70 - 19,345
MILLER, Robert James, 1944- STOICHIOMETRIC AND COMPOSITIONAL EFFECTS IN CRYOCHEMICALLY PROCESSED BARIUM FERRITE CERAMICS.
The Ohio State University, Ph.D., 1970 Engineering, chemical
University Microfilms, A XEROX Company, Ann Arbor, Michigan
THIS DISSERTATION HAS BEEN MICROFILMED EXACTLY AS RECEIVED STOICHIOMETRIC AND COMPOSITIONAL EFFECTS IN CRYOCHEMICALLY PROCESSED BARIUM FERRITE CERAMICS
DISSERTATION Presented in Partial Fulfillment of the Requirements for the Degree Doctor of Philosophy in the Graduate School of The Ohio State University
By Robert James Miller, B. Cer. E ., M.S.
* * * * * *
The Ohio State University 1970
Approved by
Adviser Department of Ceramic Engineering ACKNOWLEDGMENTS
The author wishes to express his sincere appreciation to his adviser, Dr. Ralston Russell, J r ., for his advice and guidance throughout the course of this investigation. His consultation and suggestions proved, a great asset in planning and carrying out this investigation. The advice of Dr. Morris Berg and Dr. Rodney Tettenhorst in vari ous aspects of this study proved very helpful. The technical assistance of Dr. Michael Komarmy, Robert Winberg, Sam Matarella, David Achey, Steven
Miele, Roger Brown, Carl Schaefer, Arthur Somers, and Robert Carol is gratefully recognized. They have helped to make this investigation more meaningful and complete. In addition, the author wishes to take this opportunity to express his gratitude to the AC Spark Plug Division of General Motors, who sponsored this investigation, and especially to Dr. Karl Schwartzwalder, whose efforts made the entire program possible. His guidance, advice, and even prodding were greatly appreciated. The author e g resses his gratitude to his wife, Marilyn, for her de votion and understanding during the last several months of this endeavor. Finally, the author wishes to somehow convey his utmost gratitude to
Mrs. Ruth Wick for all her help in preparing this manuscript. Without her assistance, it could never have been completed at this time.
li VITA
June 27, 1944 . . . Born — Uhrichsville, Ohio 1967 ...... B. of Cer. E ., The Ohio State University, Columbus, Ohio 1967 ...... M. S c., The Ohio State University, Columbus, Ohio 1967-1969 .... General Motors Fellow, Department of Ceramic Engineering, The Ohio State University, Columbus, Ohio 1967-1969 .... Research Engineer, General Motors Corporation, Flint, Michigan
PUBLICATION
"BaO Glasses in High Alumina Ceramics," Bull. Am. Cer. Soc., 1969.
ill TABLE OF CONTENTS
Page
ACKNOWLEDGMENTS...... il VITA ...... Hi LIST OF T A B L E S...... vl LIST OF ILLUSTRATIONS...... vil INTRODUCTION...... 1 LITERATURE REVIEW ...... 5 I. Material P reparation...... 5 n. Calcination ...... 17 III. Sintering...... • . 18 IV. Composition...... 28 MODE OF INVESTIGATION...... 34 EXPERIMENTAL PROCEDURE...... 36 I. Materials...... 36 n. Ferrite Preparation...... 39
in. Specimen Preparation...... 51 IV. Specimen Composition...... 52
V. Property Measurement...... 56 a. Magnetic P ro p erties...... 56
b. Fired Density...... 61
c. Microstructure...... 61
iv RESULTS AND DISCUSSION...... I. Material P reparation...... II. Compositional Evaluation ...... a. BaO • 6 FegOg Materials ....
b. BaO • 6Fe2Og + 2% Si02 Materials c. Varied BaO: Fe2Og Ratio Materials d. BaO -5 .5 Fe2Og Materials . . . e. BaO -5 .5 Fe2Og Additive Study .
f. Dual Additive Materials ....
III. Hot Pressing Evaluation...... SUMMARY...... CONCLUSIONS...... APPENDIX...... BIBLIOGRAPHY...... LIST OF TABLES
Table Page
1. Barium Acetate Lot A n a ly s is...... 36 2. Bismuth Trioxide Lot A n a ly s is ...... 37 3. Titanium Oxide Lot A n a ly s is...... 38 4. Calcium Oxide Analysis, Maximum Im p u r itie s...... 38 5. Lead Oxide Analysis, Maximum Im purities...... 39 6. Chemical S y s te m s...... 43 7. Batch C om p osition s ...... 45 8. Calcination Results ...... 48 9. B inders...... 48 10. Binder E valuation...... 50 11. Compositions for BaO • 6 FegOg with A d d itiv e s ...... 53 12. Compositions with BaO • 6 Fe O, + Si09 as Base M a teria l...... 7 ...... 54 13. Compositions for Evaluation of BaO: Fe2Og R a tio...... 55 14. Final Additive Evaluation, BaO • 5.5 Fe_0„ Base M a teria l...... 57
vl LIST OF ILLUSTRATIONS
Figure Page
1. Cold Bath for Freezing and Pelletizing ...... 12 2. Schematic Diagram of Freeze D ryer...... 14 3. Phase Diagram...... 15 4. Pelletizing Apparatus...... 40
5. Hysteresigraph Electromagnet...... 58 6. Hysteresis Loop Measurement T ech niqu e...... 60 7. Remanence Values for BaO • 6 FegOg + A d d itives-I...... 66 8. Remanence Values for BaO • 6 Fe^Og + Additives-II...... 67 9. Coercive Force Values for BaO ■ 6 FegOg + Additives-I . . . 68 10. Coercive Force Values for BaO • 6 FegOg + Additives-II . . . 69 11. Energy Product vs. Intrinsic Coercive Force; BaO • 6 Fe2Og + A d ditives-I ...... 70
12. Energy Product vs. Intrinsic Coercive Force; BaO • 6 FegOg + Additives-II ...... 71
13. Fracture Surfaces of BaO • 6 Fe2O g ...... 73 14. Fracture Surfaces of BaO • 6 Fe2Og + 2% Si02 ...... 75 • 15. Remanence Values for BaO * 6 Fe2Og + 2% Si02 + AlgOg . . . 76
16. Coercive Force Values for BaO • 6 Feo0 Q + 2% SiO« + 77 AV>3...... ? ...... i7 . Energy Product vs. Intrinsic Coercive Force; BaO • 6 Fe2Og + 2% SiOg + AlgOg ...... 78 18. Remanence Values for BaO • 6 Fe2Og + 2% SiOg + BlgOg . . . 79 viii
Figure Page 19. Coercive Force Values for BaO • 6 Fe00„ + 2% Si09 + Bi20 3 ...... -2 .3 ------? ...... 80 20. Energy Product vs. Intrinsic Coercive Force; BaO • 6 Fe2Og + 2% S102 + B igO g...... 81 21. Remanence Values for BaO • 6 FegOg + 2% SiOg + CaO .... 82 22. Coercive Force Values for BaO • 6 FegOg + 2% SiOg + CaO . . 83
23. Energy Product vs. Intrinsic Coercive Force; BaO • 6 Fe2Os + 2% SiOg + CaO ...... 84 24. Remanence Values for BaO* 6 FegOg + 2% SiOg + PbO .... 85 25. Coercive Force Values for BaO • 6 FegOg + 2% SiOg + PbO . . 86 26. Energy Product vs. Intrinsic Coercive Force; BaO • 6 FegOg + 2% SiOg + PbO...... 87 27. Remanence Values for BaO • 6 FegOg + 2% SiOg + TiOg . . . 88 28. Coercive Force Values for BaO • 6 Fe00„ + 2% Si09 + T i O g ...... 7 ...... 89 29. Energy Product vs. Intrinsic Coercive Force; BaO • 6 FegOg + 2% SiOg + TiOg ...... 90 30. Remanence Values for Various BaO: FegOg R a tio s...... 93 31. Coercive Force Values for Various BaO: FegOg Ratios . . . 94 32. Energy Product vs. Intrinsic Coercive Force; Various BaO: FegOg R a tio s...... 95 33. Remanence Values for Various BaO: Fe90„ Ratios with 2% S iO g ...... 96 34. Coercive Force Values for Various BaO: Fe90„ Ratios with 2% S iO g...... 97 35. Energy Product vs. Intrinsic Coercive Force; Various BaO: FegOg Ratios with 2% S iO g...... 98 36. Fracture Surfaces of BaO *6.5 F egO g...... 99 ix
Figure Page
37. Fracture Surfaces of BaO • 5.5 F egO g...... 100 38. Remanence Values for BaO • 5.5 FegOg + 2% Additive .... 102 39. Coercive Force Values for BaO* 5 .5 FegOg + 2% Additive . . 103 40. Energy Product vs. Intrinsic Coercive Force; BaO *5.5 FegOg + 2% A d d itive...... 104
41. Fracture Surfaces of BaO *5.5 FegOg + 2% T iO g...... 105 42. Fracture Surfaces of BaO *5.5 FegOg + 2% CaO...... 106 43. Fracture Surfaces of BaO • 5.5 FegOg + 2% AlgO g...... 108 44. Fracture Surfaces of BaO • 5.5 FegOg + 2% SiOg...... 109
45. Fracture Surfaces of BaO 5.5* FegOg + 2% BlgO g...... I l l 46. Fracture Surfaces of BaO • 5.5 FegOg + 2% PbO...... 112 47. Remanence Values for BaO • 5.5 FegOg + S iO g...... 114 48. Coercive Force Values for BaO • 5.5 FegOg + S iO g...... 115 49. Energy Product vs. Intrinsic Coercive Force; BaO • 5.5 FegOg + S i O g...... 116 50. Remanence Values for BaO • 5.5 FegOg + P b O ...... 117 51. Coercive Force Values for BaO * 5.5 FegOg + P b O ...... 118 52. Energy Product vs. Intrinsic Coercive Force; BaO • 5.5 FegOg + PbO...... 119 53. Remanence Values for BaO* 5 .5 FegOg+AlgOg ...... 120 54. Coercive Force Values for BaO* 5 .5 FegOg + AlgOg ..... 121 55. Energy Product vs. Intrinsic Coercive Force; BaO • 5.5 FegOg + AlgOg...... 122 56. Remanence Values for BaO • 5.5 FegOg + BlgOg...... 123 57. Coercive Force Values for BaO* 5.5FegOg + BlgOg ..... 124 Figure 58. Energy Product vs. Intrinsic Coercive Force; BaO • 5.5 FegOg + BigOg...... 59. Remanence Values for BaO* 5 .5 FegOg + Two Additives . . . 60. Coercive Force Values for BaO *5.5 Fe„0„ + Two A d d itiv e s ...... 61. Energy Product vs. Intrinsic Coercive Force; BaO *5.5 Fe2 ° 3 + Two Additives ...... 62. Remanence Values for Hot Pressed BaO *5.5 Fe_0„ + A d d itiv e s ...... 63. Coercive Force Values for Hot Pressed BaO • 5.5Fe_0„ + A d d itiv e s ...... 64. Energy Product vs. Intrinsic Coercive Force; BaO *5.5 FegOg + A dditives...... 65. Laboratory Hot P r e s s...... 66. Cross-Sectional View of Hot Press Assembly...... INTRODUCTION
Ceramic permanent magnets are recognized to possess certain advantages over metallic magnets, the most important being their high re sistance to demagnetization, an essential in small motor applications. Material costs, always an important consideration, are also lower than for metal magnets. Inherently high resistivity is not a prime consideration for these magnets, although this characteristic is quite important in the high frequency applications of such ferrites. Many techniques for preparing the magnetic ferrite material have been devised. The most popular industrial method is wet ball milling ferric oxide and an alkaline earth oxide or carbonate (1, 2). This has yielded a material satisfactory for some applications, especially those where size and/or weight limitations are not critical. However, the magnetic strength of the magnet becomes important in cases where size and weight of the motor assembly are significant. A more powerful magnet will accommodate a smaller coil assembly for the same output, thus providing a motor of much smaller mass. Various approaches have been studied in attempting to upgrade ce ramic magnets, including the preparation technique for making the ferrite powder, the sintering techniques, and the use of additives to improve the magnetic characteristics. Forming methods may also be varied to produce stronger magnets, since proper orientation of the magnetic particles yields higher property values in specific directions. Several powder preparation techniques have been used with varying success. Milling of iron oxide with either barium or strontium carbonate is a common method, with the carbonate decomposing to an oxide during calcining. Sulfates or nitrates may be used in a similar manner (1). Pro cesses which do not involve ball milling include a sol-gel process (3), vapor phase reaction (2), coprecipitation (2, 3, 4, 5), and freeze drying (6). These processes are somewhat more costly and thus are not widely used commercially. However, the freeze drying process is in its early de velopmental stages and has not been evaluated for the production of per manent magnets. Freeze drying was selected for this investigation because it elimi nated several problems associated with other processes. For example, ball milling of the raw materials is always a source of contamination for the material being milled. Milling in steel mills with steel balls intro duces excess iron into the ferrite powders, causing difficulties in control ling the B a: Fe ratio within the batches. Thus, the powder obtained after processing may be quite different from the starting components. The coprecipitation process involves some rather intricate prob lems, the foremost being realization of materials having the proper ratto of components. Concentrations, solubility products, and temperature must all be considered when using this technique. After the material is precipi tated, it must be filtered and washed to remove excess solution, which may contain unwanted cations or anions. Filtering this material often requires considerable time and is rather inefficient. Vapor phase reaction and the sol-gel process are both rather
expensive methods for producing materials. Vapor phase powders, being very light and fluffy, present difficult collection and handling problems. The sol-gel process is slow and poses problems similar to those encoun tered with the coprecipitation technique. Freeze drying, on the other hand, is a relatively simple process once suitable raw materials have been selected. The main features in clude forming a solution or slurry, freezing the solution, and drying the material in a vacuum chamber. This investigation involved a study of the feasibility of producing barium ferrite permanent magnets by freeze drying techniques, since freeze drying should produce a very pure, homogeneous ferrite over which good control could be exercised. Also, the optimum properties of the freeze dried material were ascertained by varying the BaO: Fe^Og ratio and through the use of additives in the batch composition. Two other methods used for optimizing properties were controlled sintering atmosphere and hot pressing, with hot pressing expected to yield a dense, fine-grained microstructure, conducive to optimum magnetic characteristics, x While the field of magnetic ceramics is relatively new, it is expand ing at an extremely rapid rate. And while growth has been rapid, there has been a paucity of fundamental research in the field. This investigation was designed to provide some insight into the critical parameters which affect magnetic characteristics of oxide materials and possibly improve their competitive position in the magnet industry. Present industrial research and development seems concentrated almost entirely on processing methods rather than on material considera tions. While every investigation necessarily involves some processing methodology, this investigation has dealt chiefly with the development of material properties. Thus, even though considerable effort was expended developing the freeze drying process, it is the material which is of prime consideration. LITERATURE REVIEW I. Material Preparation
The simplest method of processing is wet or dry ball milling the raw materials together, followed by a calcination step to develop the ferrite structure. Milling of oxides is fairly common in industry (1, 2, 7), since the raw materials are relatively cheap and readily available. Economos (1) indicates that wet milling may be preferred with oxide materials, because hydration of the oxides increases their reactivity. Harwood et al. (8) discussed reactivity of wet ball milled iron oxide but did not mention hydration of the material as a contributing factor. They
defined reactivity index as the ratio of ferric oxide powder reacted to the original amount present in the sample, as determined by quantitative x-ray analysis utilizing prepared standards of known compositions. It was re ported in the study that reactivity of the ferric oxide cannot be determined from surface area measurements alone. The milling operation was con sidered to include four distinct phases; namely, 1) deagglomeration of the ferric oxide particles, with an accompanying Increase in reactivity but not in surface area; 2) coating of the large grains with smaller particles; 3) the
sudden breakup of particles; and finally 4) reagglomeration, which decreases reactivity. Raw materials other than oxides may be used in ball milling, for ex ample, oxalates (2), carbonates (7), nitrates (1), or sulfates (1) in various combinations with or without oxides. Such materials are usually ball milled, dried, presintered (calcined), and ball milled again. The presintering step, usually carried out at 900° to 1000°C, converts the materials to oxides which then react to form the ferrite structure (1). Shrinkage during sintering of the formed magnet is decreased by presintering, and higher fired density may result (9). Brown (10) found that a very reactive material could be formed by milling salts (rather than oxides) and using low temperature calcination. However, low density is often a problem, which may best be alleviated by a second milling. Unfortunately, such treatment may impair the material's reactivity. Ferrite materials may be calcined a second time to develop optimum magnetic characteristics. Heimke (11) reported that during milling, the coercivity of barium ferrite tended to increase due to the more uniform fine grain size, but to decrease as a result of the creation of mechanical defects in the crystal structure. Annealing the material at 1000°C for one-half hour removed most of the defects, with a corresponding increase in coercivity. Precipitation techniques are often used in ferrite preparation, result ing in a material with high density and small grain size (3). The ferrite is generally more homogeneous than material that is mixed in a ball mill, es pecially if the material is coprecipitated rather than a single precipitation of one compound on another. A simple example of the single precipitation process is the precipi tation of barium carbonate on ferric oxide particles (12). Barium hydroxide is dissolved in water and insoluble ferric oxide added to the solution. Then, a carbonate (ammonium carbonate, for example) is added and barium carbon ate precipitates on the surfaces of the iron oxide particles. This material is
then dried and calcined to yield barium ferrite. Economos (7) reports a similar precipitation process for nickel fer rites, in which nickel carbonate is precipitated on ferric oxide. The oppo site of this is a French process (13) wherein ferric ions are dissolved in a solution containing ferrous sulfate and then deposited as ferric oxide on a finely divided divalent metal. While single precipitation processes forming one material on the sur face of another produce a fairly homogeneous material, a coprecipitation process in which the iron compound and the divalent metal compound are both formed from a solution yields even better homogeneity. Coprecipitation is a much more involved process than simple precipitation, in that two distinct materials are being formed simultaneously from a single liquid solution, and the concentration of each in the solution can affect the precipitation rate of the other. Hence., much closer control must be exercised over the process to insure formation of a good magnetic material. Some ferrites may be coprecipitated by starting with sulfates dis solved as an aqueous solution and adding sodium hydroxide to precipitate the materials as hydroxides (1). Micheli (5) reported a similar process in some detail. In this case, the process originated with a solution of soluble salts (for example, sulfates, chlorides, or nitrates) to which a solution of sodium or potassium hydroxide is added, precipitating iron and barium hydroxides. It was pointed out that the pH of the solution and the concentration of ionic species must be closely controlled for the two compounds to be precipitated in the correct ratio. Oxalates also may be produced by such coprecipitation reactions (1, 2, 7). For example, in forming nickel ferrites by this process, the follow ing reaction takes place:
2 Fe2+ + Ni2+ + 3 (C204)2“ + 6 H20 -> NiFe2(C20 4)3 • 6 H20
The ferrite structure is formed by subsequent heat treatment of the precipi tate. Formates and acetates may also be precipitated, the main concern being that the products easily decompose to form the ferrite and that the de composition by-products be gaseous, so that there is no residue to contami nate the material. Oxalate and hydroxide coprecipitations yield powders which possess higgler reactivity. Economos (7) found in the case of nickel ferrites that the oxalates had the higher reactivity, since mixing was on a molecular scale with NiFe2(C204)3 • 6 HgO crystallizing, rather than Fe(OH)2 and Ni(OH)2 in the case of hydroxides. From this example, it is seen that a material's re activity is dependent on the homogeneity of the species precipitating. A homogeneous, reactive material means low temperature ferrite formation with a minimum of sintering and grain growth. It has also been reported (2) that coprecipitated oxalates formed very small crystallites after calcination, indicating high reactivity. O'Bryan et al. (14) provide details for a coprecipitation process used to obtain good material for mlcrostructural control. An aqueous solution of each salt is prepared and added in the proper ratio to a stirred solution at
80°C to 100°C, carefully controlling the pH of the solution at all times. The solution is boiled for several hours, following which the precipitate Is fil tered, washed with ammonium hydroxide, and dried at 110°C. The water of 9 hydration is removed by prolonged heating (16 hours) at 500 °C. The authors suggested that the anions in the solution determine the porosity of fired speci mens and that the cations of the base used in precipitation determine the fine ness and uniformity of the grains in the microstructure. Further, the valence of the iron in the starting solution affects the physical characteristics of the powder, and hence the porosity and grain size. Thus, it quickly becomes evident that close control of the coprecipitation process is necessary to ob tain satisfactory and reproducible material. Other processes and materials for precipitation reactions may be found in the patent literature (4, 12, 15, 16, 17), while many other minor variations in the techniques exist. One recent development involves the use of higher pressures to trigger the precipitation reaction (18), and further work on this process to form ferrite materials is presently in progress. Among the variables which must be controlled are concentrations of all chemical species, solution pH, temperature, pressure, and time of reac tions, all of which influence the quality and amount of material produced. Coprecipitation is recognized as a good method for producing reac tive materials with fine particle size. The materials tend to be very homo geneous and contain few impurities when processed correctly. However, there are inherent drawbacks in the process which definitely detract from its initial appeal. It is obvious from the preceding examples that the process is rather complicated and requires stringent control for the successful pro duction of quality material. For instance, small amounts of impurities can have a great effect on the grain size of the powder produced (3). The physi cal form of the precipitated material, gelatinous insome cases (5), makes 10 filtering and washing difficult, resulting in tendencies for impurities to be present in the material. Decomposition of these impurities during sinter ing leads to lower density in the fired samples. Coprecipitation requires the precipitation (generally) of two differ ent crystalline phases at the same time and in the correct ratio for ferrite formation. To accomplish this task successfully, concentration (and solu bility product) of each species in the solution must be taken into considera tion. Since solubility product is temperature dependent and the solubility of each material is different and often dependent on the concentration of the other ionic species present, it is rather difficult to obtain the correct ratio of precipitating materials to form a good ferrite material. Vapor phase reaction is a more exotic method of forming ferrite materials. A typical reaction (1) involves dissolving nitrates of the appro priate metals in alcohol and igniting this solution in an oxygen atmosphere, where decomposition and reaction between materials take place. Another (19) calls for the reaction at 700°C to 1350°C of ferric oxide with barium or strontium halide in an atmosphere of water vapor and/or oxygen with other nonreducing gases. Such materials are light, fluffy, reactive, very fine, and often ex hibit a high magnetic moment (2, 3), although some calcining may be neces sary to increase their homogeneity. Since the materials are light and fluffy, collection and handling pose problems. However, the main drawback to this process is that relatively expensive chemicals are required to produce a good ferrite material. The sol-gel. process (3, 20) is another rather exotic method of forming ferrite materials. An aqueous colloidal suspension of oxides is prepared from nitrates or other salts and dispersed in small droplets at the top of a column of an organic liquid, which must be immiscible with water but yet acts as a dehydrating agent. The microspheres are then dehydrated, gela tion occurs within them, and they are collected at the bottom of the column. The material is separated from the organic liquid, dried, and fired. It is evident from this description that the process is somewhat complicated and as a result inherently slow. Also, the materials required are very expen sive and, by the nature of the process, must be water soluble. However, this method has not been thoroughly investigated and may yet prove to be feasible in certain applications. Freeze drying, the final technique to be discussed, is presently in the early stages of development. The basic process is an extremely sim ple one (3). The material to be freeze dried is frozen in some manner, the frozen material placed in a vacuum, and the water is sublimed from the sample, leaving a dry powder. Subsequent calcining may or may not be necessary, depending on the requirements of the material being pre pared. The technique utilized in this research originated at the Bell Tele phone Laboratories (6, 21). The raw materials are first dissolved in an aqueous solution, and the solution is then introduced into an Erlenmeyer flask. By increasing the air pressure in the space above the liquid, it is sprayed into a cold, swirling liquid and frozen, as shown in Figure 1. When the solution is sprayed into swirling liquid hexane cooled with an acetone-dry ice mixture, it breaks up into small spheres, which are 12
COMPRESSED AIR
STIRRER SOLUTION
HEXANE
ACETONE- DRY ICE BATH
FIBERFRAX INSULATION
Figure 1. Cold Bath for Freezing and Pelletizing 13
i frozen almost instantaneously. Thus, a very homogeneous material in the form of small spherical particles is obtained. The frozen spheres are removed from the liquid hexane by screen ing and placed in a freeze dryer, the essential elements of which are shown in Figure 2. The frozen material is placed in trays and set on shelves, which may either be heated or cooled. A condenser is present to collect the water vapor which forms in the system during sublimation. The entire process takes place in a vacuum chamber, low pressure being vital for success. Drying the frozen spheres without melting is accomplished by main taining a pressure in the system below the triple point of the frozen solution, as shown in Figure 3. The dashed lines denote phase areas for the pure water system, with the triple point at T. Addition of a soluble salt lowers the triple point, in this case to point Q, with the actual position of Q depend- » ing on the concentration of salt in the solution. By keeping the system's pressure below point Q, 1.63mmHg in this case, the ice can be evaporated from the sample without the formation of liquid water. The condenser in the system allows much more rapid evaporation from the sample than would otherwise be possible, whereas the ability to control shelf temperature in the freeze dryer gives positive control over rate of evaporation, which is accomplished as quickly as possible while still keeping the system's total pressure below the triple point. When dried, the powder is removed from the freeze dryer and may be calcined or sin tered, converting the materials from salts to oxides.
When using this method, quick freezing is necessary to preclude PRODUCT SHELF
> REFRIGERANT OR CONDENSER <2 1 HEAT
< J = 3 REFRIGERANT
VACUUM < J
Figure 2. Schematic Diagram of Freeze Dryer PRESSURE (mm Hg) 1.63- 4.58 SALT ICE e eat re tu pera TEMPERATUREtem (°C) -12 Figure
3. Phase Diagram Phase 3. 5 ■ SOLUTION SALTVAPOR+ / I c) 'c
16 segregation of the chemicals during the freezing operation (22). The result is a very fine grained, homogeneous material of high purity. The occur rence of melting during the drying process would result in larger particles and segregation of material, and thus must be avoided. Sublimation often is carried out at system pressures of 50-100 microns Hg to prevent such cir cumstances (22). Final particle size will depend not only on solution concen tration, but also the time and temperature of calcination (21). Another innovation possible with the process is the utilization of a fluidized bed in the drying operation (23, 24). With a fluidized bed, the re moval of water vapor from the ice-vapor interface is more rapid, thus lower ing drying time. Heat transfer can also be made much more efficient in the same manner, again decreasing drying time, so long as the pressure within the system is kept below the triple point pressure of the solution to avoid melting (24). A secondary bed of material, such as salt or sugar, might be used for this purpose, resulting in drying times in the 6-8 hour range. The freeze drying method of material preparation presents the advan tages of good reproducibility, little contamination, and low reaction and sin tering temperatures (3). The low reaction temperatures result from the fine particle sizes attained, often 100-5000 angstroms after calcining. In addi tion, the distribution of particle sizes obtained is rather narrow. Nucle- atlon, crystallization, and chemical changes at low temperatures should be considered when using this process, as well as the mutual solubility of the raw materials used. The principal drawback to the method at present is the lack of proper equipment, a problem which should soon be overcome con sidering the expanding interest in freeze drying. 17
II. Calcination
While calcination of ferrite materials was discussed briefly in con nection with ball milling, some further remarks are in order. Calcination before subsequent forming and sintering processes can vastly alter the properties of the freeze dried material, and it is often used to minimize ir regularities resulting from earlier processing. It is also useful in allevi ating the effects of raw material variation and its homogenization if more than one ingredient is involved (25). Decomposition reactions which take place during calcination convert the freeze dried material to the oxide form usually desired, and thus reduce shrinkage during subsequent firing. Calcination temperature may be selected by examination of x-ray patterns for the raw materials and the ferrite itself, choosing a major peak for each material. An examination of the ratio of these peak heights in any sample allows an estimate of the amount of each phase present. Based on these results, the preferred calcination temperature may be selected (26). An easier, faster method of selecting calcination temperature is based on differential thermal analysis (DTA). A sample of the batch is heated to produce a DTA pattern, which can be analyzed to determine the reactions taking place and the temperature range over which they occur. With this knowledge, a proper calcining treatment may readily be selected. In most cases it is rather difficult to evaluate the effect of calcin ing treatment. Ferrite content of the calcined material is one measure of the extent of reaction, as is the true density of the material. Surface area is also an indicator often used, with a given surface area producing a 18 corresponding set of magnetic properties for a given powder (25). The important consideration in calcining ferrite materials is whether the calcining temperature is above or below the recrystallization tempera ture (i.e ., the temperature at which the oxides are converted to the ferrite structure). Powders calcined above the recrystallization temperature are characterized by discrete crystallites, higher sintering temperatures, and high tap and mold densities (27, 28, 29). It has been reported (27) that such powders exhibit reduced grain growth activity, except in cases where the calcined powder was very finely milled and particle sizes were sharply dis tributed. Ferrites calcined below the recrystallization temperature showed lower tap and mold densities due to a narrower size distribution and many pores (29). However, lower sintering temperatures are possible with the higher reactivity material calcined at lower temperatures (28), the calcined material being in the form of polycrystalline aggregates (27). Hence, cal cination is a critical step in ferrite manufacture, having a definite effect on subsequent material properties.
HI. Sintering
Since sintering is often the final step in ceramic processing, it is ex tremely important with regard to attainment of the desired product character istics. Even if all processing prior to sintering has been carried out advan tageously, a good product will not result without proper sintering or final firing. Ambient furnace atmosphere is increasingly recognized as a most im portant consideration in sintering ceramic bodies. However, the optimum 19 atmosphere for ferrite sintering is debatable. Most authorities specify a neutral or oxidizing atmosphere, but Carter (30) advises initial sintering in a reducing atmosphere to form a nonmagnetic second phase in the body. The latter is based on the premise that the second phase is located in the grain boundaries where it will inhibit grain growth. Subsequent sintering in an oxidizing atmosphere converts the nonmagnetic phase to a second ferromag netic phase, thus increasing the magnetic strength while still inhibiting grain growth. Another patent (31) also mentions the need for a reducing agent in the body (suggesting flowers of sulfur or powdered charcoal), but this applies to the presintering stage of production. Van Hook (32), on the other hand, feels that having a highly oxidiz ing atmosphere is important. In cooling a BaO* 6FegOg melt in air, an other type of ferrite, BaO • 2 FeO • 8 FegOg, was noted, and he concluded that higher oxygen pressures were necessary to maintain the hexagonal barium ferrites, BaO • 6Fe2Og. Experimenting in atmospheres with oxygen partial pressures from 0. 01 atmosphere to 1.0 atmosphere, he found BaO • 6 Fe^Og to be stoichiometric up to 1450°C at one atmosphere of oxygen pressure. Inert atmospheres are specified in some patents, but small amounts of oxygen are often Included. Pierrot and Lescroel (33) advise sintering In a nitrogen atmosphere containing 0.01 to 1.2 percent oxygen, but cooling in a 100 percent nitrogen atmosphere. Argon or carbon dioxide are other inert atmospheres mentioned for sintering ferrites (34), with very close tolerances specified for the level of oxygen impurity present. Oxygen content in the furnace atmosphere has been related to density and thus to magnetic properties (35). Spinel ferrites fired at 1200°C were 20 not affected by oxygen content, but at 1400°C fired density was increased 6 percent by decreasing the atmospheric oxygen content from 21 percent to 10 percent. This evidence supports the view that a slightly oxidizing or in ert atmosphere is preferable for ferrite sintering. Water content of the furnace atmosphere has an affect on the sinter ing behavior of many ceramics. In work with barium ferrites, Heimke (36) found that moisture in the ambient atmosphere increased coercive force, remanence, and saturation magnetization at sintering temperatures below 1200°C. Above this temperature the effect of water vapor is greatly dimin ished. Strontium ferrites show a similar diminution effect above 900°C, but mixed barium-strontium ferrites showed improvement in properties at all sintering temperatures. Heimke felt that in some manner the water cata lyzed the healing of defects in the crystal structure, causing an improve ment in magnetic properties. The goal of sintering ferrite materials is the development of a strong, dense ceramic body which possesses good magnetic properties. In sintering ferrites, densification accompanied by grain growth occurs, followed by sec ondary recrystallization (9). This secondary recrystallization, the develop ment of some very large grains, is especially detrimental to magnetic properties. The ideal microstructure for hard ferrites is fhie grains and no porosity, with a resulting high density (37). Coarse starting material is suggested as a means of eliminating secondary recrystallization and reduc ing grain growth (9). This view is supported by Kingery (38), who feels there are a few larger than average particles in a fine-grained powder.
These particles are the source of large grains produced by secondary 21 recrystallization. If the matrix material is coarser, the chances of a few larger grains being present is greatly reduced, thus increasing the diffi culty of nucleation for secondary recrystallization. Kingery further states that secondary grain growth is proportional to 1/Dm, where Dm is the aver age grain diameter of the matrix material. Thus, coarser starting material indeed reduces the rate of grain growth. Impurities or small amounts of porosity also inhibit grain growth, since they are barriers to grain boundary movement (9, 38). Minor addi tions to improve magnetic characteristics will be considered at length in a later section. The reason for the detrimental effect of large grain size on magnetic properties of hard ferrites is explained by their magnetic structure (39). The total magnetism of a body is due to the additive effects of small mag netic domains present in the structure. The domain is an area in which all the magnetic dipole moments are aligned; however, adjacent domains are generally not aligned with each other. Separating the two domains is a do main wall, which has higher energy than the domains. Such a situation is analogous to a polycrystalline ceramic microstructure, where two grains not crystallographically aligned are separated by a higher energy grain boundary. The hysteresis loop and the magnetic properties of a material de rived from the loop are dependent upon the action of domains and domain walls under the Influence of an external magnetic field. The internal fields of the material tend to be aligned with the external field, either by rotation of the moments of the domains or by movement of the domain walls (40). 22
Wall movement is such that favorably aligned domains increase in size at the expense of less favorably aligned domains. Schieber (41) reports that the sintering of ferrites above 1250°C brings about the growth of large grains and a resulting low intrinsic coercive force (Hcl). The lowered coercive force is apparently brought about by the interaction between the grain structure and the magnetic domain structure just discussed. If the grain size is below a critical value (42), each grain has only a single domain and no domain walls are formed. As grain size in creases, it becomes energetically favorable for grains to be separated into more than one magnetic domain by domain walls. Magnetostatic energy of the domain varies with the cube of the diameter while domain wall energy varies only with the square of particle diameter (40), so increased grain diameter leads to the formation of multidomain particles. The critical grain size for single domain particles is (43)
1_ d = (2 KT k/aM 4) 2 c s
where Ms is saturation magnetization, k is Boltzmann's constant, T c is Curie temperature for the material, a is the distance between magnetic ions in the lattice, and K is an anisotropy constant. Above this critical diameter, particles develop more than one domain. For barium ferrite, this critical diameter is approximately one micron (40), while for strontium ferrite per manent magnets, it is 10 percent larger (44). Intrinsic coercive force is a measure of the resistance of the material • to demagnetization. Demagnetization is brought about by rotation of domain orientation and movement of domain walls through the crystal. Since domain 23 wall movement requires less energy than domain rotation, materials arc more easily demagnetized when domain walls are present; hence, lower coercive force results (42, 43). While a fine grain structure is favorable, high density is also a de sirable characteristic of ferrite magnets, since it leads to high remanence,
B„,r and energy product, (BxH) max (44). Stuijts (40) confirms that Br is proportional to apparent density. Hence, the ideal structure is one of small grain size with low porosity. Several heat treatments are suggested to produce a favorable micro- structure. Sintering temperatures should be kept low, perhaps by use of fluxes, in order to keep grain size below the critical diameter (41). Economos (45) reports that high firing temperatures, long hold times, and fast cooling rates are particularly detrimental to both Br and Hc-. Stuijts (46) confirms that fast firing rates with a short holding time at peak tem perature give smallest grain size and optimum properties. Porosity, which has been blamed for a decrease of Br values, may also be at least partly responsible for discontinuous grain growth in ferrites (47). High purity, very homogeneous material helps to prevent such porosity because there are many nucleation sites; hence, a fine grained material develops. Another method for obtaining a fine grained structure with low po rosity is to employ pressure sintering or hot pressing techniques. Sintering under pressure allows better microstructural control since lower sintering temperatures and shorter holding times are required, limiting diffusion and hence grain growth and secondary recrystallization (48, 49, 50). 24
Most authorities believe that pressure sintering is a diffusion con trolled process, although some believe that plastic flow is involved (51). The former describe hot pressing as interparticle bonding due to atom or vacancy diffusion, yielding high density with no dimensional change, no recrystalliza tion, and little grain growth. Since the sintering mechanism is diffusional, it is less sensitive to particle size, distribution and shape than a mechanism in which extensive nucleation and grain growth occur (50). However, the temperature and pressure of hot pressing affect densification rate (positive correlation) as does particle size (negative correlation). Those championing plastic flow as the dominant process describe hot pressing as occurring in four stages (52); namely, A) Low temperature; a little punch rod movement resulting only in better packing of particles. B) Rapid punch rod movement; at the yield point where plastic flow « begins at points of contact between particles. C) Plastic flow, the rate depending on rate of temperature rise. D) Rate of densification decreases, approaching maximum density. Remaining porosity may be removed by diffusional mechanisms. Thus, it is apparent that there are differences of opinion as to the actual mechanics of hot pressing. Two general methods of heating are associated with hot pressing. In duction heating, generally with graphite dies, is a widely used technique (50,
53, 54). This arrangement allows fast heating rates, resulting in short sin tering cycles with little time for grain growth to occur. However, the use of graphite dies poses a limitation on the maximum pressure which may be 25
b applied. Such dies must also be lined with platinum foil or other material to prevent diffusion of carbon into the ferrite structure. Finally, graphite dies must always be used in a reducing or inert atmosphere (vacuum (53) and argon (52, 54) mentioned), whereas an atmosphere containing some oxygen may be desirable for sintering ferrites. Resistance heating elements also may be used for heating hot press dies and materials (49). In this case metal (50, 55), ceramic (56), or graphite (52, 55) dies may be used. Graphite dies may be used at higher temperatures, while selected metal alloy dies can withstand higher pres sures and may be used in oxidizing atmospheres. Ceramic dies are a good choice because they are rather inert, are not greatly affected by furnace at mosphere, and maintain their strength at high temperatures. Alumina and zirconia are promising die materials. While the firing cycle used in hot pressing naturally varies with the » materials being sintered, some generalities may be noted. The complete firing cycle is rather short, especially when Induction heating is used (57). At one extreme is a cycle in which the material is heated to peak tempera ture and cooled immediately, with no hold at peak temperature (52). Other cycles involve holding times generally between 10 and 45 minutes (48, 53, 58). At the opposite extreme is a cycle in which the material is held at peak temperature until die plunger movement stops, indicating maximum density had been reached (54). Choice of pressure and the period of its application vary widely. Ex perimental pressures have ranged from 3000 psi (52) up to approximately 100,000 psi (48), at which the author claims greatly enhanced consolidation 26 compared to ordinary hot pressing and at temperatures several hundred degrees below the temperatures required in conventional hot pressing. These pressures may be applied throughout the firing cycle (52, 54) or only during certain portions of the cycle (53, 58). Several variations or modifications have been made in hot pressing techniques. A simple modification is merely squeezing the material between two punch rods rather than having it constrained in a die (51, 58). This hot forging is useful only when no shape is required, such as calcining a ma terial under pressure (also called "calc inter ing" (51)). Pressure sintering a material at the temperature of a phase change or decomposition reaction often permits the sintering to be accomplished at lower temperatures (51, 55), and Hedvall (59) reported enhanced densifica tion and interparticle bonding take place during such reactions. Another author (55) reports that the enhanced reactivity is probably due to broken » bonds and unsatisfied valence charges that exist on the surface and in the bulk of the material during these transformations. Instability of atomic po sitions may also produce a transient superplastic state. At any rate, pres sure sintering the ferrite at the recrystallization temperature might well lead to enhanced densification. A method of hot pressing isostatically has been developed (50) in which heat and pressure are applied to a cold formed body in an autoclave. Pressure is applied with a compressed gas, while the autoclave is heated by a molybdenum or silicon carbide element. Isostatic hot pressing facili tates the handling of much more complex shapes than conventional pressure sintering, while at the same time eliminating contamination of the material. The principal drawback to hot pressing as a production technique is its inherently slow rate. The time required to heat the die and material to the sintering temperature, hold, and then recool to room temperature, as well as the heat lost in such an operation make it a feasible technique only for expensive, low volume parts. Two methods have been suggested to al leviate these problems. The first (60) operates on a cyclic basis wherein material is fed into the die at the top, pressure sintered, the upper punch rod raised, more powder added, pressure sintered, and so on. Here the column of material itself acts as the lower punch rod, the column moving slowly downward during the sintering so that all material is sintered in the hot zone of the furnace. However, only rod-shaped parts or parts which may be cut from rods can be fabricated by this method. Haertling (61) has developed a nsemi-continuous hot press tunnel kiln, " in which the die is never removed from the kiln and, hence, is not cooled to room temperature after sintering each part. Pressure is applied only while the die is in the hot zone of the furnace; after the die leaves the hot zone, the sintered part is ejected from the die and cooled. The die is then returned for recycling through the furnace. While hot pressing tends to yield a higher quality product than can be produced by conventional sintering, the cost is obviously also much greater. These two factors must be weighed against each other in con sideration of pressure sintering techniques. Equipment innovations in this area are sorely needed. 28
IV. Composition
Much has been written concerning the relationship of ferrite com position to the behavior of the sintered magnet. Barium ferrite as a pure compound has a stoichiometry of 6 Fe2Og for every BaO present. Of course these are merely oxide ratios and do not imply that the FegOg and BaO are separated into different structures. Variation of this BaO: FegOg ratio from the stoichiometric 1:6 can have a pronounced effect on the physical and magnetic characteristics of the body. Stuijts (46) reports difficulties in attempting to sinter pure, stoichio metric ferrites to high densities, due to the lack of defects (vacancies) in the lattice. Vacancies produce high diffusion rates and better sintering, while their scarcity causes the diffusion rate to be rather slow. Stuijts recom- 3+ 2- mends the addition of a small excess of BaO, which produces Fe and O vacancies in the structure. Barium ferrites with BaO: Fe2Og ratios of 1:5.6 to 1:5.9 had good sintering characteristics and yielded high density samples, with 1:5.9 proving the best combination. However, the presence of vacancies led to greater shrinkages than observed for the stoichiometric material. The patent literature contains claims that good magnetic characteris tics are obtained with ratios of 1:5.9 to 1:6.4 (62), "about 1:6" (63), and 1.0:6 to 1.2:6 (64). Thus, the range between 1:5 and 1:6.4 is well covered, with best characteristics reported at 1:5.9 or 1:6. A rather complete study was undertaken by Okamura et al. (65) of the effect of BaO: FegOg ratio on pure ferrites and also on compositions to 29 which BigOg was added. In studying ratios from 1 :4.4 to 1.: 6.6, a sharp drop in intrinsic coercive force (H .) was noted for materials above approxi- Cl mately 1:6 ratio, but only in samples containing BigOg. It seems apparent that the BaO: FegOg ratio can affect the behavior of some additives and in turn affects the intrinsic properties of the ferrite itself. Many additive materials are reported to produce beneficial effects on the magnetic properties and physical characteristics of ferrite materials. Bismuth trioxide, BigOg, is one mentioned frequently (41, 63, 65, 66, 67, 68, 69), and it generally lowers the sintering temperature of the ferrite and improves its coercive force (41, 66). The BigOg melts at 820°C and aids densification through the introduction of a reactive liquid phase in the system. Subsequently, it reacts with FegOg to form pseudo-perovskite, BiFeOg, as a nonmagnetic second phase in the bodies (41, 66). Such behavior could pro duce a higher density ferrite with smaller grain size and hence higher coer cive force. In another study, BigOg raised the intrinsic coercive force (by 30 to 50%) of materials for which the BaO: FegOg ratio was less than 1:6 (65). The sharp decline in H,. for ratios greater than 1:6 in specimens containing BigOg was not observed in undoped compositions of the same barium ferrite. Strontium ferrites show an opposite effect, with Hc^ values near the same level for all SrO: FegOg ratios in samples containing BigOg (67). For un doped strontium ferrites, Hci shows a sharp increase In samples where the ratio is greater than 1:6. Lead ferrites, on the other hand, exhibit be havior very similar to the barium ferrites, only less pronounced. No satis factory explanation is given for these observed behaviorial differences, but 30 it is intimated to be related to the lower firing temperatures of strontium and lead ferrites.
The amount of BigOg to be added to barium ferrite for best magnetic characteristics is subject to debate. Various authors give amounts of 4% to 20% (63), 1. 5% to 2.5% (65), and 0. 01% to 1. 0% (68). Bradley (66) claims the optimum content depends upon grain size, grain size distribution, par ticle packing, and effectiveness of dispersion. Such reasoning probably ex plains the widely different ranges reported by various investigators. Silica additions to ferrite bodies also seem to result in a liquid phase during sintering (37, 70), reportedly a strontium-iron silicate in the case of strontium ferrites. This liquid phase penetrates the grains and forms an impurity layer, thereby inhibiting grain growth during sintering. Hence, a small amount of silica concentrated in the grain boundaries can function ef fectively as a grain growth inhibitor.
Since SiOg can have a marked effect on the microstructure and proper ties of the ferrite, small percentages (about 2% in most cases) are usually recommended (36, 71, 72, 73, 74). SiOg is also often mentioned for use in conjunction with one or more other additives, such as CaO (70, 75) and AlgOg (71). In such situations certain proportions of additives give better results than others, but no explanation of the phenomena is given. Calcium oxide or some compound which will decompose to CaO is fre quently mentioned as a constituent in ferrite batch compositions. In some cases it is added in small percentages, apparently to form a second phase in the body and act as a grain growth Inhibitor (67, 68, 70-72, 75-77). A liquid phase forms during sintering, acts as an impurity layer, and may partially 31 form a solid solution with the magnetic phase (68). The resultant micro structure consists of small, uniformly-sized grains, a situation conducive to high intrinsic coercive forces. The combined effect of additives is again mentioned frequently, most often as a combination of SiOg and CaO (70, 71, 75). CaO and TiOg (71) and CaO and BigOg (67) are also noted as possible combinations. In the latter case, CaO additions to BigOg doped samples raised the value of HQi; but CaO additions to samples not doped with BigOg had no effect on Hc^. Some researchers claim CaO may enter the ferrite lattice itself (78- 81), apparently replacing divalent cations in the structure. Formulas sug- 2+ 2+ gested include (MO). „ (CaO)„ • kFeo0„, where M is Ba or Sr , x =0.3, 1 “ X X a J and k = 4.5 to 6.2 (79), and MxCai_xFei2°19' w^ere x = 0.6 to 1.0 and M is Ba2+, Sr2+, and/or Pb2+ (80). This is not unreasonable, since Ca2+ is somewhat similar to Ba2 + or Sr2 + but smaller. However, it is obvious that the actual manner in which CaO influences ferrite materials has not been determined conclusively. Lead oxide may be used in varying amounts to improve magnetic characteristics of barium or strontium ferrites (41, 44, 72 , 74 , 79-83). Since it can form a permanent magnet material with FegOg, it is rather safe to assume that the PbO becomes a part of the ferrite lattice replacing Ba2 + , rather than forming a second phase in the body. The amount added may vary 2+ 2+ anywhere from 1.5% (83) to the complete replacement of all Ba by Pb (pure lead ferrite). A complete series of solid solutions apparently exists between lead ferrite and barium ferrite or strontium ferrite, as might be expected since the difference in ionic radii is less than 15% in both cases. 32
PbO fluxes the barium or strontium ferrite, lowering the sintering temperature and thus producing a fine grained structure (41, 44, 72) and an
increased . Such action generally decreases porosity, with a resultant increase in remanence and energy product (44, 83). Even an improvement in mechanical properties was noted in one case (72). The main drawbacks to the use of PbO are its volatility and toxicity (41), which together pose practical sintering problems. As a result, the use of PbO in ferrite production is usually avoided when a more easily
handled and suitable additive is available. Alumina is added to ferrites, usually in conjunction with SiOg as a grain growth inhibitor (71, 73, 78). Such additions probably lead to the development of small amounts of mullite or glass in the bodies, slowing grain growth. Mones and Banks (84) report to the contrary that the A1 3+ ion preferentially enters the ferrite structure in octahedral sites, thereby decreasing the magnetization of the material. A similar result is reported
for titania additions to barium ferrites (85). Gorter (81) supports this view in his patent, but no detailed explanation is presented. Many other additives are suggested, generally in the patent literature, which might improve the properties of magnetic ceramics (14, 68, 69, 71-73, 77, 84, 85). These Include LagOg, BgOg, ASgOg, SbgOg, CrgOg, MgO, NagO, SnOg, ASgOg, BeO, CdO, CeOg, CSgO, CoO, COgOg, CuO, LigO, MnO, NiO, NbgOg, and others. It is evident that almost any additive which might be named is at least mentioned in one of the various ferrite composi-
• tional patents. The anions present in raw materials may also have an effect on sintering behavior and subsequent magnetic properties. Sulfates reportedly improve reactivity during presintering, decrease porosity, and raise Hcj (44). Another patent (73) also specifies small amounts of sulfates in the batch compositions to increase magnetism and magnetic anisotropy. The
2 - - replacement of some O ions with F reportedly results in a better mag net material (86); another patent specifies small amounts of Cl- (62). Nitrates are mentioned for making a more anisotropic powder (28), and barium permanganate (BaMnO^) may be beneficial (87, 88). Development of small amounts of second phases to inhibit grain growth and pin magnetic domain walls is a well established method for increasing H^. In almost all cases previously discussed, the second phase was nonmagnetic. Methods have been suggested for introducing a secondary magnetic phase and thereby upgrade magnetic properties even more (79, 89). Regardless of the composition of the second phase, it is obvious that the use of additives can drastically affect the ferrite’s micro structure and consequently its magnetic characteristics. MODE OF INVESTIGATION
The scope of this investigation includes complete formation and pro
cessing of a hard ferrite material, including measurement of the properties
of the resulting magnets. Freeze drying was selected as the method of material preparation be cause it presented a relatively simple way of producing a homogeneous,
carefully controlled material with little chance of external contamination. The selection of favorable materials and equipment for the process led to a much better understanding of the critical parameters involved. Calcination and other powder preparation techniques were considered in some detail. A calcining temperature was selected which provided a fully converted ferrite, but one which retained some degree of reactivity to en-
% hance the sintering process. Prepressing and granulating steps were effected to increase the bulk density of the calcined powder so that a realistic green density value could be attained. Powders were prepared in which the BaO: FegOg ratio was varied from 1:5 to 1: 7, thus fully bracketing the stoichiometric 1:6 ratio. A series of compositions were prepared which included an undoped material
and samples containing various percentages of additive materials. These were gradient fired, and the best additives and optimum percentages were
determined. Next, batches were formulated for the various BaO: FegOg ratios
34 35 including an undoped composition and one containing the optimum amount of the best additive for each ratio. In this manner the optimum BaO: FegOg ratio for the freeze dried barium ferrite material was selected. Using powder of the optimum ratio, another series of compositions containing various additives in varying amounts was evaluated magnetically. Additives which did not improve the magnetic characteristics of the material were eliminated from further study; optimum additive levels were ascer tained for those dopants which improved magnetic characteristics. The best compositions were also hot pressed in an attempt to develop an optimum microstructure and corresponding magnetic properties. A resistance heated laboratory hot press was constructed for this purpose, allowing for the introduction of any desired atmosphere during pressing. Comparative results were obtained for an inert atmosphere and an air at mosphere to determine the effect of 20% oxygen on the ferrite compositions. All fired samples were ground and measured with a hysteresigraph to determine coercive force, remanence, and energy product. Fired density measurements and photomicrographs were used to further characterize the various specimens. EXPERIMENTAL PROCEDURE I. Materials
Several different iron and barium compounds were investigated in an attempt to obtain materials suitable for freeze drying. Those which failed to satisfy various requirements of the process will only be noted later when the freeze drying process is discussed in detail. The components finally selected for freeze drying were ferric oxalate,
FegfCgO^lg • 6 HgO, and barium acetate, Ba(C2H30 2)2- The ferric oxalate (stock No. 31116) was obtained from the Ventron Division of Alfa Inorganics, Beverly, Massachusettes. The barium acetate was Fisher Certified ACS Reagent Grade B-24 powder from the Fisher Scientific Company. The lot analysis for the barium acetate is given in Table 1. Table 1. Barium Acetate Lot Analysis % Insoluble Matter 0.010
Cl 0.0002 Oxidizing Substances (as NOg) - 0.005 Substances not Ppt'd. by HgSO^ 0.040 Heavy Metals (as Pb) 0.0001 Fe 0.0002
Ca 0.002
Sr 0.2
36 37
The additives used in this investigation were AlgOg, BigOg, CaO, PbO, SiOg, and TiOg. The AlgOg and SiOg were both obtained from the Oxides Division of the Cabot Corporation, Boston, Massachusettes. The AlgOg was designated Alon G, a very fine, fluffy material produced by high temperature flame hydrolysis, with a purity of 99%+. The surface area oi' this and all other additive materials was measured using a Model 2200 Sur face Area Analyzer, manufactured by Micromeritics Instrument Corporation, Norcross, Georgia. The surface area of the AlgOg was 112.5 m /gm. The SiOg was Cabot's EH-5 Cab-O-Sil, a pyrogenically formed ma terial with a purity of 99.7 to 99. 9% and a maximum of 0.05% residual chloride. Its surface area was determined to be 303.7 m2 /gm. The BigOg and TiOg were obtained from the J. T. Baker Chemical Company, Phillipsburg, New Jersey. Lot analyses of the two materials are given in Tables 2 and 3, respectively. Surface area of the BigOg was measured to be 0.2 m2 /gm while the TiOg was 7.6 m 2/gm. Table 2. Bismuth Trioxide Lot Analysis % 99.6
Cl 0.0005
NOg 0.010 so4 0.001 Pb - 0.003
Fe 0.001
Alkalies and Earths (as SO^) 0.060 38
Table 3. Titanium Oxide Lot Analysis
% HgO Soluble S a l t s ------0.050 A s ------0.00002 Fe ------0.005
P b ------0.002 Zn ------______o .003
The CaO and PbO from the Mallinckrodt Chemical Works are charac terized by listing the maximum amounts of impurities, rather than lot analy ses. These maximum impurity levels are given in Tables 4 and 5 for CaO and PbO, respectively. Measured surface areas were 1.3 m /gm for CaO o and 0.5 m /gm for PbO. Table 4. Calcium Oxide Analysis, Maximum Impurities % C l---...... - ...... 0.005 Heavy Metals (as Pb) ------0.010 Insoluble in HC1 and NH4OH P p t . 1.500 F e ------0.100 Ignition Loss ------5.000 Mg and Alkali Salts (as MgO) — ------1.000
N03 ------° . 050 Sulfur Compounds (as SO^) ------0.100 39
Table 5. Lead Oxide Analysis, Maximum Impurities % C l ...... 0.005
C u ...... - - 0.002
Insoluble in C H g C O O H ------0.150 F e - ...... 0.005 Ignition Loss ------0.500
NOg - ...... 0.010 A g------0.00015
Not Ppt'd. by H gS------0.200
II. Ferrite Preparation Selection of materials for preparing the freeze dried powder served to point out the critical parameters of the freeze drying process itself. At first glance, the process appears very elementary, giving the Illusion that almost any chemicals can be used without difficulty. The selected chemicals are first dissolved to form a water solution, thereby allowing mixing to take place on a molecular scale. This solution is placed in an Erlenmeyer flask (Figure 4), and the pressure of the1 air above the liquid is increased. This forces the solution through the glass outlet tube and attached hose and through a nozzle at the end of the hose as a fine stream. This fine stream is directed into a beaker of pentane, which is cooled by an acetone-dry ice bath. A stirrer in the system createsa vortex in the cold pentane. As the liquid stream is injected into the vortex,
t it breaks up into small droplets striking the swirling pentane. The droplets freeze almost instantaneously, forming very small frozen spheres. For 40
COMPRESSED AIR
SOLUTION STIRRER
HEXANE
ACETONE- DRY ICE BATH
FIBERFRAX INSULATION
Figure 4. Palletizing Apparatus 41 optimum, freezing, the ratio of pentane to injected material should be at least 3*1. After the entire batch is sprayed into the pentane bath and frozen, the spheres are separated from the pentane by screening and placed in 3 stainless steel trays. These trays are set on the shelves of the freeze dryer (Virtis No. 10-010, Repp Division, The Virtis Company, Inc., Gardner, New York), and a vacuum is drawn on the entire system (Hyvac 14, Central Scientific Company, Division of Cenco Instrmnents Corpora tion). While the system is under vacuum, a condenser coil in the freeze dryer is maintained at approximately -50°C to condense the water vapor as it sublimes from the frozen sample. About 6 hours after the sample is first placed in the freeze dryer, the electric power is turned on to heat the shelves. This is actually done only after the pressure in the system has stabilized at about 50 to 100 microns, as measured with a McLeod pressure gauge. This insures that the water vapor will not evaporate from the sample at a disastrously high rate, which otherwise would raise the total pressure of the system and cause melting of the sample. The shelf temperature is raised to 90°F by the heaters, at which temperature the sample is dried for about 60 hours at a pressure less than
100 microns. While the drying time was established by trial and error, a thermocouple probe placed in the center of the sample itself could be used to determine when the sample is dry, since the sample will remain colder than the shelf so long as ice is still present. After drying, the sam ple is of course removed from the vacuum chamber for further processing. 42
To form a barium ferrite material successfully by such a procedure, the chemicals selected must be mutually soluble in water and the resultant solution must have a moderate freezing point. If either the barium or iron compound is insoluble, mixing is limited by the particle size of the insoluble compound and cannot be accomplished on a molecular scale. Also, a solu tion which has a very low freezing point usually melts during the drying op eration. While this does not seem to be critically dependent on the original concentration of the solution, melting occurs sooner when the solution is more concentrated. Thus, it is apparent that only certain ionic species are suitable for freeze drying. Several chemical combinations were studied in selecting a system for freeze drying barium ferrites, as listed in Table 6 along with the diffi culties encountered for each system. The use of versene acid (a chelating agent) to form a soluble ferric versene salt was an attempt to prevent a re- 3+ - action between Fe . and (CgH^Og) resulting in a precipitate. However, the synthesis reaction was rather sluggish, and no fruitful results were attained. The main problem involved in synthesis was the difficulty of driving the reaction to completion. Most synthesis reactions took the following form: AX + BY -> AY + B+ + X" An Incomplete reaction leaves A+ and Y” species in the solution as impurity ions in subsequent processing. The complexities involved and the uncertainty of the Ba: Fe ratio when such material is used led to the abandonment of further attempts at chemical synthesis. 43
Table 6. Chemical Systems
Barium Component Iron Component Problems Encountered
BaClg FeClg Melting.
Ba(N03>2 Fe(NOg)g. 9H20 Melting.
BafCgHgOglg Fe^C2H3°2^2 Iron compound not A available; synthesis attempts failed.
Ba
Ba(C2H30 2)2 Ferric versene salt Synthesis of versene compound rather difficult. Ba(CH02)2 Fe(CH02)3 Synthesis of com pounds, presence of impurities.
Ba BaCOg FeS04 * 7 HgO BaSC>4 precipitated. BaO FeS04 • 7 HzO Green precipitate formed. Ba The selection of ferric oxalate, Fe^CgO^g • 6 HgO and barium ace tate, Ba(C2H302)2, as raw materials for the freeze drying process posed some special processing requirements. For example, attempts to dissolve the two materials in the same solution simultaneously failed due to the for mation of a precipitate of unknown composition. Subsequently, however, it was found that dissolving the compounds separately and then mixing the re sulting solutions yielded quite satisfactory results. While a metastable precipitate forms when the solutions are mixed, it redissolves with subse quent stirring to yield a homogeneous solution. An x-ray diffraction pattern for the precipitate revealed it was completely amorphous, so no identifica tion could be made. Batch size for the freeze drying process was 300 grams of material plus 750ml of distilled water. The proportions of ferric oxalate and barium acetate of course varied with the BaO: FegOg ratio of the calcined ferrite desired (Table 7). The ferric oxalate was dissolved in 650ml of water, with the barium acetate dissolved in the remainder. Such an arrangement allowed both materials to be dissolved relatively easily, although the ferric oxalate required some heating to facilitate solution (solution temperature never exceeded 40°C). After the material was frozen and dried, it was calcined to convert the oxalate-acetate mixture to oxides and subsequently to the barium ferrite structure. Calcining also reduced the fired shrinkage of the material, which was prohibitively high for the uncalcined powder. Differential Thermal Analysis (DTA)and Thermogravlmetrlc Analysis (TGA) were run on various freeze dried powders, with similar curves Table 7. Batch Compositions Ferrite Ratio Fe2 1: 5.0 90.45% 9.55% 1:5.5 91.24% 8.76% 1:5.9 91.79% 8.21% 1:6.0 91.91% 8.09% 1:6.1 92.04% 7.96% 1:6.5 92.49% 7.51% 1:7.0 92.99% 7.01% * Listed as weight percentage of dry batch. 46 observed in all cases. Any adsorbed water present in the sample was lost in the range up to 100°C. Between 200°C and 400°C a very large exothermic peak (possibly two peaks together) on the DTA plot indicated the loss of COg and HgO from the sample as the oxalates and acetates were converted to oxides. The TGA curve showed a 62.4% weight loss during this reaction (calculated value 64.8%). A small exothermic DTA peak at 750°C was pro duced by the conversion of the BaO and FegO^ to barium ferrite (lower than reported for ball milled powders (90)), with no concomitant weight loss on the TGA curve. No other reactions were indicated below 1200°C. All DTA and TGA results were obtained using a system consisting of DTA-A1, TGA-A2, and CR-C2 controller units from Harrop Laboratories, Columbus, Ohio. Since an essentially complete conversion to the ferrite structure was desired, calcination temperatures of 1400°F, 1800°F, 2000°F, and 2300PF were evaluated. The material to be calcined was placed in 8-lnch x 8-inch x 3-inch alumina saggers to a depth of 1 inch and fired in a glo-bar heated furnace with a Barber-Colman cam controller. Calcination was effected on a 24-hour schedule with the material heated to the holding temperature at a rate of 300°F per hour. Stoichiometric ratio material (1:6.0) was selected for the calcination study. Properties measured on the calcined powders were surface area, pressed density, fired density, and magnetic properties. Surface area was measured by nitrogen adsorption with the Micromeritics Instrument noted previously. Pressed density and fired density were determined by standard techniques; the magnetic measurement procedures will be described later. 47 Samples were dry pressed at 25,000 and 50,000 psi using a 7/16- inch die and a Carver Model B laboratory press. However, a pressure of 50,000 psi was found to produce laminations in the specimens, so 25,000 psi was used for all subsequent pressing. The properties obtained for each cal cining temperature are summarized in Table 8. Microstructures of fired specimens for powders calcined at 2000°F and 2300°F showed excessive grain growth with large amounts of porosity, indicating rather poor sintering characteristics. Similar photographs for powders calcined at 1400°F and 1800°F showed much finer grains, again interspersed with porosity. Photographs of the .powders themselves revealed that large, well formed grains were already present before sintering in the case of materials calcined at 2300°F and 2000°F. The 1800°F grains were very small and not well formed whereas no grains were discernible at 1200X magnification in the 1400°F powder. Evaluation of the foregoing led to the selection of 1800°F as the cal cining temperature for the freeze dried ferrite materials in all subsequent investigations. A program to select a proper binder and to specify other processing steps for the preparation of ready-to-press powder was carried on concur rently with the calcination study. The problem was essentially one of low pressed densities due to the inherently low bulk density of the freeze dried powder. Binder selection was based entirely on density, since the unfired samples were acceptably strong even without binder additions. Several binders were considered, with the three finally selected for trial being listed in Table 9. Polyvinyl alcohol (PVA) was mentioned as a 48 Table 8. Calcination Results Calcination Temperature Property Measured 1400°F 1800°F 2000°F 2300°F Surface Area (m2/gm) 13.2 12.3 3.8 2.2 Pressed Density @ 25,000 psi (gm/cc) 2.42 2.60 2.65 2.77 Fired Density (gm/cc) * 3.80 3.78 3.72 3.29 (BxH)' 'max (x 10® g-oe) * 0.59 0.91 0.85 0.76 Hci (oe)* 1569 2513 2171 2395 Sintering Temperature (<>F)* 2360 2280 2280 2225 * Optimum values determined by thermal gradient study. Table 9. Binders Binder Number Source Polyvinyl Alcohol 51-05 E. I. du Pont de Nemours & Co., Inc. Wilmington, Delaware Ammonium Stearate DC-100-A Nopoco Chemical Company Newark, New Jersey Polyethylene Glycol 20,000 Union Cai’blde Corporation Detroit, Michigan 49 binder in the ferrite literature (91, 92), ammonium stearate w as recom mended as being a good lubricant as well as binder, and polyethylene glycol (PEG) had been found to work well in ferrite bodies. Two percent binder additions were introduced in the batches (method described in specimen preparation section) and pressed densities of the fer rite materials were determined (25,000 psi). Since PVA and PEG showed promise, batches containing only 1.0% of these binders were prepared and evaluated in the same manner. Results of the binder trials are shown in Table 10. Ammonium stearate specimens tended to laminate during pressing, so this binder was rejected. PVA and PEG both worked well, with PEG yielding slightly better pressed density for both percentages. Also, the PEG binder was much easier to use since it dissolved in cold water, where as heating was required to dissolve PVA. Hence, PEG in a 2% amount was selected as the binder for use in all subsequent studies. Even though the PEG binder raised the pressed density into the range of 2.6 gm/cc, a somewhat higher pressed density was desired to promote good sintering characteristics with accompanying high fired density and moderate shrinkage. Milling the material (25 gm) in a 1-quart steel mill with 0.5-inch steel balls for two hours produced only minor improvements in compaction. Also, noting that such milling was a potential sou rce of iron contaminants, no further milling studies were conducted. Prepressing the ferrite powder was eventually selected as a suitable method of increasing pressed density. In this approach thecalcin ed material is first screened through a 35 mesh Tyler screen and then pressed at Table 10. Binder Evaluation (Pressed Density (gm/cc)) ______Binder______0%_____ 1%_____ 2% PEG 2.14 2.41 2.60 PVA 2.14 2.39 2.52 Ammonium Stearate 2.14 — 2.59 51 100,000 psi using a 1-inch diameter steel die and a Model FS-160 press (Riehle Testing Machines Division of Ametek, Inc., East Moline, Illinois). Finally, the pressed material is broken up and passed through a 28-mesh screen. Such techniques raised the pressed density in materials containing 2% PEG binder into the 2.95-3.05 range. m . Specimen Preparation Batches for each composition studied were made up of 30. 0 gm pre pressed ferrite material, the selected amount of additives and 12.0 gm PEG binder solution. The binder solution was made up of 47.5% distilled water, 47.5% Solox alcohol (ethyl alcohol), and 5% PEG. The batch was placed in a Waring blender, and a sufficient amount (~200 ml) of 50% distilled water-50% alcohol solution was added to yield a low viscosity mixture. The mixture was blended at high speed for 10 min utes, poured into an evaporating dish, and dried. When dry, the batch was granulated through a 28 mesh screen and was then ready for either dry pressing or hot pressing. Dry pressed specimens were formed at 25,000 psi to a thickness of approximately 0.150 inch using 7/16-inch diameter die. Nine such speci mens were pressed for use in each thermal gradient study. Thermal gradient firings were made in a cam-controlled Model TG-DA Thermal Gradient Furnace (Robert L. Stone Company, Austin, Texas). A heating rate of 115°F/hour to the hold temperature (2350°F in most cases) and a 1-hour soaking period were used. Hot pressed specimens were formed at 10,000psi using a 0 .625-inch 52 diameter zirconia die. Zirconia was selected after a study of various die materials, conducted in an attempt to eliminate the sticking problems en countered. Dylon AE Graphite Coating (Dylon Industries, Cleveland, Ohio) was applied on the die surface to further prevent sticking during hot pressing. A complete description and drawings for the hot press designed and built especially for this investigation are included in the Appendix. Samples were heated at a rate of approximately 700°F/hour to the holding tempera ture, at which point pressure was applied. The hold period was 10 minutes, following which the furnace was shut off, pressure was released, and the specimen was allowed to cool naturally to room temperature. Following sintering, all test specimens were ground to obtain paral lel sides, as needed for the measurement of the magnetic properties. * * IV. Specimen Composition Three compositional groups were considered in evaluating the ferrite materials and additives. In the Initial phase of the evaluation, stoichiometric BaO • 6 FegOg powder was used as the base material. A thermal gradient study was made with the compositions listed in Table 11, all values being weight percents. The results of this initial study led directly to the second series of compositions listed in Table 12. In this case the base material was BaO • 6 FegOg with 2% SiOg present in every composition to enhance sin- . tering and Increase the fired density. A second compositional group evaluated consisted of ferrite materials with varying BaO: Fe2Og ratios, both undoped and with 2%S102- Table 13 is 53 Table 11. Compositions for BaO • 6 FegOg Ferrite with Additives Composition Number Additive Amount of Additive 1 None ___ 2 PbO 2.0% 3 CaO 2.0% 4 Si02 2.0% 09 5 o 0.5% to 6 Si02 4.0% 0.5% 7 Bi2°3 2.0% 8 Bl2°3 9 A12°3 2.0% 10 Ti02 2.0% 54 Table 12. Compositions with BaO • 6 FegOg + SiOg as Base Material Composition Number Additive______Amount of Additive 0.5% 12 A12°3 2.0% 13 A12°3 4.0% 14 A12°3 0.5% 15 Bl2°3 2.0% 16 Bi2°3 4.0% 17 Bi2°3 18 CaO 0.5% 19 CaO 2.0% 20 CaO 4.0% 21 PbO 0.5% 22 * PbO 2.0% 23 PbO 4.0% 24 TiOg 0.5% 25 TiOg 2.0% 26 Ti02 4.0% 55 Table 13. Compositions for Evaluation of BaO: FegOg Ratio Composition Number Ferrite Material 2% SiO? Added 27 BaO • 5 Fe2Og No 28 BaO * 5 Fe2Og Yes 29 BaO *5.5 FegOg No 30 BaO* 5.5Feg03 Yes 31 BaO * 5.9 FegOg No 32 BaO* 5.9F e20 3 Yes 33 BaO • 6.1 Fe2Og No 34 BaO*6.1Fe2Og Yes 35 BaO * 6.5 Fe20 3 No 36 BaO* 6 .5 FegOg Yes 37 BaO * 7 Fe2Og NO 38 BaO • 7 FegOg Yes 56 a tabulation of these bodies. The third phase of the study had as a base the material of optimum ratio as determined in the preceding study; namely, BaO *5.5 FegOg. A series of compositions with various additives and combinations of additives were analyzed, as shown in Table 14. V. Properly Measurement a. Magnetic Properties All magnetic measurements were made using a D. C. Recording Hysteresigraph (F. G. de Roza Electrical Engineering Co., Detroit, Michigan). The general configuration of the electromagnet’s pole pieces is shown in Figure 5. The specimen to be tested is placed between poles of the electro magnet, and the upper plate is lowered until the specimen is held securely (no air gap between specimen and the poles of the electromagnet). A magnetic field of at least 10,000 oersteds is applied to the speci men to saturate it; the current is then shut off and the applied field dissi pates. Next, a negative field is applied and raised to a value of approximately 10,000 oe. This action causes an opposing field to be induced in the sample, and the effect of this Induced field is sensed by a coil buried in the face of the bottom pole piece. The rate of change in the flux bounded by the coil induces a voltage in that coil. A second buried coil provides a measure of the applied field around the sample. Integration and amplification of these two signals produces a plot of a portion of the magnetic hysteresis loop which is charac teristic of the material being tested. Decreasing the negative applied field to zero results In completion of the third and fourth quadrants of the hysteresis 57 Table 14. Final Additive Evaluation, BaO • 5.5 Fe^O^ Base Material Composition Number Additive (s) Amount of Additive(s) 39 PbO 2.0% 40 PbO, Si02 2. 0%, 2.0% 41 TiOz 2.0% 42 2.0% A12°3 43 CaO 2.0% 44 2.0% Bl2°3 45 Si02 0.5% 46 sio2 4.0% 47 PbO 4.0% 48 PbO 0.5% 49 A12°3 0.5% 50 A12°3 4.0% 0.5% 51 Bi2°3 52 4.0% Bi2°3 53 Bi2Og, PbO 1.0%, 1.0% 54 1.0% Bi2°3’ A12°3 1.0%, 1.0% 55 Bi2°3« Sl02 1.0%, 56 S102, PbO 1.0%, 1.0% 57 sio2, ai2o3 1.0%, 1.0% 58 Al2Og, PbO 1.0%, 1.0% Figure S Kyatentgnph Klectnxuguet. 59 loop; similar application of a positive field finishes the entire loop. The complete loop is plotted on an x -y recorder in a test time of about 30 seconds. Measuring the values of the magnetic properties from the hysteresis loop is accomplished graphically, as illustrated in Figure 6. First, two sets of parallel lines are drawn in conjunction with the sides of the hystere sis loop, forming a parallogram. The intersection of the diagonals of the parallelogram locates the center of the hysteresis loop. The remanence, Br , of the magnet is determined by measuring the vertical dimension through the center as shown and multiplying by a constant value which de pends on amplification and other equipment factors. Similarly, the coer cive force, H , is determined by measuring the horizontal dimension and c multiplying by a different constant value. Next, a line is drawn parallel to the two sides of the parallelogram as shown, intersecting the center of the hysteresis loop and a third side. A vertical line is drawn from this intersection with the third side, and intrin sic coercive force, H is determined from the x-distance between the ci vertical line and the center of the hysteresis loop. The proportionality con stant between this measurement and Hcj once again is dependent upon the test equipment configuration. The maximum energy product, (B x H)max, may be ascertained by multiplying the B_ value by the H value and dividing the resulting product 17 C by 4. Such a technique is made possible by the fact that the second quad- rant of each of these hysteresis loops closely approximates a straight line. 3 r Figure 6. Hysteresis Loop Measurement Technique { Oi © 61 b. Fired Density The fired density of ferrite samples was measured in a standard method, similar to ASTM C-373. Samples were weighed dry on a Sartorius analytical balance (Sartorius-Werke, Germany), placed in water under a vacuum for 2 hours (instead of boiling them), and reweighed suspended in water and then simply as saturated. Density is calculated according to the formula Bulk specific gravity = ^/(w - S) where D is dry weight, W is wet weight, and S is suspended weight. . c. Microstructure Microstructure of the fired ferrite specimens was examined by means of a scanning electron microscope, Model JSM-2 (Japan Electron Optics Laboratory Co., Inc.). The specimens were fractured in two directions, parallel and perpendicular to the direction in which they were pressed. The surfaces were then metallized with gold in an SC-3 High Vacuum Evaporator (Optical Film Engineering Co., Philadelphia, Pa. - now a division of The New York Air Brake Co., Camden, N. J .), to produce clearer photographs. RESULTS AND DISCUSSION I. Material Preparation The results of the material preparation study presented in detail earlier clearly indicated that not all chemical species may be freeze dried effectively. However, only broad guidelines can be given since no attempt was made to accurately determine the freezing point of each solution. To be successful, the solution being freeze dried must have a moderately high freezing point. For example, melting was observed when an attempt was made to freeze dry a solution of ferric sulfate and barium acetate. Since the melting point of this solution was approximately -15°C, it is indicated that the freezing point should at least be above this value. Melting seems to be independent of the initial solute concentration, presenting a somewhat paradoxical situation, since a solution’s freezing point is normally recognized to be dependent on solute concentration. The explanation is apparently related to the fact that water is removed during the drying process, thereby yielding a frozen solution whose solute concen tration increases with the passage of time. The melting observed is prob ably associated with the solute concentration, since the time which elapses before melting occurs is inversely related to the initial solute concentration. It thus appears that the success of a freeze drying operation ap parently depends more on the physical characteristics of the saturated 62 63 solution of the chemical system involved rather than the characteristics of the initial solution. While the specific parameters involved were beyond the scope of this investigation, a detailed investigation of the freeze drying technique would be useful to the selection of favorable chemical systems for freeze drying. Selection of the calcination temperature for production of the experi mental ferrites involved a consideration of surface area, microstructure, sintering characteristics, and evaluation of DTA results. The first step was the differential thermal analysis of the freeze dried powder, noting the temperatures at which various reactions and phase changes occurred. Since a material fully converted from oxides to the ferrite structure was desired, it was felt this information would be helpful in effecting optimum magnetic properties and low firing shrinkage. The normal procedure involved calcination of the freeze dried powder above its indicated recrystallization temperature and then evaluating both pressed and fired densities. Pressed density was considered important be cause sintering characteristics are recognized to be influenced by the com paction of the pressed part. It was actually observed that low pressed density resulted in low fired density no matter what calcination temperature was used. Maximum densities attained in the gradient fire were 3.53 gm/cc, 3.97 gm/cc, 4.51 gm/cc, and 3.87 gm/cc for material calcined at 2300°F, 2000°F, 1800°F, and 1400°F respectively. Surface area and a gradient study were used to evaluate the reactivity of the calcined material. Surface area decreased as calcining temperature was increased, with a sharp drop noted between 1800°F and 2000°F. Magnetic properties showed peak values with material calcined at 1800°F. Material calcined at 2300°F sintered poorly, indicating low reactivity. The low firing temperature at which optimum magnetic properties were attained indicated that relatively large grains were already present in the calcined material and that sintering at higher temperatures produced undesirable grain growth. At the other extreme, the material calcined at 1400°F was fine grained and seemed rather reactive, but low magnetic property values resulted for this material, perhaps due to the large amount of porosity in the pressed parts before firing. As noted earlier, the maximum density attained in gradient firing this material was 3.87 gm/cc (vs. 4.51 gm/cc for 1800°F calcined material), so high density sintering was not possible. While evaluation of the reactivity of a material is, at best, an obscure pro cess, research with respect to hard ferrites is needed in an attempt to increase reactivity and lower sintering tempei'ature. Attainment of a high pressed density with the powder calcined at 1800°F posed a definite experimental problem. As previously mentioned, the first attempts resulted in densities of about 2.15 gm/cc, which yielded fired specimens of very low density. Various binder additions finally led to the selection of polyethylene glycol (PEG) 20,000, raising the pressed density to 2.6. Further enhancement of pressed density was accomplished by prepressing the batch material at 100,000 psi. Subsequent granulation produced a powder with a bulk density over twice that observed before pre pressing; and when this material with binder was pressed, green density • values of 2.95 to 3.05 were realized. Fired density was likewise improved, as were the magnetic properties. Careful microstructural examination 65 revealed no trace of the granular structure of the material, so the operation was deemed successful and became a part of the standard powder prepara tion technique. The entire powder preparation process as evolved in this investiga tion yielded a very homogeneous material with very low impurity levels, and provided the mechanism for the subsequent determination of the effects of stoichiometry and other interesting compositional variations. II. Compositional Evaluation The magnetic properties were naturally of prime interest in evalu ating various freeze dried compositions. Plotting the hysteresis loop of each sample permitted the determination of remanence (Br), coercive force (Hc), intrinsic coercive force (Hcj), and energy product ( (BxH)max). Of these, the latter two are most Important in motor application for automotive uses and will be examined more closely. Fired density was measured for most specimens in an attempt to establish its relation to magnetic properties. Samples which were particu larly significant or interesting were examined with the aid of the Scanning Electron Microscope. In most cases 2700X magnification produced the best photographs for observation of the ferrite microstructures. a. BaO • 6 Fe2Og Materials The first material produced in this study was the stoichiometric bar ium ferrite, BaO* 6Fe2Og, for which magnetic data results are presented in Figures 7 through 12. The standard material (no additives) possesses a relatively high B„, but H and H , are low indicating excessive grain growth, r c ci 2400 -i Undoped 0.5% SiO, 2200 - 2% PbO 2% Si0o 2000 B 2% CaO 1800 . 1600 - SiO, 2lb0. 2200 2300 2400 T °F Figure 7. Remanence Values for BaO • 6Fe20g + Additives-I caOS 2200 Undoped 2000 * 0.5% BLO, 1800- 2% TiO, 1600 - 1400 - 21'00 2200 2300 2400 Figure 8. Remanence Values for BaO • 6 Fe2Og + Additives-II 2000 ' 1800- 1600 - SiO H CaO 1400 0.5% SiO, 2% PbO 1200 4% SiO, Undoped 2100 2200 2300 2400 T °F 05 Figure 9. Coercive Force Values for BaO • 6 Fe2Og + Additives-I 00 1800- 1600- 1400* HL 1200. Undoped 1000- 2% TiO, 2100 2200 2300 2400 T ° F Figure 10. Coercive Force Values for BaO* 6Fe00„ + Additives-II ' « O o> CO to a W o X X max 0.60- 0.70- 0.80- 0.90. l.OOl Figure 11. Energy Product vs. Intrinsic Coercive Force; BaO* 6Fe20g + Additives-I BaO* 6Fe20g Force; Coercive Intrinsic vs. Product Energy 11. Figure 1500 2000 2500 2%CaO Hci 004000 3000 Undoped 2%PbO 504500 3500 0.5%SiO, 2%SiO, SiO, 5000 o co a 33 H 0.80 O X max 0.40- 0.60- 0.90- 1 . 00 -, iue1. nryPoutv. nrni orieFre BaO 6• Force; Coercive Intrinsic vs. Product Energy 12. Figure 1500 2% A1„0, 2000 0.5%Bio0, 2500 2% TiO, Undoped 3500 2%Bi00, 004500 4000 eO^ + ^ Fe^O Additives-II 50*00 72 As explained earlier, large grains with more than one magnetic domain have muchlower coercive force than single domain particles, and this would seem to be the reason for the increase in HQl with decreasing temperature in all gradient fires. Figure 13 shows the grain structure of this stoichiometric material. The grains are small platelets, well developed and rather thin, accounting for the Hcj of 2950 oersteds by the fact that most grains are probably single domain. The structure shows a considerable amount of porosity, verified by the fact that densities for samples of optimum magnetic properties were 3.8 to 4.0 gm/cc. Two micrographs are shown in order to isolate any orienta tion in the ferrite structure. Such orientation would of course lead to anisotropy of the magnetic properties. Orientation may occur in various ways (93, 94), and the degree of orientation may be gauged by various tech niques (95, 96). In the case of the standard BaO* 6 Fe^Og, no orientation was evidenced in the.photomicrographs. The compositions in this group, BaO • 6 FegOg with various percent ages of additives, are characterized by low Hc values throughout, indicating relatively poor sintering behavior. Densities of these compositions were generally less than 4.5 gm/cc, substantiating this conclusion. Br values tended to decrease with the use of additives. Also, the low B„r values com- bined with low Hc produced lower energy products in most cases when additives were used. A notable exception to the foregoing was the effect of added silica on magnetic properties. Additions of 0.5% and 2% SiOg raised the values of both energy product and intrinsic coercive force. Apparently the SiOg acted W pwlS. FnefcraawfiMMofBaO^FegQg. (*) Parallel to praaaiaf dlrectioi. (b) PerpouUoaUr to pnMiag diraotloi (2700 X). 74 both as a density mg agent and a grain growth inhibitor (Figure 14), since maximum energy product was observed at slightly lower temperatures than in the case of the undoped material. Higher densities were produced by the silica additions, 4.7 -4 .8 gm/cc in the case of 2% SiOg and about 4.6 gm/cc in the case of 0.5% SiOg. Hc increased slightly for 2% SiOg, again due to the inhibited grain growth, while 0.5% SiOg had no appreciable effect on Hc. On the other hand, 4% SiO„ produced improvement only in H ., with a defi- A C l nite degradation in the other magnetic properties. Br was decreased by adding SiOg to the ferrite body, and the drop in B„ increased with increasing SiO„. This decrease is probably due to the r « formation of a glassy silica-rich second phase in the ferrite body. This is substantiated by the fact that Br is decreased in almost every case when additives were used in this study. No other compositions in this series were particularly promising. In general, magnetic properties were rather low, probably due to poor sin tering of the bodies. The densities were generally less than 4.2 gm/cc, ranging as low as 3.3 gm/cc. Specimens containing 2% BigOg exhibited densities up to 4.6 gm /cc, with a higher energy product than other compo sitions in the series, accompanied by a reasonably good Hcj value. b. BaO • 6 FegOg + 2% SiOg Materials Since the above results indicated that poor sinterability and low density were detrimental to the magnetic properties, an additional series of compositions was prepared in which the base material was BaO * 6 FOgOg plus 2% SiOg. Various dopants were added to this base material, and the results are shown in Figures 15 through 29. 78 Figire 14. Fractal® svfaces of BaO* SFegOg 48% SiOj. (a) Parallel to presaiag dlreottaa. (b) Pccpeadlcatar to preaatag diiaettoa. (8T00 X). 2200-1 2100- Undoped 2000 " 1900- 0.5% 1800 ‘ 2100 2200 2300 2400 T °F Figure 15. Remanence Values for BaO • 6 FegOg + 2% SiOg + AlgOg 1900 - 1800 - 1700 * Undoped 1600 - 1500 - 0.5% 2100 2200 2300 2400 T °F Figure 16. Coercive Force Values for BaO * 6 Fe2Og + 2% SiOg + AlgOg 1. 0 0 , Undoped COo r t 0.80- X 0.5% aX a S' X a 0.70- 0.50- 2200 2600 3000 3400 3800 4200 4600 5000 Figure 17. Energy Product vs. Intrinsic Coercive Force; BaO* 6 FegOg + 2% SiOg + Al^Og 2200 n 2100 * Undoped 2000 B_ 1900 * 1800 0.5% 2100 2200 2300 2400 T °F Figure 18. Hexnanence Values for BaO * 6 FegOg + 2% SiOg + BigO^ CO 1900-, 1800- 1700- Undoped 1600 - 1500 - 0.5% 1400 - 2100 2200 2300 2400 T °F Figure 19. Coercive Force Values for BaO • 6 FegOg + 2% SiOg + Bi2°3 o00 rH 0.80 o & 0.70- * X max 1.00 "I 1.00 0.60- 0.90- iue2. nryPoutv. nrni orieFre BaO2%+ * 6FegOg Force; SiOg BigOg+ Coercive Intrinsic vs. Product Energy 20. Figure 50 00 2^ 3 25^0 20*00 15b0 H o ci 5 o 50 00 4^ 50*00 45^0 40*00 35^0 Undoped 0.5% 00 2200 2100. Undoped 0.5% 2000 - 1900 ■ 1800 ' 2ibo 2200 2300 2400 T °F Figure 21. Remanence Values for BaO* 6Fe2Og + 2% SiOg + CaO 2000- 1900- 1800- 1700- Undoped 1600- 2300 2400 6 Fe2°3 + 2% Si°2 + Ca° 00 co co o a W 0. i 0 .8 0 * X max 0.50 H 0.60 0.70 0.70 0.90-1 1 . 00 -, A Figure 23. Energy Product vs. Intrinsic Coercive Force; BaO* 6Fe20g + 2%+ SiOgCaO + BaO* 6Fe20g Force; Coercive Intrinsic vs. Product Energy 23. Figure 2000 2400 2800 Hci 3200 4% 3600 00 4400 4000 Undoped 4800 0.5% oo tft 2200 . 2100 - Undoped 2000 B_ 0.5% 1900 1800 2100 22*00 23*00 2400 T °F Figure 24. Remanence Values for BaO • 6 FegOg + 2% SiOg + PbO 00 cn 2000 1900- 1800 . 1700 * \ Undoped 1600 0.5% 1500 2100 2200 2300 2400 T °F Figure 25. Coercive Force Values for BaO • 6 FegOg + 2% SiOg + PbO a>00 <0 .0 - 0.70 a K 0.80 2 X x max 0.60 .0 ' 0.90 0 0 . 1 M Figure 26. Energy Product vs. Intrinsic Coercive Force; BaO* 6Fe2Og + 2%+ BaO* 6Fe2Og SiOg PbO + Force; Coercive Intrinsic vs. Product Energy 26. Figure 5 ) 2ot)0 15\)0 50 00 3500 3000 2500 . H ci Undoped 00 50 5000 4500 4000 0.5% i i i I -a 00 2200 n 2100 - Undoped 2000 * 1900 - 0.5% 1800 1700 21*00 T 2200 2300 T °F Figure 27. Remanence Values for BaO • 6 Fe 2 ° 3 + 2% SiOg + TiOg 1900 _ 1800- 1700- Hc Undoped 1600 - 1400- 0.5% 4% 41------«— 1400 2100 2300 2400 Figure 28. Coercive Force Values for BaO • 6 FCgOg + 2% SiOg + TiOg 00 CO •© a X X X max .0 * 0.50 .0 - 0.70 0.80 - 0.80 0.90- l.OOn Figure 29. Energy Product vs. Intrinsic Coercive Force; BaO 6• Fe^Og Force; 2%+ Coercive SiOgTiOg + Intrinsic vs. Product Energy 29. Figure 1500 2000 5030(00 2500 . H ci 3500 4000 Undoped 4500 5000 © to 91 In this series of compositions, magnetic properties were still rather low, although different additives produced a range of property values. Alu mina caused a decrease in Br , probably due to the formation of an alumina- silica nonmagnetic phase. Hc values were also decreased, with nearly equal values for all three additions. Energy products decreased with increasing amounts of alumina, all values being lower than for specimens containing only 2% SiOg. was increased by the addition of AlgOg, probably due to inhibition of grain growth by the secondary phase development. Bio0„ produced similar decreases in properly values, except that , U O H . values were also lower when bismuth was added. Decreased densities, ci in the 4.6 -4.7 gm/cc range, probably are responsible for the lower mag netic property values. CaO and TiOg additions yielded no properties of promise, with ex tremely low values in all cases. Densities were very low, ranging from about 4.5 gm/cc for 0.5% additions to 4.2 gm/cc for 4% additions. Hence, these additions were not beneficial to either the magnetic properties or the sintering characteristics of the ferrites. The results for lead oxide varied, with compositions containing 0.5% and 4% PbO having lower properly values than samples with no lead present. However, the composition containing 2% PbO exhibited properties slightly better than for lead free samples. This may be partly due to the entry of 2+ some Pb ions into the ferrite lattice, which is entirely possible since the lead ferrite structure is similar to barium ferrite and the two ions do not differ critically in size. However, some sort of lead silicate second phase must also be formed (possibly lead monosilicate), since increasing the 92 amount of PbO in the bodies caused a corresponding increase in H,.. c. Varied BaO: FegOg Ratio Materials Next, two series of compositions were evaluated in which the BaO: Fe203 ratio was varied. Two sets of samples were prepared, one with no dopants being added and the other series containing 2% SiOg. The results are shown in Figures 30-35. It is immediately obvious that compositions with BaO :FegOg ratios greater than 1:6 were definitely of lower quality than for lower ratios. The densities of these materials never exceeded 4.3 gm/cc, with most values being less than 4.0, indicating very poor sintering behavior. X-ray analy ses revealed the presence of iron oxide in the fired specimens, which ap parently is detrimental to magnetic properties. The addition of 2% SiOg did not significantly improve the properties of these ferrites. Of the remaining ratios evaluated, the 1 :5.5 ratio produced the best magnets. Even though sintering characteristics were not optimized (densi ties near 4.6 gm/cc), Br values peaked at over 2500 gauss, with energy products greater than 1.1x10 6 g-oe. A photographic comparison of the 1:6.5 and 1:5.5 materials is shown in Figures 36 and 37. While the grain sizes are fairly comparable, more orientation is evident in the 1:6.5 material. Many more platelet edges are evident in the surface parallel to the pressing direction. Also, the hexagonal grains are less well formed in the 1:6.5 material, possibly due to the presence _of free FOgOg in the body. The 1:5.0 ratio material was nearly as good as the stoichiometric ferrite, but was inferior to the 1:5.5 material. Specimens for these three 2800 1 2400 1 1:6.5 2000 1600 - 1 : 6.1 1200 2400 2100 2200 2300 T op- CO Figure 30. Remanence Values for Various BaO: Fe20 3 Ratios CO Figure 31. * 1 . 2 0 1 1.00 O. o 0.80 a X 0.60 a 1 : 6.1 0.40 400 800 1200 2obo 2400 2800 3200 Figure 32. Energy Product vs. Intrinsic Coercive Force; Various BaO: Fe2Og Ratios C7ICD 2200 1:5.5 2100 * 2000 * 1:5.9 1900 - 1800 : 1700 . 1600 - T T 2100 2200 2300 2400 T°F Figure 33. Remanence Values for Various BaO • Fe2Og Ratios with 2% Si02 CO C3 1900 n 1800- 1:5.5 1600 1500 2100 2200 2300 2400 T °F Figure 34. Coercive Force Values for Various BaO: Fe20 3 Ratios with 2% SiOg <1CO CD © X ' max .0 * 0.60 0.70 - 0.80 0.90 1.00 Figure 35. Energy Product vs. Intrinsic Coercive Force; Various BaO: FegOg Ratios with 2%with BaO:Ratios FegOg SiOg Various Force; Coercive Intrinsic vs. Product Energy 35. Figure 2000 3000 Hci 4000 1:5.5 1:7 .9 1 5000 : 6.1 00 CO 99 Figaro 96 Fnctara aarfacea of Ba0<6.5 FegOj. (a) Parallel to preaalag diroctioe. (b) Perpeediealar to preaalag directioa (8700 X). too Fiiue 37. Fracturem iC m m s of B«0«.5 FejOj. (a) Parallel to preaalag direction, (b) Perpeadlcalar to preaalag dlreeliae (2700 X). 1 0 1 ratios containing 2% SiOg all had similar properties. However, since the undoped 1: 5.5 material was definitely superior, it was selected as the base material for the final compositional study. d. BaO • 5.5 FegOg Materials The next series of compositions contained 2% amounts of the six additives under investigation, with the results being shown in Figures 38- 40. It is revealed that additions of TiOg and CaO decreased all magnetic properties as compared with the undoped BaO *5.5 FegOg material. Sin- terability was also very poor, with densities never exceeding 4.2 gm/cc for these specimens. Photomicrographs are shown in Figures 41 and 42. The two fracture surfaces of the TiOg material present a striking contrast. Parallel to the pressing direction, well defined individual grains are easily discerned, with considerable porosity in evidence. However, perpendicular to the pressing direction, a very glassy appearing fracture surface is presented. This is the first case of obvious orientation in the ferrite structure. The CaO containing specimens also showed some degree of orienta tion. The fracture surface perpendicular to the direction of pressing shows mainly flat platelet surfaces, while considerably more platelet edges are evident in the other fracture direction. This would indicate that magnetic properties measured parallel to the direction of pressing would be lower than when measured perpendicular to pressing direction. However, once again the magnetic properties were too low to warrant further considera tion of this composition. Based on the foregoing results, TiOg and CaO 2600 Undoped 2100 2000 • 1900 • PbO 1800 • Bio0, SiO, 1700 ■ CaO Undoped 1500 - TiO 2 1400 ' 2100 2200 2300 2400 T ° F Figure 39. Coercive Force Values for BaO" 5 .5Fe20g + 2% Additive 1.25 1.15 ' PbO 0.85 - Undoped CaO SiO, TiO, 2000 3000 4000 5000 H . Figure 40. Energy Product vs. Intrinsic Coercive Force; BaO * 5 .5FegOg + 2% Additive 104 105 Figure 41. Fractare anrfacea of BaO • 5.5 FegOg. +25 TIOj. (a) Panlld to preaalag dlreetioa. (b) Perpeadlcolar to preaalag dlrectloa (2700 X). 106 (b) Figure 42. Fracture surfaces of BaO <5.5 FegOg +2% CaO. (a) Parallel to preaalag dlrectloa. (b) Perpeadlcular to preaalag dlrectloa (2700 X). 107 were eliminated from further additive studies. A1„0„ added to the base material led to a decrease in B , but a 2 3 r ’ substantial increase in H and H .. Apparently the A1_0„ acted as a grain C Cl a o growth inhibitor in raising H and H ., although its presence diluted the C Cl magnetic phase, producing the observed decrease in Br. Photomicrographs of this material (Figure 43) show little evidence of any grain orientation. The grains appear extremely well formed, with the small grain size reflec ting the effect of the AlgOg in the body. Densities were slightly better than for the undoped material, ranging between 4.6 and 4.75 gm/cc. However, some porosity is still visible in the ferrite microstructure. SiOg produced results similar to AlgOg, causing a decrease in Bf and an increase in Hci; however, unlike AlgOg, SiOg did not increase Hc, with the result that lower values of (BxH)' max resulted for this composition. In this case, the photomicrographs (Figure 44) clearly evidence the pres ence of SiOg. The grains are generally much smaller (causing the high Hcj values) and not well formed. The somewhat glassy appearance of the struc ture is due to the development of the SiOg phase between the ferrite platelets. The densities were not high, generally 4 .5 -4 .6 gm/cc, probably due to the large pores in the structure, such as those visible in the fracture surface perpendicular to the pressing direction. The reason for these large discontinuities is unknown. The presence of Bio0o likewise caused a decrease in B , although a o r the decrease was much less than in the previous cases. A sizeable increase was observed in H c , resulting in an improvement in the energy product over that of the undoped material. Hcl was also improved, but not as much as 108 (*) Figwe 43. Pnotore ratfaces of BaO* 5.5 Fa^Qs +8*j A10j < (a) Parallel to prasslag dliectioe. (b) Pwpeadlcalar to presslag dtraottoo (8700 X). Figure 44. Fractare earfacea of BaO* 5.5 FqOg +1% lUO]. (a) Parallel to preeslag dlrectloa. (b) Perpeadicalar to preaalag dlrectloa (8700 X). 110 in the case of SiOg additions to the ferrite. The photomicrographs (Figure 45) again show the effect of the addi tive in retarding grain growth of the ferrite, as well as the somewhat ragged appearance of the grains. Densities were also improved, reaching 4.7 to 4.8 gm/cc for specimens showing optimum magnetic characteristics. While less porosity is evident in the photos, small open pockets may still be seen. BigOg seemed to act as a flux, promoting sintering of the ferrite at somewhat lower temperatures. This resulted in the decreased grain size and higher HQ and values observed. The decrease in Br expected from the presence of a nonmagnetic phase was apparently counteracted in part by the increase in density, the net result being only a slight decrease in Br . Lead oxide was the only additive to produce an actual increase in B , reaching a maximum of 2531 gauss. The H values also were the r c highest attained for this series of compositions, peaking at 2024 oersteds. Accordingly, energy products were also highest for this material, reaching 1.25x10 g-oe. The only negative result was the low Hcl values, although they were 300 to 400 oersteds higher than for the undoped material. How ever, the values were still generally less than 3000 oe. Photomicrographs of the PbO-contalning material are shown in Figure 46. The grains are very well developed and somewhat larger than those of structures containing SiOg and BigOg, with no second phase material visible. It is clearly indicated that some degree of orientation is present, the platelets being oriented with their faces perpendicular to direction of pressing. Densities were nearly the same as for the undoped material, averaging near 4.6 gm/cc. I ll Figve 45. Fracture m tkeei of BaO* 5.5 FegOg +2% BlgOg. (a) Parallel to preasiag dtiecttaa. (6) P op—dlcalar to preaalag flw cttoa (1T00 X). I ll Flgnre 46. Fiutm nthoci otBdM.8 F(|Os PM. (a) Parallel to preaalag direetioa. (b) Perpeadicalar to preaalag dlrectloa. (2700 X). 113 Considering all the data, it appears that the Pb2+ ions are actually entering the ferrite structure, probably replacing Ba2+ ions. This appar ently improves the intrinsic magnetic properties of the ferrite, with little second phase material having been formed. A combination of PbO with a grain growth inhibitor should, it seems, give the optimum ferrite properties. e. BaO • 5.5 FegOg Additive Study Based on the foregoing results, it was decided that a further study of four additives was warranted; namely, PbO, AlgOg, BigOg, and SiOg. Compositions were prepared in which each dopant was added in amounts of 0.5%, 2%, and 4%, and the results for these compositions are presented in Figures 47 through 58. For SiOg, an addition of 0.5% gave the optimum magnetic proper ties. Br decreased with increasing additive content, probably due to dilu tion of the magnetic phase with a nonmagnetic phase. Hc was also highest for the 0.5% additive composition, yielding the maximum energy product. Even though for the 0.5% SiOg composition is lower than for 2% and 4% SiOg, the value is still rather high and is in any case satisfactory. Lead oxide gave optimum results for a 2% addition, with Br , H c , and (Bx H)max all peaking for this composition. increased moderately with increasing additive content, indicating that PbO did have at least a limited effect in decreasing grain growth in the ferrite material. However, the differences in did not indicate any real advantage was gained through an increase in PbO to 4%. 2100 2200 235o 2400 T °F Figure 47. Remanence Values for BaO *5.5 FegOg + SiOg 114 2000 1 1900 ' 1800 H 1700 Undoped 1600 — i— 2100 2200 2300 2400 T ° F Figure 48. Coercive Force Values for BaO* 5.5F e20g + SiOg 115 1.20 - i 1.10 ■ Undoped e 0.70 2000 2500 3000 3500 4000 Figure 49. Energy Product vs. Intrinsic Coercive Force; BaO* 5.5F e20g + SiOg 116 2100 2200 2300 2400 T °F Figure 50. Remanence Values for BaO .5.5 Fe20 3 + PbO 2100 2000 ' 1900 - 0.5% 1800 ■ 1700 ■ Undoped 2100 2200 2300 24*00 T °F Figure 51. Coercive Force Values for BaO * 5.5 Fe^O ^ + PbO 0 0 Undoped Figure 52. Energy Product vs. Intrinsic Coercive Force; BaO • 5.5 Fe90„ + PbO 2600. 2400- Undoped 2200 * 2000 * 1800- 1600 * 2100 2200 2300 2400 T °F Figure 53. Remanence Values for BaO • 5.5 Fe2Og + AlgOg 120 2000 1900 < 1800 ‘ 0.5%, 1700 - Undoped 1600 - 2100 2200 2300 2400 T °F Figure 54. Coercive Force Values for BaO *5.5 FegOg + Al^Og 121 1.15 1.05- co X | 0.95 • Undoped 0.85- 0.5% 0.75- 2000 3000 50004000 Hci Figure 55. Energy Product vs. Intrinsic Coercive Force; BaO • 5.5 Fe2Og + AlgOg 2500-1 2400- Undoped 2300- 0.5% B.r 2200- 2100- 2100 2200 2300 2400 T °F Figure 56. Remanence Values for BaO • 5.5 FegOg + BigOg 123 2100 2000 ‘ 0.5% 1900 ' 1800 ' 1700 - Undoped . 2100 2200 2300 2400 T °F Figure 57. Coercive Force Values for BaO • 5.5 Fe^O ^ + BigOg & X max .5 - 1.05 1.10 1.15 - 1.15 1 1.25 . 0 2 ' - Figure 58. Energy Product vs. Intrinsic Coercive Force; BaO* 5.5FegOg + BigOg+ BaO* 5.5FegOg Force; Coercive Intrinsic vs. Product Energy 58. Figure 5030 4000 3500 2500 Undoped 3000 Hcl t 125 126 Additions of alumina to the ferrite resulted in a decrease in B„r as the percentage of alumina increased, again apparently due to the formation of a secondary phase which dilutes the ferrite phase. Hc peaked for the 2% AlgOg addition, as did the energy product. HQi increased with increasing AlgOg content, further indicating the development of a second phase in the bodies. It is indicated that 2% AlgOg is optimum. The addition of 4% BigOg decreased the magnetic values except for the usual increase in HQl as more second phase formed. Addition of 0.5% and 2% Bl„00 resulted in similar B values, with the 2% composition sin- £ o r tering at a substantially lower temperature. The Hc results were also similar, with the 2% composition being slightly higher. The energy product for the 2% composition was about 5% better than for 0.5% BigOg. However, a 0.5% addition induced the higher H . value, due possibly to Cl better sintering (and larger grains) in the 2% composition. For 4% BigOg, the H . value was similar to that measured for the 0.5% addition, ci f . Dual Additive Materials The above results for the 1:5.5 material with BigOg, AlgOg, PbO, and SiO„ dopants revealed the possibility that a combination of additives a might yield higher magnetic values. A series of compositions were thus prepared using the four additives in pairs, with all compositions containing 2% additives, 1% of each Individual dopant. The results are presented in Figures 59-61. None of these compositions yielded outstanding magnetic values, al though some were better than the undoped material. The Sl-Pb and Al-Pb Figure 59. 2000 1900 ' Bi2°3 + A12°3 A120 3 + Pbo 1800 SiOg + PbO s io 2 + a i2o3 1700 ■ 1600 * Bi2°3 + Pb° Bi2°3 + Sl02 2200 2300 2400 T °F Figure 60. Coercive Force Values for BaO *5.5 Fe2Og + Two Additives 128 CO e H o 1.00 K X max 0.95 ' .5 ' 1.05 1.10 1.15-, Figure 61. Energy Product vs. Intrinsic Coercive Force; BaO* 5 .SFegOg +Two Additives .SFegOg BaO* 5 Force; Coercive Intrinsic vs. Product Energy 61. Figure • 2500 30*00 . H + PbO 3500 4obo B i00 0 + A100, + 0 i00 B 129 130 compositions had the highest Br values, probably due to the entry of lead into the ferrite lattice. Hq values were highest for the Bi-Al and Pb-Al compositions, although all compositions had peak H values in the 1850- w 1910 oersted range, the values being higher than for undoped material but lower than the 2% PbO composition (the best reported). Most interesting is the fact that all compositions yielded very nearly the same values. Highest Hcj values were obtained for Al-Bi, Si-Bi, and Al-Si com positions, which is not surprising since these three additives yielded the highest Hci values when used singly. AlgOg, SiOg, and BigOg were the best grain growth inhibitors evaluated in this study. Only the composition containing the Pb-Al additive had an energy product higher than the undoped 1:5.5 material. Apparently the lead raised the energy product, while the alumina was slightly detrimental to it. However, the presence of AlgOg did raise the Hc^ value to approximate that of compositions containing only lead oxide. Accordingly, it once again appears that the optimum composition is a matter of compromising individual properties to obtain the best overall result. III. Hot Pressing Evaluation Four dopant compositions were hot pressed along with an undoped sample of the BaO* 5.5F6gOg ferrite. The four additives, 2% Al^Og, 2% PbO, 0.5% BigOg, and 0.5% SiOg were used in the BaO • 5 .5FegOg ferrite base. No binder was added to these compositions, thus eliminating any burn-out problems during the pressure sintering operation. The five compositions were hot pressed at 10,000 pst in all cases. 131 with the pressure being applied during the high temperature hold period. A hold period of 10 minutes was selected for use, since it was determined that extending the hold period to 20 minutes resulted in no significant benefit. Results of the hot pressing study are presented in Figures 62-64. In most cases, Br decreased with decreasing temperature, but magnetic values were nevertheless much higher than was the case for dry pressed samples. The composition containing PbO again yielded the high est values, peaking at over 2800 gauss. BigOg doped material exhibited the next highest values, but they were slightly below those of the undoped material. The values for SiOg and AlgOg doped materials were decidedly lower. H values showed no general trend as temperature changed. How- c ever, the undoped composition and the one containing SiOg yielded the lowest values reported. BigOg doped material had intermediate values, with PbO and AlgOg compositions showing the maximum values (above 2450 oersteds). All Hc^ values decreased with increasing temperature, indicating progressive grain growth. SiOg proved to be the best grain growth inhibi tor, followed by BigOg and PbO, and finally the undoped and AlgOg doped materials. In this series, the material containing PbO had the highest energy product, 1.81x10 g-oe, which correlated with the results for dry pressed compositions. BigOg yielded the next highest energy product, reflecting ' the result of Inhibited grain growth accompanied by a relatively low drop In 2800 2% PbO Undoped 2700 2% PbO-N( 2600 0.5% SiO, *r 2500 -I 2400 ' 1700 1800 1900 2000 T °F Figure 62. Remanence Values for Hot Pressed BaO • 5. SFegOg + Additives 132 2600 2400 i 2% PbO 2200 1 0.5% SiO, H 2000 2% PbO-N, 1800 i Undoped 1700 1800 1900 2000 T °F Figure 63. Coercive Force Values for Hot Pressed BaO *5.5 FegO^ + Additives 133 1.90 1.70 ■ 2% PbO- 1.50 2% PbO 1.30 Undoped 0.5% SiO, 1.10 2200 2400 2600 2800 3000 3200 3400 3600 Hci Figure 64. Energy Product vs. Intrinsic Coercive Force; Hot Pressed BaO • 5 .5Fe2Og + Additives 135 B . Both Si0o and Alo0 o doped materials had energy products lower than r a a O for the undoped material. PbO and BigOg doped materials were then hot pressed to the same temperatures in a nitrogen atmosphere to determine the effect of atmosphere on the magnetic properties. A nitrogen flow rate in the furnace of 20 scfh was maintained throughout the hot pressing cycle. The results for these two materials are also included in Figures 62 - 64. It is indicated that no significant differences resulted for the nitrogen atmosphere. Property values seemed to be randomly higher or lower for a given sample lot. The only result worth mentioning Is the maximum energy product of PbO doped g material, 1.88x10 g-oe. Two sets of hot pressed specimens, those containing 0.5% BigOg and 2% PbO (the latter fired in Ng), were further characterized by measur ing magnetic properties on a surface machined parallel to the pressing direction. Properties in this direction were generally lower than for the perpendicular direction, indicating a highly oriented material. Values of B r were decreased an average of 46%, as were H c values, with (BxH) max subsequently reduced by 73%. On the other hand, values increased an average of 15%. However, this is expected since magnetization is taking place in the magnetically hard direction of the grains. This means that the material is harder to magnetize and demagnetize in this orientation. Hence, the hot pressed material is highly oriented and very anisotropic. Densities of most hot pressed specimens were considerably higher than for conventionally sintered materials. The undoped material and com positions containing PbO and AlgOg all showed densities slightly above 136 4.9 gm/cc, while the material containing BigOg peaked near 5.0 gm/cc. Only the SiOg-containing material had lower density values in the 4.5 to 4.6 gm/cc range. It is apparent that the applied pressure led to enhanced sintering in these ferrites without producing excessive grain growth. Hot pressing could be a very useful process for the sintering of such ferrites (if an anisotropic material is desired), since a combination of high density and small grain size yield the optimum magnetic properties. However, the practical and economic problems involved in using hot pressing as a pro duction technique are well known. SUMMARY A process was developed for freeze drying barium ferrite materials to achieve high purity and homogeneity, starting with a chemical solution of ferric oxalate and barium acetate, the only satisfactory combination dis covered. The results of this study along with the prior art suggest that most ceramic materials might well be freeze dried, providing a proper chemical system can be found. Freeze di'ying thus becomes a plausible production technique for making ceramic materials when high purity and/or homogene ity are of utmost importance. The rapidly expanding technology in this field is momentarily providing improved equipment, thus increasing the economic feasibility of the process. The principal obstacles to acceptance of the freeze drying approach are raw material and manufacturing costs, the former being especially germane with respect to development of hard ferrite materials. Neverthe less, future research designed to establish controlling parameters in the freeze drying process (solubility, vapor pressure of the solution, melting point, particle size, time-temperature relationships, total system pres sure, etc.) should provide insight into the possibilities for cheaper raw materials and improved drying equipment. The selection of 1800°F as the preferred calcination temperature in this study was a compromise between pressed density and reactivity. Pre pressing and granulating the material, using 2% polyethylene glycol as a 137 138 binder, was found effective in improving pressed and fired densities of the ferrite materials. As noted earlier, reactivity of the batch material is a characteristic that is very difficult to define precisely. Any evaluation based simply on sintering of the material is somewhat ambiguous, since processing vari ables are introduced. Surface area measurements on the other hand provide some insight into the reactivity, since most reactions are critically depend ent upon the amount of reaction surface present. However, an improved specification is needed, perhaps Involving a quantitative x-ray or thermal analytical technique. Further research in this area, not necessarily re stricted to ferrites, is urgently needed. A series of compositions were first studied in which the base ma terial was stoichiometric BaO • 6 FegOg. These materials had rather poor sintering characteristics, with a composition containing2% SiOg producing the best results. A subsequent study of dopants added to the silica-contain ing composition did not solve the low density problem and resulted in rather undistinguished magnetic properties. Of these compositions, the best proved to be one containing 2% PbO. These results demonstrate that rela tively pure, stoichiometric ferrites have rather poor sintering character istics, as previously noted in the literature (46); however, more attention might be directed to very pure materials in which the stoichiometry is closely controlled to determine the role of minor amounts of impurities in the sintering process, since a better understanding might well lead to a more effective selection of sintering aids. A study of ferrite materials representing various BaO: Fe2 (- >3 ratios revealed that compositions representing greater than the stoichiometric 1 :6 ratio had relatively poor magnetic properties. X-ray diffraction analyses indicated the presence of free iron oxide in these compositions, apparently accounting for their poorer quality. Of the other materials evaluated, a 1:5.5 ratio yielded the best magnetic properties, followed by 1:6 and 1:5. These findings are significant in that various literature sources re port good magnetic properties at various ratios from 1 :5 to 1 :6.4, with optimum properties at 1 :5 .9 or 1:6 (62-64). The discrepancy between this prior art and the 1:5.5 optimum ratio of this investigation might be due to iron oxide introduced during processing by earlier investigators; ball milling might well be especially detrimental in this respect. Analyti cal techniques for determining the BaO: ratio of fired samples are generally not sufficiently precise to allow pinpointing the optimum ratio; obviously, new methods of analysis are needed. However, the low order of contamination coincident with the freeze drying process supports a belief that the ratios reported herein are probably more accurate than those pre viously reported. When the BaO *5.5 FOgOg material was used as the basis for an additive study, TiOg and CaO were Immediately eliminated from further consideration due to the poorer properties resulting. AlgOg and SiOg were found to lower Br values and Increase H ci ., due to the formation of a second, nonmagnetic phase in the bodies. Similar effects were observed for compositions containing BigOg, but the effects were not as extreme. Lead oxide produced higher values of Br , Hc, and (B x H)max, but was 140 lower than for most other doped compositions (although higher than for the undoped material). Some grain orientation was observed in the PbO-doped specimens. It is obvious that a second nonmagnetic phase decreases the Br of a magnet, with a resulting decrease in energy product unless compensated by an increase in Hc . A magnetic second phase might not produce this dele- terious effect. Hc seems to be a function of both grain size and density, while Hcj is more critically dependent on grain size alone. Energy product is dependent upon B and H , but the shape of the hysteresis loop is a defi- r c nite factor in its maximum value. Further study should be directed toward the determination of factors controlling the shape of the loop, allowing ma terials to be designed to optimize the properties desired. This investigation revealed that CaO and TiOg are detrimental to all hard ferrite magnetic properties. SiOg, AlgOg and BigOg function effectively as grain growth inhibitors, but only BigOg produced an increase in energy product. Two percent PbO produced the best observed magnetic properties, with Br = 2530g, Hc = 2020oe, = 2950oe, and (®x H)max = 1.25 x 10® g-oe. Since the compositions containing PbO may be considered as barium-lead ferrites, perhaps a hybrid material would yield optimum magnetic properties. A study of Pb-Ba-Sr ferrite materials could pos sibly lead to the development of better hard ferrites than those based on only one of these divalent cations. No combinations of two additives (PbO, SiOg, AlgOg, and BlgOg) were found outstanding. However, this portion of the investigation was 141 rather cursory, and more attention should be directed toward optimizing the amount and combination of additives to the base ferrite materials. The five selected compositions which were hot pressed in this study exhibited both higher densities and magnetic properties. Once again, the composition BaO • 5 .5FegOg + 2% PbO produced the best magnetic proper ties, with B r = 2800 g, H c = 2450 oe, H ci . = 3175 oe, and (BxHV max = g 1 .8 x 1 0 g-oe. The undoped BaO *5.5 FegOg had better properties than SiOg and AlgOg doped materials when hot pressed. Hot pressing was obviously advantageous in producing high density materials, overcoming a problem encountered throughout the investigation. Accompanied by small grain size, these specimens were markedly superior to dry pressed specimens. Hot pressing in a 100% Ng atmosphere as com pared to air did not significantly affect the magnetic properties; however, atmosphere might play a more significant role in conventional sintering operations when longer times and higher temperatures are involved. Per haps sintering and hot pressing in a vacuum would lead to still better, higher density magnets. Commercially available barium ferrites typically exhibit Br = 2100- 2250g, Hc = 1700-1850oe, Hcl = 3200-3300oe, and (BxH)max = g .90-1.10x10 g-oe. The sintered BaO’ 5.5 FegOg specimens generally fell within this range, with compositions containing 2% AlgOg and 0.5% SiOg being near the upper limits, and those containing 2% BigOg and 2% PbO exceeding these values. Such compositions are certainly worthy of further consideration. 142 In summary, this investigation has produced a method of freeze drying barium ferrite materials which seems to merit serious considera tion for production use. Once raw material costs are reduced, the process becomes especially attractive, mainly due to its simplicity. Problems in effective process and product control might well be minimal using this method of material preparation. In this investigation it was possible to produce materials of con trollable BaO: Fe2Og ratio, and a 1:5.5 value was determined to be optimum. Hopefully, some insight has been provided into the selection of additive materials for producing specified magnetic properties. Finally, it has been shown that forming of ferrite magnets by hot pressing tends to result in better magnets than can be realized through dry pressing and sin tering as normally effected. CONCLUSIONS The following conclusions appear justified by the results of this in vestigation of barium ferrite ceramics as herein reported: 1. A cryochemical freeze drying method of preparation in conjunction with controlled calcination will produce barium ferrite material of high purity and homogeneity amenable to the production of ceramic magnets. 2. The successful application of the freeze drying method of preparing a high quality ferrite material requires the selection of mutually soluble chemicals producing a chemical solution having a freezing point within a prescribed temperature range. 3. Barium ferrites of high quality may be produced by freeze drying a controlled solution of barium acetate and ferric oxalate. • 4. The reactivity, sinterability, and magnetic behavior of freeze dried barium ferrites are strongly dependent on powder calcination tempera ture, whereas improvement may also be realized by predensifylng the calcined products prior to final forming and sintering. 5. Processing parameters such as material reactivity, density attained in pressure forming, and firing conditions are of critical importance as regards final density, structure and magnetic behavior of barium ferrite ceramics. 143 144 6. Either conventional furnace sintering or hot pressing may be used to produce useful barium ferrite ceramics, with the maximum densities and magnetic properties being realized through hot pressing, which yields sintered structures of finer grain size, although oriented in character. 7. The stoichiometry of barium ferrite compositions is an important factor in the production of magnetic materials having optimum charac teristics, and a BaO : FegOg ratio of 1:5.5 provides optimum proper ties, surpassing even the stoichiometric barium hexaferrite; a greater excess of BaO and particularly an excess of FegOg over the stoichio metric amount results in poorer magnetic qualities. 8. Oxide additions in amounts from 0.5 to 4.0% can significantly affect sintering behavior, microstructure and magnetic behavior of barium ferrite ceramics, with amounts of 2% or less producing superior magnetic properties. 9. Lead oxide as an additive in barium ferrite ceramics improves over all magnetic properties; silica and bismuth oxide tend to inhibit grain growth, and to increase intrinsic coercive force while deteriorating remanence values; and both calcium oxide and titanium oxide are generally detrimental to magnetic behavior. 10. The best hard ferrites produced in this study had a composition of BaO • 5.5 FegOg plus 2% by weight of PbO, and these were produced using a calcination temperature of 1800°F and final sintering tempera tures of 2225 °F for dry pressed specimens and 1800°F for hot pressed specimens. 145 11. The optimum magnetic properties realized in this investigation were B = 2800 gauss, H = 2450 oersteds, H . = 3175 oersteds, and r c ci fi (BxH) = 1.81 x 10 gauss-oersteds for hot pressed specimens, ' 'max and Br = 2470 gauss, Hq = 2020 oersteds, Hc- = 2950 oersteds, and (BxH) = 1.25 x 10 gauss-oersteds for dry pressed specimens. ' 'max APPENDIX The hot press constructed for this investigation is shown in Figure 65. The press is hydraulically operated, using air pressure on an oil reservoir. The machine is built with the furnace and insulation mounted on the lower ram. The shelf to support the insulation is built separately from the lower plate which supports the die, furnace, and fused quartz tube sur rounding the furnace, although it is still attached to the lower ram. In this arrangement the insulation moves with the furnace and die, so that the die remains at the center of the insulation (Figure 66). The insula tion was formed by tamping Fiberfrax FC-25 mix into a two part sleeve which surrounds the fused quartz tube containing the die and furnace. The quartz tube, the die, and the furnace were all supported on a brass plate fastened to the lower ram. Another brass plate was placed at the top of the quartz tube as well. A brass bellows was used where the upper punch rod passed through this upper plate, thus maintaining a gas tight environment within the quartz tube. Electrical, gas, and thermo couple lead-ins were through the lower brass plate, once again main taining the gas tight environment. The two brass plates were water cooled by copper tubing wrapped around the edges of the plates. The fused quartz tube was wrapped with a thin nickel sheet to reflect heat back into the inner chamber, thus reducing heat loss through the tube. V 147 Figure 65 Laboratory Hot Press. 148 BRASS BELLOWS BRASS PLATE COOLING COILS QUARTZ TUBE ALUMINA PUNCH ROD QUARTZ FURNACE END PLATE ZIRCONIA PUNCH ROD •FURNACE ZIRCONIA DIE Figure 66. Cross-Sectional View of Hot Press Assembly The furnace was mounted on an alumina tube at the center of the quartz tube. It was made by machining grooves in the outer surface of a 3 .5-inch diameter recrystallized alumina tube (Triangle HR grade, Morganite, Inc., L. I ., N. Y.) and wrapping the tube with . 036 inch di ameter platinum wire. The wire was cemented with Alundum cement to hold it in place. The ends of the furnace were fused quartz, with holes drilled for the punch rods and thermocouple wires. The punch rods from the upper and lower rams to the furnace were alumina, while the die and punch rods inside the furnace were zirconia. With this arrangement, the insulation and quartz tube must be re moved each time a sample is introduced or removed. 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