Series A: Yield strength versus aging temperature of >90% cold rolled FeCo-2.8%V annealed for 2 hrs and furnace cooled. (Pinnel et al. 1976)

-193- Thornburg (1969)f had similar problems trying to explain a ductility peak between 650 and 710°C for Fe50%Co2%V. These samples were furnace cooled at 900°C/hrf Figure 4.4.4. Josso also observed this 7% elongation peak at close to 700°C, Figure 4.4.6. These anomqlies may be associated with partial recrystallisation, but this peak is at a higher temperature than the hardness peak , Figure 4.4.6.

The only references to report on mechanical properties of Fe50%Co2%V at temperatures below 600°C are Couto and Ferriera (1989), and Josso (1974). Of great significance is a hardening peak at 550°C. This is too low a temperature and too great a magnitude for order hardening and Couto and Ferriera backed up by microscopy, attribute it to precipitation of C0 3 V, Figure 4.4.8. The temperature range over which hardening is observed is in line with the ^-phase field illustrated by Martin and Geisler (1952), Figure 1.5.1. (page 40). The K precipitates more quickly in a cold worked material and so previous work by Moine et al. and by Stoloff et al. employing short heat treatments on recrystallised material appear not to have produced precipitation.

4.4.4.0 Explanation of hardening in this work.

These experiments were on heavily cold rolled material and it is clear, that precipitation hardening is being observed over this temperature range; order hardening is too small an effect, and occurs at higher temperatures.

4.4.4.1 Variation of hardness with anneal temperature.

The shape of these curves^are similar for both the Fe35%Co0.3%Ta and the Fe50%Co2%V, and the temperature range of the peak represents the (a + Y) phase region in the FeCoV phase diagram of Martin and Geisler. The two phase region in the FeCoTa phase diagram must be at similar temperatures but for lower Ta concentrations. Precipitation is clearly visible at grain boundaries by microscopy. This implies that tantalum is less soluble in Fe35%Co than vanadium is in Fe50%Co. Persiano (1986, p.23), considered the solubility of several elements, and indeed from electronegativity and atomic radii information, tantalum is less soluble in FeCo than vanadium.

-194- TEMPERATURE(K)

Figure 4.4.8: Microhardness of 90% cold rolled FeCo-2%V alloy as a function of heat treatment temperature. (Couto and Ferriera 1989, 2hr anneal, Josso 1974, lhr anneal)

TIME (hours)

Figure 4.4.9: Isothermal hardening of 90% cold worked FeCo-2%V alloy. (Couto and Ferriera 1989)

-195 The precipitation, which is responsible for the low temperature hardening effect, may also account for the difference in magnetic anisotropy between Rotelloy 5 samples with and without additional anneals in this two phase region (Section 4.2.). It is interesting that the hardness increase m the Fe35%Co0.3%Ta^is only slightly less than alley that in the Fe50%Co2%V, which could mean either that a similar amount of precipitate comes out of solution and most of the vanadium remains dissolved, or, that the precipitates are of the optimum size for hardening, (although their volume fraction is less in the case of the Fe35%Co alloy). At higher temperatures the material softens due to a

combination of dissolution of some of the precipitate, (according to n Martin and Geisler) and coarsening of the precipitates and recrystallisation (according to Couto and Ferreira).

4.4.4.2 Variation in hardness with isothermal annealing.

After annealing at 550°C, the Fe35%Co0.3%Ta samples reach maximum hardness after 2 hours. In the microhardness tests on Fe50%Co2%V samples there is a hint of a small peak in hardness after 1% hours but this may be due to experimental scatter. The levelling off in hardness indicates that substantial recrystallisation, and/or precipitate coarsening is not occuring in under 6 hours. ^ ) ( ^ •'S After annealing at 650°C,jt>the Fe35%Co0.3%Ta alloy behaves in a similar manner to that at 550°C, though a systematic temperature error may be present as recrystallisation softening would be expected and is indicated by the hardness vs temperature curves; i.e. the anneal temperature was closer to 630°C (see Section 3.2).

The more pronounced hardness peak below 1 hour, in the microhardness curve for isothermal annealing at 650°C in Fe50%Co2%V is more difficult to explain. It is present to a lesser extent in the macrohardness measurements but completely absent from the Fe35%Co0.3%Ta results. The same effect was noticed by Couto and Ferreira, Figure 4.4.9. The hardening side of the peak was attributed to precipitation as before and the softening side of the peak was attributed to coarsening of precipitates. This was illustrated with TEM. This explanation would fit the present results^ fche coarsening not being observed in the Fe35%Co0.3%Ta because of the limited amount of ternary present.

-196- A more detailed and speculative argument could go as follows.

Fe50%Co2%V: At 650°C the alloy rapidly becomes ordered and globular precipitates form on grain boundaries while elongated precipitates form coherently with the matrix in the <111> direction. This is the hardest structure, with dislocation movement being hindered within grains and at grain boundares. As the time at this temperature is increased, both the elongated precipitates and the globular precipitates coarsen. The interiors of the grains are now softer. This softening will be enhanced by recovery, but large scale recrystallisation would be reflected by a softening ko below the start condition, which is not observed.

Fe35%Co0.3%Ta: Here the precipitation of Y phase is not coherent and precipitation will occur at grain boundaries and on dislocations, however, there is not enough ternary to allow for substantial coarsening.

-197- 4.5.0 DSC Discussion.

4.5.1 The derivation of the threshold temperature for atomic movement.

Now that it has been established that the movement of atoms to form an ordered sample is being detected by the '550 anomaly' peak, a threshold temperature, Td, for atomic diffusion can be found by extrapolating the peak temperature to a heating rate of 0°/min, Figures 4.5.1. to 4.5.3. If the order/disorder temperature, Tc, is also known, then there is a temperature range, Tc - Td,j[which the hot rolled alloy must be quenched through in order to retain disorder. Table 3.6.2 shows the results for a selection of alloys:

Table 4.5.1: Threshold temperatures and ordering ranges.

Sample Tc (°C) Td °C) Tc - Td Eact 20°C/min 0°C/min (°C) kJ/mol

Fe50%Co 720 550 Fe50%Co2%V 710 524 488 222 195

Fe35%Co 640 589 Fe35%Co0.3%Ta 640 580 520 120 140 Fe35%Co0.3%Ta4%Ni 635 566 525 110 135

Errors are large in Tc - Td and are due mainly to the bluntness of the Td peak but it is clear that the temperature range over which the Fe35%Co base alloys can order is almost half that of the Fe50%Co alloys. The temperature range is also at a higher temperature for the equiatomic base alloys and hence faster kinetics of ordering would be expected. All in all, from this data it would be expected that disorder is easier to retain in the Fe35%Co alloys.

-198- mJ/sec iue451 Eprmna S uvso '5504.5.1:curves forFigureanomqly'ofDSC Experimental Fe50%Co2%V atdifferentrates.Fe50%Co2%Vheating m J /s e c /m g iue452 Eprmna S uvso '5504.5.2: forFigurecurvesanomaly' ExperimentalofDSC Fe35%Co0.2%Ta atdifferentrates.Fe35%Co0.2%Taheating Heating rate (deg/min) ramp rates give ramp rates to temperatures. ordering threshold Figure 4.5.3: of different Extrapolation peaks at ordering Figure 4.5.3: Peak temperature (°C)

- 201 - 4.5.2 Ordering activation energies from Td peak vs Heating rate graphs.

Kissinger 1957 and Augis and Bennett 1978 produced, and reviewed methods of determining the activation energy of phase changes in metals. The basic rate equation is assumed to be of the form

dX = A(l-X)n exp(-E/KT) (4.5.1) dt where X is the fraction of the material reacted, n is the reaction order, and the other symbols have their usual meanings. The reaction order, n, determines the shape of the peak and is often determined by fitting the experimental peak to known shapes corresponding to different n values. In the Kissinger method the value of n is assumed to be 1, and a plot of In(heating rate) vs In 1/Td gives a straight line of gradient -E/K. Augis and Bennett reviewed Kissingers method and conclude that the peak position is independent of the peaks shape and hence n. They plotted In(rate) vs l/(T-To), where To is the temperature of the sample at the start of the experiment, which is valid for any n; a very similar value for the activation energy but a .'a ic-mfLt ryuJaLUc*fr different intercept. Due to the errors^present in this newly developed technique the original Kissinger method is sufficiently accurate to give E values for ordering.

Figure 4.5.4. shows these graphs for two Fe35%Co base alloys and FeCo2%V. The activation energies are lower for the 35%Co alloys than for the 50%Co alloy and are given in Table 4.5.1.

4.5.3. Mathematical modelling of low temperature ordering.

Probably the most important consequence of being able to monitor this 550°C anomaly is the information it will yield on ordering kinetics. The effect of ternary additions on the kinetics can then be investigated in detail. The theories of ordering kinetics have been developed by many workers since Bragg and Williams 1934, and Bethe 1935. Some of the more recent models applied to FeCo under non equilibrium conditions are reviewed here.

-20 2 - -£0 Z- ln(rate) Figure 4.5.4: Graph tothe4.5.4:Figureenergiesactivation Graphobtainordering by theKissingerby method. /Pek temperature) eak l/(P M E T H O D 1

Yokoyama et al 1971, applied Bragg and Williams theory. If an alloy is at a temperature, T, whose degree of order is slightly different to the dynamic equilibrium at temperature T and corresponds to the equilibrium degree of order at a different temperature 6, then, they assumed that the rate of approach to equilibrium is given by:

de/dt = -(0 - T)/E (4.5.2) where E, the time of relaxation, is the time taken for the departure from equilibrium to be reduced by 1/e th of its initial value. The time of relaxation obeys:

£ = A exp(E/kT) (4.5.3) where A is 10-1213, E is the activation energy for atomic interchange and taken as 2.33eV (225 kJ/mol) in this paper, k is Boltzman's constant.

So as to make future comparisons more easy, equations 4.5.2, and 4.5.3 can be rearranged. If $ is the heating rate, then:

de/dt = de/dT . dT/dt

= de/dT . $ (4.5.4) and hence 4.5.2 and 4.5.3 can be combined by substituting for E, and using 4.5.4, to give:

de/dT = 1/A$ . exp(-E/kT) . (T - 0) (4.5.5)

This equation has to be solved numerically by a Runge-Kutte type calculation.

- 2 0 4 - M E T H O D 2

Due to language difficulties it is not possible to give a detailed account of the theory of Tahara et al. (1983). However the Bragg and Williams approach of method 1 seems to be followed fairly closely. The two starting equations in this case are:

dS/dt = 1/2 . F(T) .S. (Se2- S2) (4.5.6) where S is the long range order parameter, Se is the equilibrium value of S at temperature T, and:

F(T) = B exp(-E/RT) (4.5.7) where B is the jump attempt frequency equal to 0.97 x 1013/s, E is the activation energy, equal to 230 kJ/mol., and R is the gas constant. These equations combine to give:

dX/dT = B/4> . exp(-E/RT) .X. (Xe - X) (4.5.8) where X = S2, Xe = Se2 and $ is the heating rate. Tahara showed that dX/dT can be related to resistivity, R, and specific heat, Cp.

dX/dT = const. /\£lk (Tahara) (4.5.9) and:

£

Where

- 2 0 5 - Temperature/ K

Figure 4.5.5: Variation of the temperature derivative of electrical resistivity caused by the ordering or disordering, a

-206 Temperature / K

Figure 4.5.6: Variations of a

-207- M E T H O D 3

This is the model used by Matsuda et al. (1971). They extended the theory of Bragg and Williams to a binary alloy containing vacancies. It was then assumed that atom movements occur by two mechanisms, a direct interchange of atoms, and an indirect mechanism involving an interchange through vacancies.

The difference between Method 3 and Methods 1 and 2 is that Methods 1 and 2 assumed that it took a certain time for the alloy to reach equilibrium and values for this time were found by experiment. The mechanisms of atomic diffusion were not discussed in detail. This model starts from an atomic level and considers the change in energy when an atom or vacancy moves from one lattice point to another. This can occur via a direct interchange of atoms or by interchange via a vacancy. In iron cobalt alloys Matsuda found that both mechanisms needed to be taken into account for his model to fit the experimental data. The equilibrium degree of order was obtained by minimising the free energy.

The results of this model agree with those of Yokovama (1971), and are similar to those of Takagi and Oguchi (not reviewed here).

In addition, this was the first paper to consider iron cobalt binary alloys of non equiatomic composition and the '550 peak' temperature for these alloys was predicted.

METHOD 4

Chamberod et al. (1972) also approached the problem from an atomic level. The Yang-Li-Hill approximation was used because it took into account, not only the nearest neighbour bonding, but the next nearest neighbour bonding. If only the nearest neighbour bonding is considered, the results, like those of Matsuda et al, are analogous to those of Bethe (1935) and of Takagi (1956).

The aim of this model was to explain theoretically the anisotropy in equiatomic binary iron cobalt, but it also provides an equilibrium curve for L.R.O. and for specific heat (at constant volume), versus temperature. In attempting to fit theoretical to experimental results

- 2 0 8 - for magnetic anisotropy, it became necessary to vary the bond energy linearly with temperature. Having done this the theoretical fit was very good.

A useful equation used later and which relates the difference in bond energy between the ordered and disordered states, v, to Tc is:

Tc = v/ln(z/z-2) = v/ln(4/3) (4.5.11) where v = (v(aa)+v(bb)-2v(ab))/2 and z (=8) is the coordination number. Hence a lower Tc implies a smaller difference in bond energy between the ordered and disordered state.

Because this model inherently applies to equilibrium conditions the 550 anomoly is not described and does not appear.

METHOD 5

The work by Sato and Kikuchi in 1976, is probably the most accurate model (certainly the most complicated) of the iron cobalt binary system. Again a pair approximation is used based on the vacancy mechanism of atomic migration and using the path-probability method of time-dependent cooperative phenomena. Perhaps the major advance over previous reports is the incorporation of more than one relaxation time, which is neccessary to explain the overshooting of the equilibrium state. Unfortunately their computer experiment was much simplified because of the excessive computer time required for accurate calculations, and hence the results are only qualitative.

In this work the equilibrium condition was determined by minimising the Grand Potential, not the Helmholtz free energy as in previous papers. The Grand potential is the Helmholtz free energy plus a term representing the chemical potential. The order/disorder temperature was found to be the same as equation (4.5.11), though it is pointed out that this equation is only approximate as it does not account for the presence of vacancies.

The authors then investigated the variation of the system's state with time when the system was in a non equilibrium starting state. To do

-209- this, 4 variables were introduced, 2 were used to describe the long range order and 2 were used to describe the short range order.

Associated with these 4 variables describing the state, are 4 relaxation times. If the system is started in a non equilibrium state, then each variable will relax to the equilibrium state in an exponential manner described by its respective relaxation time. This assumes that the temperature is held constant and that the relaxation times are constant for small deviations from the equilibrium conditions.

The relaxation times are obviously temperature dependent, (c.f. E a exp(E/kT), eqn. 4.5.3.), but the authors go on to stress that the relaxation times are also affected by the degree of order, and call this the cooperative effect.

The results of this computer experiment do not match the shape of experimental results as well as some of the other models but it does afford an explanation of the following effect.

Suppose, upon cooling at a certain rate, the system freezes into a state with a certain value of thermodynamical variable (say the volume VI) at low temperatures. The corresponding temperature (fictive temperature) T1 is sought; this means that the equilibrium volume at T1 is VI. Suppose that the system is rapidly heated to T1 and is kept there isothermally. When the isothermal change of volume with time is measured at Tl, it is observed that V first increases from VI and then decreases back to the equilibrium condition, VI, instead of remaining constant.

4.5.4 Fitting a theoretical curve to the experimental data.

For this the method of Tahara, Method 2, is used to analyse the present results, in particular use is made of the expression:

dX/dT = B/4> . exp(-E/RT) .X. (Xe - X) (4.5.8)

Firstly, we need to use Oyedele and Collins neutron diffraction data (c.f. Section 1) to give the equilibrium value of order with

- 2 1 0 - temperature for Fe50%Co and Fe35%Co:

Se = (1-T/Tc)° (4.5.12)

Xe = Se2= (1-T/Tc)2P (4.5.13)

Xe can now be calculated. The values of E derived from the experimental results using Kissinger's Method were used. 4> is the heating rate and is obviously known, which leaves B, X and T. The starting condition of the sample is of very low order at room temperature, so we can put in a value of X=0.05 at T=300K for Fe50%Co and X=0.025 for Fe35%Co and use Runge Kutte to provide a theoretical curve for different values of B. The start values of X were chosen to be in the same ratio as maximum order squared for Fe50%Co and Fe35%Co, i.e. 12:0.72. These start values of X, were arbitrarily selected, and the curve shape was sensitive to these values. Too small a value of X and dX/dT remained close to zero throughout, too large a value and the initial difference Xe-X became significant. In this case, B was adjusted to give a best fit for a ramp rate of 20°/min (0.3333 K/s) and then the curves, and hence peak temperatures, for different ramp rates were drawn, keeping the same value of B. A spreadsheet was used for this calculation and Table 4.5.2. shows the parameters used. Figures 4.5.7. and 4.5.8. show the calculated graphs of dX/dT vs T.

Table 4.5.2: Parameters for best fit curves.

Fe35%Co(1) Fe35%Co(2) Fe50%Co Order/disorder temp. (K) 908 908 983 Critical exp. (i 0.315 0.315 0.308 Jump frequency, B (/s) 2.25 x 107 8.6 x 1010 5.0 x 1011 Act. energy, E (kJ/mol) 140 195 195 Runge Kutte step (°C) 1 1 1 Start X 0.025 0.025 0.050

The activation energies for Fe35%Co(l) and Fe50%Co of 140 and 195 kJ/mol, are derived from applying Kissinger's method to the experimental results. Fe35%Co(2) uses the value*irr of 195 kJ/mol to see if the poor agreement with the experimental results^could be improved.

- 2 1 1 - .2-1 5 6 '/ ) Y \ 'M

Figure 4.5.7: Calculated curves of dX/dT vs temperature for Fe35%Co using Tahara's model.

Figure 4.5.8: Calculated curves of dX/dT vs temperature for Fe50%Co using Tahara's model. - 2 1 2 - Table 4.5.3. shows the temperatures of the calculated peaks and the Kissinger energy obtained from these theoretical peaks.

Table 4.5.3. Calculated *550 peak* temperatures.

Ramp Fe35%Co(1) Fe35%Co(2) Fe50%Co Rate $°/s Theory Expt.(°C) Theory Theory Expt.(°C)

5 535 (-4) 539 543 (+4) 10 559 (+19) 540 562 (+22) 507 (+10) 497 20 580 (0) 580 580 (0) 524 (0) 524 30 589 (-15) 604 589 (-15) 534 (+7) 527 40 594 603 542 50 548 (0) 548

'Input energy’ 140 195 195 'Output energy’ 177-205 206-216 209 (kJ/mol)

The difference between experiment and theory is shown in brackets. The input^energy is the value shown in Table 4.5.2. which was used to produce the fitting curves, and the output energy is the activation energy given by a Kissenger analysis of the theoretical peaks.

These calculated peak temperatures can be put into the Kissinger equation to see if they give a straight line, and if the gradient of that line gives the same value as that used to generate it. Figure 4.5.9. shows that the lines are fairly straight and that the Fe50%Co gradient represents E=209kJ/mol. (c.f 195 kJ/mol. expt.). The activation energy for the Fe35%Co alloy is 177-205kJ/mol which is not self-consistent with the energy used to obtain it, (c.f. 140kJ/mol. expt.).

Equation 4.5.8. does fit the experimental peaks quite well for the Fe50%Co alloy, as can be seen from the difference in brackets. The fit is not very good for the Fe35%Co base alloy and the line of/Figure 4.5.9. is not straight. This may be because the activation energy value

-213- Figure 4.5.9: Activation energies derived fromthecalculated4.5.9:Figureenergiesderived Activation DSCpeaks Ln (HEATING RATE) usingMethod. Kissinger’s -214- was incorrect, so a value of 195 kJ/mol., the same as the Fe50%Co base alloy, was tried with a corresponding value of B, shown in Table 4.5.2. Again, Table 4.5.3. shows that the fit has not improved.

4.5.5 Why £ derived by Kissinger method does not agree with E derived from fitted equations.

The Kissinger method originates from an equation of the form:

dY/dt = A (1-Y) exp(-E/RT) where Y is the amount of material reacted and is between 0 and 1. Tahara's equation written in the same form with respect to t and hence losing the heating rate, $, is:

dX/dt = B X(Xe - X) exp(-E/RT)

Here Xe is a function of T. In our use of the Kissinger method we have assumed that X(Xe-X) behaves in a similar manner to (1-Y), and this, along with factors in Section 4.5.6. is responsible for the inconsistency.

4.5.6 Reasons why theoretical fit to Fe35%Co base alloys is poor.

Firstly, the experimental peaks are small and, hence there position is not accurately determined. It has been shown that they exist, and as the technique is developed further, improved resolution will be achieved (see Section 5).

As far as the the mathematical fitting goes, Tahara’s equation does not account for changes in vacancy concentration, which he himself mentions in other papers, (c.f. Girifalco 1964, Tahara 1978), so a Cv(T) term is required. This is not a simple term but consists of exponentials and energies of vacancy formation and migration.

Chamberod points out that thermal vibrations must be accounted for at these temperatures, and reports that the difference in bond energy, v varies with temperature in the Fe50%Co binary as v - 3/2QkbT where Q is found by experiment, and = 0.467 for the equiatomic binary. If this

-215- change in bond energy involves a greater weakening of the v(aa) and v(bb) than of the v(ab) bonds, then there must be a corresponding change in the jump attempt frequency, and hence B is a function of temperature. Both of these effects will become greater at higher temperatures, and hence the higher temperature peaks of the Fe35%Co base alloys will be more greatly affected than those of the Fe50%Co alloys. A more complete equation is of the form:

dX/dt = B(T).Cv(T).X(Xe(T) - X) exp(-E/RT) (4.5.14)

4.5.7 Shape of the theoretical curves.

Figures 4.5.7.& 4.5.8. show the theoretical curves for Fe35%Co and Fe50%Co. When comparing them to the experimental results in Figures 4.5.1 & 4.5.2, there is a glaring discrepancy in the shape of the curves. The theoretical ones get smaller with heating rate, the experimental ones get larger. The theoretical curves plot dX/dT vs T, which is equivalent to specific heat, Cp vs T, whereas the experimental curves plot power supplied vs T, or J/sec vs T. To convert this to Cp vs T, the power must be divided by the heating rate, or, equivalently, as far as fitting goes, the theoretical curves can be multiplied by the heating rate. As can be seen from Figures 4.5.10.& 4.5.11., the fit is considerably improved.

4.5.8 Lattice parameter and diffusion energy reinforcement of lower ordering energy.

The most significant result of this work is that the Fe35%Co orders more rapidly than the Fe50%Co base alloys in agreement with Eymery et al. (1974), and over a narrower temperature range, and this is attributed to a lower activation energy. A lower activation energy may have been expected for other reasons. The lattice parameter is larger for the Fe35%Co base alloys, implying that atoms are being held together less tightly. In addition, Hirano and Cohen, 1972, give a lower diffusion energy in the BCC phase for cobalt in Fe29at%Co (Q=47kcal/mol =197 kJ/mol) than in Fe50at%Co (Q=59kcal/mol =248kJ/mol).

-216- iue451: acltdseii et curves4.5.11:Figurefor CalculatedspecificFe50%Co. heat

SPECIFIC HEAT (ARBITRARY UNITS) % SPECIFIC HEAT (ARBITRARY UNITS) 4.5.10: Calculated specific heat curves4.5.10:for Calculatedheatspecific Fe35%Co. -217- However, not all workers agree with this lower value of activation energy for ordering in Fe35%Co binary alloys. In particular Yokoyama et al. (1971), used resistivity measurements to measure order in a range of FeCo binary alloys, and found that Fe33at%Co had the greatest change in resistivity between the ordered and disordered states and used the greatest change in resistivity with temperature and an Arrhenius equation to deduce that the activation energy for ordering was higher in the Fe33at%Co binary (255kJ/mol.) than in the Fe50at%Co binary, (214kJ/mol.). If this value of E is used in Tahara's model for Fe35%Co, the peaks occur at 550, 565, 580, 588 & 594 °C for heating rates of 5,10,20,30 & 40°C/min, B=5.8xl014 and Kissingers method gives a good straight line with a gradient corresponding to 280kJ/mol. Unusually, he reports that the resistivity of the Fe33at%Co alloy increases upon ordering but that of the Fe50at%Co decreases upon ordering.

Asano et al. (1966), studied the position of the specific heat relaxation peak of a fully ordered sample approaching the equilibrium order, the opposite of the relaxation of the disordered state to the equilibrium order state studied in this project. He reports that the peak is independent of the composition on the iron rich side of the binary alloys. Matsuda fitted his mathematical model to Asano's data and found good agreement.

4.5.9 Consequences of finding a full mathematical model of the system.

The ultimate aim of this type of investigation is to form a complete mathematical model of the system, with a justification for every term that occurs in the equations. The position of low temperature disorder to order peaks could then be used to find the values of parameters such as E & B. If the model is accurate, a negative heating rate (i.e. cooling rate) could then be used with a starting condition in the disordered a phase field. It would then be possible to predict the degree of order in a sample cooled at a known rate.

Unfortunately the '550 peaks' are not yet well enough defined in the 35%Co alloys and Tahara's equation is not yet sufficiently accurate for this type of prediction.

-218 5: Conclusions

5.1 Conclusions and implications.

5.2 Factors that determine cold reliability.

Fe35%Co based alloys can only be cold rolled if the hot rolled material is quenched to retain disorder, and if the grain diameter is kept to less than s25pm. Larger grain sizes can be rolled if the sheet is warmed to slOO°C.

Grain sizes can be reduced in the hot rolled sheet if a ternary^is added such that fine preciptates are formed at the grain boundaries of the hot rolled sheet. V, Nb, Ta form C03X precipitates as well as compounds of C, N, 0 and H. It has been shown that an increase in ternary addition or a quaternary addition, results in a reduction in grain size. This is illustrated by the addition of Ni to both Fe50%Col.8%V and Fe35%Co0.3%Ta base alloys which results in smaller grain sizes, and an easily cold rolled material.

Ternary additions may also enhance ductility by interacting with interstitials such as C, 0, N, and H that are detrimental to reliability because of their impedance to dislocation motion. These interstitials can be removed from solution by the formation of compounds with V, Nb or Ta and thus impart ductility to the alloy.

Further work.

Precipitates of V, Nb and Ta compounds improve ductility and V in, particular increases resistivity; however, the precipitates are paramagnetic and reduce both susceptibility and saturation magnetisation. In a.c. applications where the components are not under demanding mechanical stresses, such as stators in motors or generators, it would be an advantage to have a small grain size during production but a larger grain size to enhance the magnetic properties in the final component.

To achieve this, a precipitate could be used to restrict grain growth during production but then be removed in a heat treatment of the final component. MnS is added to Fe3%Si steels to control grain size and texture during production but the sulphur is removed by annealing in

-219- hydrogen leaving the Mn in solution. The advantage to FeCo would be two-fold. Not only would the larger grain size improve susceptibility, but Mn in solution has a higher magnetic moment than Fe or Co atoms and aligns parallel to the matrix, thus raising saturation. In FeCo alloys the removal of sulphur may be difficult because temperatures in the gamma phase field may be required.

5.3 Rolling and recrystallised textures.

The two main cold rolled and recrystallised textures of iron cobalt alloys are {100l<110> and 11111<211>. For high 092%) degrees of deformation and if precipitates and consequently small grains are present, the 11001<110> component can be increased; in the case of FeCol.8%V3.5%Ni, an almost pure 11001<110> texture can be achieved.

Annealing temperature is important in determining the final recrystallised texture. The strongest 11001<110> textures were achieved at s50°C above Tc. In the Fe35%Co0.2%Ta alloys higher annealing fieM temperatures m the BCC phase/resulted in a stronger 1 1111<211> component at the expense of the U00}<110> component.

5.4 Precipitation hardening in cold rolled alloys.

Between 400 and 650°C both the Fe35%Co and Fe50%Co alloys harden due to the formation of a precipitate. Above 650°C the alloys soften due to a dissolution of the precipitate. On prolonged annealing within the precipitate and BCC phase field, coarsening occurs which also results in softening. This is important because these alloys have a high Curie temperature and are often used in high temperature applications, where precipitation will occur resulting in a change in both magnetic and mechanical properties.

5.5 Saturation magnetisation.

The 35%Co alloys all have higher saturation magnetisations than 50%Co alloys with similar ternary additions, and in the cold rolled, annealed and quenched condition the Fe35%Co alloys are inherently more ductile than the Fe50%Co base alloys. Greater ternary additions can be added to the Fe50%Co and Fe35%Co alloys to produce improved ductility via grain

- 2 2 0 - refinement with the disadvantage of greater precipitation of a paramagnetic phase which decreases the saturation. Again there is a compromise here. Ductility enables the rotors of aircraft generators to operate under greater centrifugal stresses and hence leads to greater power production.

5.6 Magnetic anisotropy.

In Fe50%Col.8%V3.5%Ni (Rotelloy 5) 96% cold rolled sheet, annealed at 450° and then 760°C, with a strong {100}<110> texture, a significantly lower permeability and greater coercivity was observed in the <110> roll direction compared to 45° to the roll direction, i.e. the <100> direction. The magnetic properties of the 45° sample were comparable to those of a thicker randomly orientated sample of identical composition but sheet with identical texture from a single 760°C anneal did not exhibit these characteristics.

It is suggested that this effect is due to precipitation. This effect is only present if the precipitates are formed during a preliminary anneal of 1.2hrs at 450°C. Further work is required to ascertain whether the precipitates have to nucleate in the <100> domain environment which exists up to 650°C or if the ordered matrix at 450°C is required.

This work would involve magnetic annealing to control the domain environment and TEM to observe the domain structure and determine the shape of precipitates. Magnetic properties in the transverse direction must be compared with those in the roll direction to determine if the anisotropy is a consequence of rolling or crystallographic texture.

This effect is observable to a lesser extent in the Fe35%Co alloys but these alloys do not have the same strength of texture.

Further work.

If orientated precipitates can be retained in the final product such that they are influential in increasing resistivity perpendicular to the field and improving mechanical stength, but not creating magnetostrictive stresses at the precipitate matrix interface leading

- 2 2 1 - to domain wall pinning, then greater generator efficiencies would be acheived. Clearly resistivity measurements in different sheet directions are required.

At higher operating frequencies higher initial permeability is required and this has been achieved in Supermendur through magnetic annealing. Further improvements are possible if the crystallographic easy direction is aligned with the annealing field.

5.7 Methods of measuring LRO in Fe35%Co alloys.

Differential thermal calorimetry.

DSC has successfully been used to monitor the change in LRO from a metastable disordered state to the equilibrium state of order, and the Fe35%Co alloys order more rapidly. The activation energy for ordering in Fe50%Co2%V of 195kJ/mol, is in close agreement with previous workers. In the Fe35%Co base alloys the peaks were not as well defined. The results do show that the activation energy for ordering is lower in the Fe35%Co base alloys (sl40kJ/mol) but the technique needs to be developed further. Samples have been cycled at various heating rates and the ’550 anomqly' is still observable. A further improvement would h? be^monitor the 550 peak for different heating and cooling rates.

Once the technique has been refined the possibilities for determining the effect of ternaries on ordering kinetics are extremely promising. The compositions and quench rates necessary to preserve the required degree of order could then be calculated.

X-ray lattice parameter and superlattice line measurements.

The change in the lattice parameter of Fe35%Co alloys on ordering is too small to be used as a method of determining order.

Similarly the 100 superlattice line of the Fe35%Co base alloys can be detected using Co Ka X-rays, but only for high degrees of order. It is not possible to measure low degrees of order with any degree of accuracy and hence, this method is not suitable for determining isothermal ordering rates, and ordering activation enegies for alloys

- 2 2 2 - with different ternary additions.

Resistivity.

It has been reported by Yokoyama (1971) that resistivity is sensitive to the degree of order in the Fe35%Co base alloys. His results were not in agreement with those of this report, and it was not clear if changes in the microstructural features had been accounted for. Further work should include an assessment of using resistivity to determine LRO in the Fe35%Co alloys.

6.1 Suggestions for further work.

6.1.1 Improvements in magnetic properties for DC applications.

Here the saturation magnetisation is paramount and in the Fe35%Co base alloys can only be improved upon by the addition of manganese atoms in solution. Grain size is largely unimportant.

6.1.2 Improvements in magnetic properties for AC applications.

In static components a large grain size would lower the grain boundary density and hence would facilitate domain wall movement. Here there is a difference between the Fe35%Co and the Fe50%Co base alloys. The domain structure is in part determined by the magnetostriction of the material which is greater in the Fe50%Co material than in the Fe35%Co material. For optimum coercivity and permeability, a critical domain size is required and a study of the Fe35%Co domain structure is necessary. If appropriate, the domain size can be reduced while retaining large grains, by the introduction of surface scratching. This relatively new field used in silicon steels is not yet fully explored and if a large grain size is achieved it may be appropriate to FeCo alloys.

Amorphous melt spun ribbon has no grain boundaries to impede domain wall motion and would be a natural development of this domain control technique. Large domains are likely to result which can be broken up by surface scratching. Melt spinning would also avoid the problems encountered with cold rolling.

223- Alternatively, the ribbon could be powdered and sintered in a similar manner to the production of rare earth permanent magnets. The iron - cobalt printer head mentioned in the introduction, was constructed by sintering. Magnetic alignment of the powder may also be appropriate.

6.1.3 Alternative methods of observing the '550'anomqly.

It is clear that any further investigations of the Fe35%Co base alloys would benefit from a quick and inexpensive method of long range order determination. The dynamic thermal cycling method of the DSC, developed in this thesis shows promise. However, rather than measuring specific heat changes upon ordering, greater sensitivity may be achieved by monitoring a change in mechanical properties upon ordering.

A new technique called Dynamic Mechanical Thermal Analysis, DMTA, offers just such a possibility. Originally developed for glasses, the technique monitors the phase lag of the strain behind the stress, as the sample is sinusoidally stressed while the temperature is increased at a predetermined rate. Although viscous flow is not observed in BCC iron cobalt, the change in Young’s modulus between the ordered and disordered state may be observable.

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* Contains information on alloys with non-equiatomic composition.

+ Recommended reading.

-229- Appendix 1: Bragg Williams Theory and the Long Range Order Parameter.

It is important to understand/^Bragg Williams theory of ordering because the majority of experimentalists compare their results to this theory. This summary of the Bragg Williams theory [1934] is limited to the body centred cubic alloys relevant to iron cobalt.

Definition of the long range order parameter, S.

Let us initially assume a perfect lattice of BCC iron with N atoms of iron. This constitutes two interpenetrating simple cubic lattices, a-sites and p-sites. If atoms of Fe are now replaced with atoms of Co, they can either be positioned on an a-site (defined as a position of order) or on a p-site (a position of disorder). Let there be n positions in a crystal block in which such replacements are to a-sites. In iron cobalt alloys n=N/2, because only cube centt& s are characteristic of the superlattice. Let r be the fraction of Co atoms on a-sites; r=%, for randomly arranged Fe50at%Co and r=0.33 for randomly arranged Fe33at%Co.

It must be remembered that we are dealing with simple lattices and that the a and p-sites are identical as regards the phase pattern, and that we can only define the a positions of order as a set distinguished by having a higher proportion of replacements than other similar sets of phase sites.

We define the degree of order, S as follows: Let p be the probability that an a-position is occupied by a replacement, then,

S = Actual value of p - Value of p for full disorder Value of p for full order - Value of p for full disorder

However, it is intended to compare the degree of order in the Fe50at%Co alloy to that in the Fe33at%Co alloy. The above equation will give a value of S=1 (i.e. "Actual value of p" = "Actual value of p for full order") for both compositions, so the values used in the denominator must be those used for Fe50at%Co. i.e.

-230- S = (Actual value of p - Value of p for full disorder) (Val. of p for full ord - Val. of p for full disord) in 50%Co

= (p - r) (1 - r)so*co

= 2 (p - r)

For Fe50at%Co : S - 2(p - 0.5) 0

For Fe33at%Co : S = 2(p - 0.33) 0 < S < 0.67 (Fe35wt%Co)

The neutron diffraction study of Lyashenko et al. (1962) reported values for S of 0.7 and 1 for Fe35wt%Co and Fe50wt%Co respectively.

-231- Acknowledgeaents.

I would like to thank Dr.R.Rawlings and my fellow PhD students for their assistance through useful debates throughout my PhD, as well as Prof. Pashley for the provision of laboratory facilities. In addition, I am grateful to Richard Sweeney, Ian Button and Graham Briers for technical support, and to R. Baxter, Keith Bor tin and the workshop staff, without whom my studies would not have been possible. Finally, I would like to thank Dr.R.Majorf Dr.V.Samadian and Mr.W.Gemel of Telcon Metals Ltd. for useful discussions and for the financial backing of this project.

-232-