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The manufacture and characterisation of aligned fibres of the ferroxplana ferrites Co2Z, 0.67% CaO-doped Co2Z, Co2Y and Co2W
Journal of Magnetism and Magnetic Materials 186 (1998) 313—325
The manufacture and characterisation of aligned fibres of the ferroxplana ferrites Co Z, 0.67% CaO-doped 2 Co Z, Co Y and Co W 2 2 2 R.C. Pullar!, S.G. Appleton", M.H. Stacey!, M.D. Taylor!, A.K. Bhattacharya!,* ! Centre for Catalytic Systems and Materials Engineering, Department of Engineering, University of Warwick, Coventry CV4 7AL, UK " Structural Materials Centre, Defence Evaluation and Research Agency, Farnborough, Hants GU14 6TD, UK Received 26 September 1997; received in revised form 9 February 1998
Abstract Gel fibres of Co Z, 0.67% CaO-doped Co Z, Co Y and Co W ferrite were blow spun from an aqueous inorganic sol 2 2 2 2 and collected as an aligned tow blanket, with an alignment comparable to that found in commercial fibres. The fibres were then heated to produce the desired ceramic phases, characterised by various techniques and their ferromagnetic resonance spectra measured. Single phase Co Z was found by X-ray diffraction to form at a relatively low temperature of 2 between 1200°C and 1250°C, and the material exhibited the expected microwave properties. Furthermore, an addition of 0.67% CaO was found to promote the formation of Co Z at an even lower temperature of below 1200°C and delay the 2 exaggerated platy grain growth, which is normally encountered at the onset of formation of the Co Z phase and which 2 results in a mechanically weakened fibre. ( 1998 Elsevier Science B.V. All rights reserved. PACS: 75.50.Gg; 82.70.Gg; 76.50.#g Keywords: Hexagonal ferrites; Ferroxplana; Fibres; Sol—gel; Grain growth; Ferromagnetic resonance
1. Introduction This investigation into the characterisation and microwave properties of aligned Co Z, calcium2 doped Co Z, Co Y and Co W ferrite fibres is part 2 2 2 of an ongoing programme to study the development of ceramic fibres, manufactured by aqueous
* Corresponding author.
sol—gel routes. The inorganic fibres research group at Warwick have previously reported both random [1] and aligned hexagonal ferrite fibres [2] made from an aqueous sol—gel route which show improvements in microstructure over conventional materials fired at equivalent temperatures. The many advantages of sol—gel processes for manufacturing fibrous materials compared to bulk polycrystalline powders and the potential enhancement of their magnetic properties have been discussed in
0304-8853/98/$19.00 ( 1998 Elsevier Science B.V. All rights reserved. PII: S 0 3 0 4 - 8 8 5 3 ( 9 8 ) 0 0 0 9 8 - 5
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our previous publications. The characterisation and morphology of these ferrite fibres is reported here, as well as a comparison of the microwave permeabilities of the fibrous hexagonal ferrites with literature data on the same phases in powder form, where available. Among the new class of planar hexagonal ferrites discovered between 1952 and 1956 by Philips [3,4] were Y ferrite (Ba M Fe O ), W ferrite 2 2 12 22 (BaM Fe O ) and Z ferrite (Ba M Fe O ). 2 16 27 3 2 24 41 Crystallographically these comprise basic units of hexagonal barium M ferrite and cubic spinel ferrites in various combinations, and they retain a hexagonal structure, usually with the direction of magnetisation parallel to the c-axis. However, if the metal M"cobalt(II), then Co Y, Co W and Co Z 2 2 2 are formed, known as ferroxplana ferrites, so called because their preferred direction of magnetisation is at an angle to the c-axis [4]. Co Y has a cone of 2 magnetisation at an angle to the c-axis below !58°C [5], above which temperature it has a preferred plane of magnetisation perpendicular to the c-axis until the Curie point is reached at 340°C [6]. Likewise Co Z prefers the basal plane at room 2 temperature following a transition to this from a cone of magnetisation at !53°C, but it then undergoes a further change to magnetisation parallel to the c-axis from 207°C to the Curie point at 400°C [5]. Co W has a cone of magnetisation at 2 70° from the c-axis at room temperature and remains so up to 180°C, at which point this angle steadily decreases until the direction of magnetisation is parallel to the c-axis at 280°C, where it then stays up to the Curie point at 490°C [7,8]. This means that at room temperature these materials are magnetically soft, because although a large amount of energy is needed to move out of this plane or cone, the magnetic vector can easily rotate within the preferred plane or cone [9]. Therefore these materials are of little use as permanent magnets despite their spontaneous magnetisation and high thermal stability, but they have a much higher permeability and ferromagnetic resonance up to the GHz region compared to the 300 MHz ceiling encountered with the spinel ferrites [10], and this brings them into the microwave region useful for inductor cores and uhf communications [9]. As the crystals of these materials are
naturally isotropic within the basal plane, the permeability of compacted polycrystalline components can be raised even further by the alignment of individual crystals in a preferred direction, without a lowering of the limiting frequency or the loss maximum [11]. These properties make the ferroxplana ferrites potentially ideal for use in non-reciprocal microwave devices, provided the resonance line-width can be reduced to single crystal values and that insertion losses can be reduced, by exploiting the effect of gyromagnetic resonance that appears when the material is subjected to a low power microwave field perpendicular to the axis of magnetisation. The production of these materials as fibres in an aligned form will enable the product to be incorporated into a composite matrix, opening up a new area of novel materials and applications exploiting the useful magnetic and microwave properties of these ferrites. Common commercial methods of producing the ferroxplana ferrites are similar to those for all the hexagonal ferrites [12], with the resulting grain diameter exceeding the critical domain size of 1 lm [13,14]. Low porosity and submicron size are each desirable properties in microwave ferrite components, but the commercial methods used to obtain both, such as hot pressing, cannot be used in the production of a fibrous material [15]. There have been many studies to reduce both grain size and porosity of spinel ferrite powders using nonstoichiometric additives, but very little involving the ferroxplana ferrites. Therefore, in an attempt either to promote the formation of Z ferrite at a lower temperature or to improve the microstructure of the ceramic material a calcium oxide doped Co Z fibre was also produced. 2 2. Experimental 2.1. Sample preparation The spinning solution was made from an acid peptised halogen stabilised iron(III) hydroxide sol. Stoichiometric amounts of the barium and cobalt salts were dissolved into solution with an organic liganding agent, before blending with the iron sol to give the required Co Z composition. Spinnability 2 was bestowed by the addition of a small amount of
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polyethylene oxide spinning aid and after preparation the overall Fe/anion ratio, including the salt addition, was 1.45 : 1. The fibres were produced by a modified proprietary blow spinning process [16] in which the spinning solution is extruded through a row of holes on either side of which impinge parallel jets of attenuating air. The fibres were gelled by mixing in a stream of hot secondary air and then collected on a rotating drum as an aligned blanket. After collection the fibres were removed from the drum and stored in a circulating oven at 110°C. The gel fibres were then heat treated in a muffle furnace, firstly being pre-fired to 400°C at 100°C/h to remove water and organic compounds, and then further heat treated at 200°C/h up to 1000°C in a re-crystallised alumina vessel for three hours. The fibres were further fired to between 1100 and 1300°C on a platinum foil sheet for three hours to form Co Z ferrite. Platinum was found to be neces2 sary for the vessel material when firing above 1100°C to prevent cobalt migration into the alumina boat and the subsequent loss of stoichiometry. The calcium-doped Co Z fibres were produced 2 as above, but with a calcium salt also added to the iron(III) sol along with the other salts in the ratio of 1 : 10 of Ca : Ba, which was equivalent to CaO"0.67% by weight of the final ceramic. These fibres were also spun and collected as an aligned tow blanket, and then stored and fired under the same regime. Co Y and Co W ferrites were also prepared by 2 2 similar methods as previously reported [1,2], and they too were both collected as an aligned tow blanket. The Co Y and Co W gel fibres were then 2 2 first heated to 400°C and then fired to 1000°C in a recrystallised alumina vessel, and the Co W gel 2 fibres were also heated on a platinum sheet at temperatures above 1000°C, as above. All the fibres were then characterised and compared, and their ferromagnetic resonance (FMR) frequencies measured. 2.2. Characterisation 2.2.1. Photon correlation spectroscopy (PCS) Particle size measurement of the sol above the 3 nm diameter range was measured on a Malvern
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Instruments Lo-C autosizer and series 7032 multi-8 correlator, using a 4 mW diode laser, 670 nm wavelength. 2.2.2. Assessment of fibre alignment The fibres were collected as a blanket on a high speed rotor, in a manner similar to that used for ‘Safimax’ alumina fibres [17]. There were differences compared to ‘Safimax’, which was an aligned blanket with 90% of the fibres within $10° and all within $20°. 1. A small proportion of fibres crossing the general alignment was estimated using an optical microscope at 40]magnification. The number of aligned fibres in a field was counted, together with the few crossing the alignment in the same field. Counts were made in up to 5 separate fields summing to several hundred generally aligned fibres and up to about 50 crossing the alignment. 2. The direction of the generally aligned fibres was analysed by traversing the electron micrographs with a protractor normal to the axis of alignment and measuring the deviations of at least 100 individual fibres. Two sets of data were taken from opposite sides of each micrograph. 2.2.3. X-ray photoelectron spectroscopy (XPS) The XPS analysis was performed using a Kratos XSAM 800 spectrometer fitted with a dual anode (Mg/Al) X-ray source and a multichannel detector. The spectrometer was calibrated using the Ag3d5@2 line at 397.9 eV and the AgMVV line at 901.5 eV. Al K radiation (1486.6 eV) was the exciting source a (120 W) and spectra were collected in the high resolution mode (1.2 eV) and Fixed Analyser Transmission (FAT). The Kratos DS800 software was used for data acquisition and analysis. 2.2.4. X-ray fluorescence spectrometry (XRF) The elemental composition of the samples was measured on a Philips PW2400 sequential X-ray spectrometer fitted with a rhodium target end window X-ray tube, and Philips X-40 analytical software. The samples were analysed in the form of a fused bead, where 1 g of sample was fused with 10 g of lithium tetraborate flux at 1250°C for 12 min and then cast to form a glass bead.
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2.2.5. X-ray powder diffraction (XRD) measurement X-ray powder diffraction patterns of the samples treated at various temperatures were recorded in the region of 2h"10—80° with a scanning speed of 0.25°/min on a Philips PW1710 diffractometer using Cu K radiation with a nickel filter. Cell a parameters were calculated and further refined using linear regression procedures applied to the measured peak positions of all major reflections up to 2h"90° with the Philips APD 1700 software. This software was also used to calculate the average size of the crystallites in a sample using the Scherrer equation. 2.2.6. Scanning electron microscopy (SEM) Scanning electron micrographs and analysis of the morphology of the samples was carried out on a Cambridge Instruments Stereoscan 90 SEM operating at 15 keV. Conducting samples were prepared by gold sputtering fibre specimens. 2.2.7. Surface area and porosity measurements Surface areas and pore size distributions of the fibres were performed on a Micrometrics ASAP 2000 using N as the adsorption gas. Samples were 2 degassed at 300°C for 6 h prior to analysis. 2.2.8. Microwave measurements The complex permeability (k*) spectra of the ceramic ferrite fibres were determined using established techniques. The fibres were milled to a fine powder and dispersed in paraffin wax to a volume loading of 30% ferrite. The wax/ferrite mixture was die-pressed to form a toroidal specimen which was inserted in a coaxial cell. Measurements of transmission and reflection were made between 0.5 and 18 GHz using a Hewlet Packard HP8510C vector network analyser coupled to an HP8515A S-parameter test-set. The values of k* for the ferrite/wax composite were calculated from the measured S-parameters using the transmission line method of Nicholson and Ross [18]. The intrinsic ferrite properties were isolated using the Lichtenecker effective medium expressions: DkD "» DkD #(1!» ) DkD , %&& & &%33*5% & 8!9 tan d "tan d #tan d , %&& &%33*5% 8!9
where DkD , DkD and DkD are the modulus of %&& &%33*5% 8!9 permeability of the composite, the ferrite and wax respectively. » is the volume fraction of ferrite and & tan d is the loss tangent, defined by the relation tan d"kA/k@, where k@ is the real permeability and kA is the imaginary permeability.
3. Results and Discussion 3.1. Doped iron(III) sol and spinning sol characterisation and stability The stoichiometric mixture required to produce Co Z contains a relatively high proportion of metal 2 salts as dopants, which not only contribute nothing to gel formation but actually destabilise the sol itself. The stability and the resulting size of the sol particles are also sensitive to the preparative techniques and conditions employed, and PCS enabled us to measure and control the properties of the sol to a certain extent. The PCS data indicated that the average particle size of the Co Z precursor sol was 2 6.8 nm, with a polydispersity of 0.690 and a weight average of 4.8]104 amu. The volume distribution of the sol particles has a direct effect on the spinning process, and even a small number of large particles can severely impede or even nullify the spinnability of the sol. Therefore volume distribution is considered a more relevant measure than the Z average, and the volume average was found to be 8.1 nm, with an upper limit of 40 nm. The Ca-doped Co Z precursor sol had a slightly 2 improved volume average of 6.2 nm and an upper limit of 35 nm, but this appeared to be an insignificant difference that had no observable effect during the spinning process. The actual spinning sol was also studied, after the addition of the spinning aid and concentration to the required viscocity. This showed a doubling in average particle size and weight average, but despite this the sol remained spinnable due to only a slight increase in the upper limit of particle size. The sizes of the Co Z and 2 Ca-doped Co Z were similar to those of the Co Y 2 2 and Co W sols, as reported previously [1,2], and 2 all are compared in Table 1. It must be considered
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Table 1 Characterisation of the iron(III) sol with the addition of stoichiometric amounts of metal salts for Co Z, 0.67% CaO-doped Co Z, 2 2 Co Y, Co W and the Co Z spinning sol 2 2 2 Formula of sol
Z Av. (nm)
Vol. Av. (nm)
Vol. Tail (nm)
Polydispersity
Co Z sol 2 Ca-doped Co Z sol 2 Co Z Spinning sol 2 Co Y sol 2 Co W sol 2
6.8 7.1 11.2 8.4 6.8
8.1 6.2 14.1 7.5 7.5
40 35 45 40 35
0.690 0.578 0.741 0.524 0.679
Wt. Av.]104 amu 4.8 5.2 10 7.4 4.4
that the technique is unable to detect particles below the 3 nm threshold, and therefore these measurements may be greater than the actual true figures. After addition of the spinning aid both the Co Z and Ca-doped Co Z sols were spun and 2 2 collected in aligned form, as were the Co Y and 2 Co W fibres, and all were then calcined to form the 2 ceramic ferrite fibres. 3.2. Alignment of the fibres The proprietary ‘Safimax’ fibre process [17] was modified to allow the collection of more friable gel fibres. The collection rotor was operated well below, rather than slightly above the fibre generating velocity, and an open diverging rather than a converging air duct was used. As a result 5—15% of the ferrite fibres cross the general alignment because of looping on to the rotor. The remaining 85—95% are well aligned as can be seen from Figs. 1 and 2, which show the Co Z and Co Y aligned ferrite 2 2 fibres respectively. Both the protractor measurement on micrographs or direct measurement at 40]magnification are viewing deviations set into the fibre on a 1—2 mm scale. These cannot be removed by subsequent tensioning and can affect the packing into composites. The Co Z and Ca-doped Co Z fibres behaved 2 2 identically during the spinning procedure, and the addition of calcium conferred no advantage to the alignment or quality of the gel fibres. The fibre diameter was between 5 and 10 lm in both cases, and the gel fibres were smooth and parallel sided. 97.7% of the fibres were found to be within$20° of the axis of alignment, and 75.6% were within $10°.
Fig. 1. Photograph demonstrating the alignment of 0.67% CaO-doped Co Z fibres fired to 1200°C. 2
Fig. 2. Photograph demonstrating the alignment of Co Y fibres 2 fired to 1000°C.
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The Co Y gel fibres were identical to those re2 ported previously [1], and were equally smooth and even-sided with a diameter of 4—8 lm. They also possessed excellent alignment, similar to that of the Co Z fibres, of 96.1% within $10° of the 2 axis of alignment and 75.6% within $20°, as can be seen in Fig. 2. The Co W gel fibres were as 2 described previously [1] with a diameter of 3—5 lm and a comparable degree of alignment. 95—98% of the fibres reported here are within $20° of the axis of alignment and 75—85% within $10°, and this compares reasonably well with ‘Safimax’ alumina [17], a development product, where all the fibres are within $20 and 90% within $10°. The present fibres are sufficiently well aligned at present to allow a programme of composite evaluation to start. The problems of looping and the slight waviness in the generally aligned fibres will be removed through further optimisation of the spinning process, such as reducing the differential between the fibre velocity and the slower rotational speed of the collection drum. 3.3. Fibre characterisation and morphology Because the ferroxplana compounds have such complex chemical compositions, a deviation in either the stoichiometries or oxidation states of the components can have an adverse affect on their magnetic properties. These were confirmed for the Co Z and Ca-doped Co Z fibres, after both were 2 2 calcined at 1000°C to remove any halide or organic contaminants remaining from the spinning-sol precursor. The XPS analysis of the fibres showed the oxidation state of the surface iron to be Fe(III) with a binding energy of 709.7 eV for the main Fe 2p peak, and the XRF elemental analysis for the oxides BaO, Fe O and Co O confirmed their com2 3 3 4 position to be Ba Co Fe O . The Co Y and 3 2 24 41 2 Co W fibres were likewise characterised by the 2 same techniques and shown to be of the correct stoichiometry, agreeing with previously reported results [1,2]. The XRD patterns for the Co Z fibres taken 2 between 400°C and 1250°C are shown in Figs. 3 and 4. At 400°C hematite had formed as the only identified crystalline phase from the amorphous background and this remained the sole component
Fig. 3. XRD patterns of Co Z fibres fired to (a) 400°C, (b) 600°C 2 and (c) 800°C for 3 h.
at 600°C (Fig. 3a and b). However, at 800°C the hematite peaks had either been replaced or hidden by the appearance of BaFe O (BaM ferrite) and 12 19 CoFe O (cobalt spinel, S ferrite) as the major 2 4 phases (Fig. 3c). Due to the presence of halides in the sample below 1000°C these XRD patterns were taken using a glass sample holder, and this is responsible for the large background hump between 2h"15—40°. The formation of these two phases between 600°C and 1000°C, coincidental with the loss of hematite, paralleled the formation processes observed for the Co Y and Co W fibres previously 2 2 [1,2], and as with these fibres, there was no discernible surface microstructure up to these temperatures. Nitrogen adsorption measurements taken at 800°C gave the surface area of the fibres as 7.9 m2 g~1 and a porosity of 0.07 cm3 g~1, with an average pore size of at least 40 nm.
R.C. Pullar et al. / Journal of Magnetism and Magnetic Materials 186 (1998) 313—325
Fig. 4. XRD patterns of Co Z fibres fired to (a) 1000°C, (b) 2 1100°C, (c) 1200°C and (d) 1250°C for 3 h.
The XRD patterns taken at 1000°C or above used a polished silicon wafer to mount the sample, and therefore contained no such background hump. By 1000°C the spinel had been lost due to the formation of Co Y, and the two ferrites BaM 2 and Co Y appeared to co-exist in equal propor2 tions (Fig. 4a). This is an unavoidable step in the production of Co Z [19], which is stoichiometri2 cally made up from one unit of each of BaM and Co Y; the crystallisation of Co Z has never been 2 2 reported without undergoing either the formation or mixing of these two phases first, and then only when there is a further increase in temperature to at least 1200°C. The relationships between the components of this system are notoriously complex [20] and the precise mechanism is unknown for the
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formation of Co Z, although it has been suggested 2 that it proceeds via a topotactic reaction between the interfaces of the phases involved [21]. The fibres still had smooth parallel sides at 1000°C, although a rapid increase in grain size had already been observed above 800°C, the microstructure consisting of what appeared to be a mixture of needles 0.5—3 lm long and large hexagonal plates which were submicron in thickness but between 2 and 5 lm in diameter (Fig. 5). This observed mixture of two distinct crystalline mixtures was not seen during the formation of previous ferrite fibres, and it is quite possible that the ‘needles’ were in fact just hexagonal plates viewed edge-on. By this temperature the surface area had decreased to 1 m2 g~1 with a further loss of porosity, indicating a high degree of sintering concomitant with the formation of the Co Y phase. 2 These two phases persisted to over 1100°C (Fig. 4b) with a further increase in both the size and number of the hexagonal plates but the material retained a fibrous nature despite this. The formation of Co Z at 1200°C (Fig. 4c) seemed to coincide 2 with the loss of BaM, but not Co Y, which co2 existed with Co Z in approximately equal 2 amounts. Further more it remained a major phase above 1200°C until the formation of single-phase Co Z, and although this is concurrent with the 2 published data it remains unclear why the two precursor ferrites should not decompose together
Fig. 5. SEM micrograph of Co Z fibres fired to 1000°C for 3 h. 2
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equally to form Co Z. Accompanying the creation 2 of Co Z was a rapid and exaggerated form of 2 crystal growth, which compromised the integrity of the fibres due to the formation of large hexagonal plates 1—2 lm thick and up to 20 lm in diameter, several times the actual diameter of the fibres (Fig. 6). Amazingly, although the material had been mechanically weakened, the elongated growth occurred along the fibre axis and they remained as discrete, separate fibres of between 4 and 7 lm in diameter despite resembling a stack of building bricks about to topple over. This is similar to the rapid grain growth seen in the formation of Co W 2 fibres previously [2], and was correspondingly accompanied by a drastic lowering in surface area and surface porosity to negligible values. This exaggerated grain growth is another common feature of the crystallisation of Co Z, sometimes resulting in 2 a grain size as large as hundreds of microns [22] along with a corresponding loss of porosity, but it has not been noted to occur at temperatures as relatively low as these. It has been reported to happen after 10 h at 1260°C and only 10 min at 1310°C, but no growth at all has been previously observed even after heating for 24 h at 1220°C with conventionally prepared ceramic specimens [23]. Although low porosity is a desirable property in microwave ferrites, so is a small grain size. Furthermore, the rapid growth process may also enclose pores within the structure of the material between old grain boundaries, possibly explaining why
Fig. 6. SEM micrograph of Co Z fibres fired to 1200°C for 3 h. 2
densities less than the theoretical maximum are obtained even with fully sintered Co Z [23]. Therefore 2 a method of preventing or limiting the exaggerated platy growth process would be extremely advantageous in the manufacture of these materials. From the XRD data at 1250°C the fibres appeared to consist of only Co Z (Fig. 4d), a relative2 ly low temperature for single phase Co Z to occur, 2 but it was thought unlikely that 100% pure Co Z 2 had been formed, as the system is so complex and there are usually either traces of Co Y from the 2 formation process, or Co Z begins to decompose to 2 Co W [21]. Due to the sharing of many peak 2 positions the signals of small amounts of these compounds are extremely difficult to distinguish from the major phases in the very convoluted and similar XRD spectra of the hexagonal ferrites. The average crystallite size of the Co Z ferrite fibre was 2 calculated to be 0.2 lm using the Scherrer equation on the 100% peak at 2h"32.72°. There was no evidence of the Co Z fibres decomposing to give 2 Co W either at temperatures up to 1300°C or 2 through a prolonged heat treatment of 24 h, as has been previously reported to happen with firing times of as little as four hours at 1200°C [21]. There was also no worsening in the morphology of the fibres during the transition from a mixed to a single phase material, although on firing to 1300°C some individual fibres fused together into a larger crystalline mass. The 0.67% CaO-doped Co Z fibres behaved 2 exactly as the undoped Co Z fibres up to 1100°C 2 (Fig. 7a), consisting of equal amounts of BaM and Co Y ferrite and having a similar morphology to 2 the Co Z at this temperature. However, at 1150°C 2 the Co Z phase had already begun to form with the 2 corresponding loss of the BaM phase, with the XRD data indicating roughly equal amounts of Co Z and Co Y (Fig. 7b), just as with the Co Z 2 2 2 fibres at 1200°C. However, the microstructure of the Ca-doped fibres was drastically different, with the runaway platy growth process seemingly controlled resulting in more equiaxial looking grains, and maintaining the fibrous characteristics of the fibres to a higher degree (Fig. 8). The hexagonal grains were 1—3 lm across the hexagonal plane and around 1 lm thick, and the fibres were decidedly more robust and handleable as a result, even as
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Fig. 7. XRD patterns of 0.67% CaO-doped Co Z fibres fired to 2 (a) 1100°C, (b) 1150°C, (c) 1175°C and (d) 1200°C for 3 h.
a thin web, although they were still not as mechanically strong as fibres fired to lower temperatures. By 1175°C, Co Z seemed to be the major of the 2 two components, and the presence of Co Y was not 2 indicated at all by the XRD pattern at 1200°C, indicating that Co Z was the sole phase at this 2 temperature (Fig. 7c and d), although again caution must be used as it is unlikely that pure single phase Co Z ferrite was actually obtained. Unfortu2 nately at this stage the exaggerated grain growth had accompanied the full transformation to Co Z, 2 resulting in the now familiar large, plate-like crystals seen also at this temperature in the undoped Co Z fibres (Fig. 9). Nevertheless, this is a very low 2 temperature for Co Z to form as the major com2 ponent, and the average crystallite size derived from the 100% peak at 2h"32.69° was 97 nm, less than half the size of the undoped Co Z fibres at 2 1250°C. An addition of calcium oxide and silica to spinel ferrites can cause an increase in the resistivity and a corresponding decrease of eddy current loss due to the additives segregating at grain boundaries [24] and forming an insulating layer upon cooling. An addition of CaO to manganese-zinc spinel ferrites has also been reported to decrease grain growth rate and sintering rate by the same segregation process [25]. It is suggested that the rate of growth in the plane perpendicular to the c-axis of the Co Z is being slowed by the segregation of 2
Fig. 8. SEM micrograph of 0.67% CaO-doped Co Z fibres 2 fired to 1150°C for 3 h.
Fig. 9. SEM micrograph of 0.67% CaO-doped Co Z fibres 2 fired to 1200°C for 3 h.
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calcium at the grain boundaries, allowing the growth along the c-axis to ‘catch-up’, thus preventing the exaggerated grain growth stage and resulting in the more equiaxed grains observed. The increase in temperature to 1200°C causes the dissolution of the additive in to the crystal lattice resulting in the onslaught of the exaggerated growth seen at this point. The aligned Co Y fibres were found to be identi2 cal in microstructure and composition to the non-aligned fibres reported previously [1] at the corresponding temperatures. The XRD data indicated that they formed a mixture of a-barium ferrite, cobalt ferrite and M ferrite between 600°C and 800°C, which then proceeded to yield a single phase XRD pattern for Co Y at 1000°C. The crys2 tallite size was calculated to be 56 nm from the Scherrer Equation using the 100% peak at 2h"41.19°. The fibres demonstrated themselves to be made of a densely packed mass of randomly oriented thin hexagonal plates between 1 and 3 lm in diameter and 0.1—0.5 lm in thickness, and were noticeably stronger and more handleable than any of the other fibres which had been fired to produce the Co Z or Co W phases. 2 2 The Co W fibres investigated also behaved like 2 those reported previously [2], forming a mixture of M ferrite and cobalt spinel ferrite at 1000°C which then produced at 1200°C what appeared to be single phase Co W from the XRD data. The forma2 tion of the Co W was also accompanied by the 2 exaggerated grain growth process, forming fibres consisting of lozenge-shaped hexagonal plates, some over 10 lm long, with a resulting loss of mechanical strength, as reported before. Despite this the calculated crystallite size was found to be 60 nm, much smaller than that for the Z ferrite fibres, using the 100% peak at 2h"34.54°.
magnetisation within the ab plane as well as the stiffness out of the ab plane. The corresponding anisotropy fields are H and H respectively, and ( h these are connected to the anisotropy constants K , 1 K (H ) and K (H ). The resonance condition is 2 h 3 ( given by 2pf "cJ(H H ) 3%4 h ( and literature values for H and H will therefore h ( give an indication of the FMR frequency expected for the Z, Y and W hexaferrites studied here. Typically H is much larger than H . For Co Z, h ( 2 H "13 000 Oe and H values have been reported h ( ranging from between 16 and 112 Oe [26]. These compute to f between 1.3 and 3.4 GHz. For 3%4 Co Y, H "28 000 Oe and H &150 Oe [27] re2 h ( sulting in f "5.7 GHz. H of Co W has been 3%4 h 2 reported as being 21 200 Oe [28], but no H values ( are given for these samples. The natural FMR frequency is, therefore, unclear. The measured spectra of k@ and kA for the Co Z 2 fibres fired to 1250°C are shown in Fig. 10. The calculated f was in excellent accord with our 3%4 measurements, and the values of both of k@ and kA at resonance agree closely with those previously reported for conventionally prepared specimens of Co Z [26,29], which for this ferrite all report a res2 onance frequency of approximately 1.3 GHz. This suggested that the sample was single phase Z ferrite, as was indicated by the XRD data discussed earlier.
3.4. Microwave measurements The natural ferromagnetic resonance (FMR) frequency is defined as the point at which kA is a maximum and which coincides with the half peak value of k@, where k@ is the real permeability and kA is the imaginary permeability. In materials with planar anisotropy the natural FMR frequency will be governed by the rotational stiffness of the
Fig. 10. FMR permeability components k@ (real permeability) and kA (imaginary permeability) for Co Z fibres fired to 1250°C. 2
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The spectra for the Ca-doped Co Z fibres fired to 2 1200°C are shown in Fig. 11. Whereas there was a peak in the kA spectrum at 1.3 GHz which suggested the presence of Co Z ferrite, the permeability 2 levels were much lower than for either standard pure Co Z ferrite samples or the Co Z fibres at 2 2 1250°C. It is possible that the reduction in the permeability at low GHz frequencies could be either an intrinsic effect due to the addition of calcium, or caused by the lower reaction-firing temperature used for this sample resulting in a mixedphase ferrite containing a small proportion of Co Y. 2 The Y ferrites are known to form at a lower temperature [13] than the W or Z ferrites, and this would seem consistent with the FMR spectra for the Co Y fibres fired to 1000°C, shown in Fig. 12. 2 A broad resonance in kA was seen, peaking between 6 and 7 GHz, and there was also a possible small resonance at 1.3 GHz, and these results appear to agree with those of bulk polycrystalline samples reported previously [30,31]. As mentioned above, there is no report of the natural FMR frequency of Co W ferrite. However, 2 the k@ and kA spectra of the Co W fibres fired to 2 1200°C (Fig. 13) were in good agreement with values reported in the frequency range 0.03—2 GHz [19]. This reference, which reports measurements made on isotropic polycrystalline samples, shows k@
323
Fig. 12. FMR permeability components k@ (real permeability) and kA (imaginary permeability) for Co Y fibres fired to 2 1000°C.
Fig. 13. FMR permeability components k@ (real permeability) and kA (imaginary permeability) for Co W fibres fired to 2 1200°C.
Fig. 11. FMR permeability components k@ (real permeability) and kA (imaginary permeability) for 0.67% CaO-doped Co Z 2 fibres fired to 1200°C.
to be fairly constant at about 3 GHz and kA starting low and rising with increasing frequency, reaching a maximum at 1—2 GHz. There is no resonance reported, whereas a slight peak in kA was seen in the FMR spectra of the Co W fibres. This suggested 2 the minor presence of a Co Z phase, despite there 2 being no evidence of this in the XRD pattern [2], but this is consistent with previous work on the phases present and the difficulties encountered when attempting to produce pure W or Z ferrite phases [32].
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4. Conclusions Stable iron(III) sols were successfully produced with cobalt and barium doping, giving precursors which are stoichiometrically equivalent to Ba Co Fe O , Ca Ba Co Fe O , 3 2 24 41 0.3 3 2 24 41 Ba Co Fe O and BaCo Fe O . These sols 2 2 12 22 2 16 27 were all found to have a similar particle size to our previously reported ferroxplana-analogue doped precursor sols. Gel fibres of all the ferrites were successfully spun and collected as an aligned tow blanket. The alignment of these fibres was of a degree comparable to that found in commercially produced materials, which was maintained after subsequent heat treatment to form the ceramic ferrite product. The Co Z formulation gel fibres were found to 2 form pure Co Z ferrite at a temperature of 1250°C, 2 and the FMR spectra agreed that the material was indeed single phase, an achievement considering the complex range of similar phases which can co-exist at these temperatures. The seemingly unavoidable exaggerated platy grain growth process was still observed however, compromising the structural integrity of individual fibres. The addition of 0.67% CaO improved both the morphology and formation temperature of the Co Z fibres significantly. The Z phase was a major 2 component of the material at only 1150°C, and a much improved microstructure was observed with no exaggerated grain growth, resulting in smaller, more equiaxed grains even up to 1175°C, by which point Co Z was the dominant phase. It is 2 suggested that was is due to the calcium segregating at the boundary edges and slowing the rate of grain growth in the plane perpendicular to the c-axis. This allowed the growth to proceed in a slower and less two dimensional manner, until the dissolution of calcium into the lattice occurred and the runaway grain growth proceeded as usual. At 1200°C the exaggerated platy growth had occurred and Co Z appeared to be the only phase present by the 2 XRD pattern at this temperature, but FMR data indicated that the microwave properties were not as good as with the undoped Co Z fibres, which 2 could be either an intrinsic property of the calcium doped ferrite or due to the retention of some Co Y 2 as a minor phase.
The aligned Co Y and Co W fibres showed 2 2 identical compositions and morphologies to those reported earlier [1,2], with the Co Y fibres demon2 strating the better microstructure and hence mechanical properties at formation of the ferrite phase, and the Co W fibres demonstrating the exag2 gerated platy growth phase associated with the formation of Co W and Co Z ferrites. The FMR 2 2 spectra of Co Y and Co W fibres were also meas2 2 ured and both agreed with previous results published on polycrystalline forms, although there is very little published data to refer to regarding the high frequency properties of Co W ferrite. The 2 FMR spectra also indicated that the Co W fibres 2 may contain some Co Z phase as an impurity, 2 which is unresolved in the XRD spectra, but this is not unusual for samples of Co W and it is present 2 only as a very minor component.
Acknowledgements Our thanks to D. Croci for surface area measurements and R.C. Reynolds for the XPS and XRD characterisation (both at the Centre for Catalytic Systems and Materials Engineering, Department of Engineering, University of Warwick) and R. Burton for the XRF analysis (Materials Research Institute, Sheffield Hallam University). The data contained herein is published with the permission of the controller of Her Brittanic Majesty’s Stationery Office.
References [1] R.C. Pullar, M.D. Taylor, A.K. Bhattacharya, J. Mater. Sci. 32 (1997) 365. [2] R.C. Pullar, M.D. Taylor, A.K. Bhattacharya, J. Mater. Sci. 32 (1997) 873. [3] J.J. Went, G.W. Rathenau, E.W. Gorter, G.W. Van Oosterhaut, Phil. Tech. Rev. 13 (1952) 194. [4] H. Jonker, H.P. Wijn, P.B. Braun, Phil. Tech. Rev. 18 (1956) 145. [5] J. Smit, H.P.J. Wijn, Ferrites, Philips Technical Library, Eindhoven, 1959, pp. 204—207. [6] J. Smit, H.P.J. Wijn, Ferrites, Philips Technical Library, Eindhoven, 1959, p. 197. [7] D. Samaras, A. Collomb, S. Hadjivasilou, C. Achilleos, J. Tsoukalas, J. Pannetier, J. Rodriguez, J. Magn. Magn. Mater. 79 (1989) 193.
R.C. Pullar et al. / Journal of Magnetism and Magnetic Materials 186 (1998) 313—325 [8] A. Collomb, B. Lambert-Andron, J.X. Boucherle, D. Samaras, Phys. Stat. Sol. A 96 (1986) 385. [9] C. Heck, Magnetic Materials and their Applications, Butterworths, London, 1974, pp. 511—517. [10] I. Gordon, R.L. Harvey, R.A. Braden, J. Amer. Cer. Soc. 45 (1962) 297. [11] A.L. Stijts, H.P.J. Wijn, Phil. Tech. Rev. 19 (1958) 209. [12] H. Stablein, in: E.P. Wohlfarth (Ed.), Ferromagnetic Materials, vol. 3, North-Holland Physics Publishing, Amsterdam, 1982, pp. 462—535. [13] L.M. Catelliz, K.M. Kim, P.S. Boucher, J. Can. Ceram. Soc. 38 (1969) 57. [14] S. Ram, J.C. Joubert, J. Magn. Mag. Mater. 99 (1991) 133. [15] A.L. Stuijts, in: R.M. Fulrath, J.A. Pask (Eds.), Ceramic Microstructures, Wiley, New York, 1968, pp. 460—467. [16] M.J. Morton, J.D. Birchall, J.E. Cassidy, (ICI), UK Pat. 1360200 (1974). [17] M.H. Stacey, M.D. Taylor, (ICI), Eur. Pat. 318203 (1987). [18] A.M. Nicholson, G. Ross, IEEE Trans. Instrum. Meas. 19 (1970) 377. [19] E. Neckenburger, H. Severin, J.K. Vogel, G. Winkler, Z. Angew. Phys. 18 (1964) 65.
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[20] M.A. Vinik, Russ. J. Inorg. Chem. 10 (1965) 1164. [21] S.I. Kuznetsova, E.P. Naiden, T.N. Stepanova, Inorg. Mat. 24 (1988) 856. [22] E.M.C. Huijser-Gerits, G.D. Rieck, D.L. Vogel, J. Appl. Crystallogr. 3 (1970) 243. [23] E.M.C. Huijser-Gerits, G.D. Rieck, J. Appl. Crystallogr. 9 (1976) 18. [24] T. Akashi, Japan J. Appl. Phys 30 (1961) 708. [25] M. Paulus, Phys. Stat. Sol. 2 (1962) 1181 and 1325. [26] J. Smit, H.P.J. Wijn, Ferrites, Philips Technical Library, Eindhoven, 1959, pp. 278—282. [27] J. Smit, H.P.J. Wijn, Ferrites, Philips Technical Library, Eindhoven, 1959, p. 206. [28] R.A. Braden, J. Gordon, R.L. Harvey, IEEE Trans. Magn. 2 (1966) 43. [29] M. Sugimoto, in: E.P. Wohlfarth (Ed.), Ferromagnetic Materials, vol. 3, North-Holland Physics Publishing, Amsterdam, 1982, pp. 428—433. [30] H.S. Belson, C.J. Kriessman, J. Appl. Phys. 18 (1959) 175S. [31] C.R. Buffler, J. Appl. Phys. 33 (1962) 1360. [32] M.A. Vinnik, A.I. Agranovskaya, N.N. Semenova, Russ. J. Inorg. Chem. 12 (1967) 18.
The manufacture and characterisation of aligned fibres of the ferroxplana ferrites Co Z, 0.67% CaO-doped 2 Co Z, Co Y and Co W 2 2 2 R.C. Pullar!, S.G. Appleton", M.H. Stacey!, M.D. Taylor!, A.K. Bhattacharya!,* ! Centre for Catalytic Systems and Materials Engineering, Department of Engineering, University of Warwick, Coventry CV4 7AL, UK " Structural Materials Centre, Defence Evaluation and Research Agency, Farnborough, Hants GU14 6TD, UK Received 26 September 1997; received in revised form 9 February 1998
Abstract Gel fibres of Co Z, 0.67% CaO-doped Co Z, Co Y and Co W ferrite were blow spun from an aqueous inorganic sol 2 2 2 2 and collected as an aligned tow blanket, with an alignment comparable to that found in commercial fibres. The fibres were then heated to produce the desired ceramic phases, characterised by various techniques and their ferromagnetic resonance spectra measured. Single phase Co Z was found by X-ray diffraction to form at a relatively low temperature of 2 between 1200°C and 1250°C, and the material exhibited the expected microwave properties. Furthermore, an addition of 0.67% CaO was found to promote the formation of Co Z at an even lower temperature of below 1200°C and delay the 2 exaggerated platy grain growth, which is normally encountered at the onset of formation of the Co Z phase and which 2 results in a mechanically weakened fibre. ( 1998 Elsevier Science B.V. All rights reserved. PACS: 75.50.Gg; 82.70.Gg; 76.50.#g Keywords: Hexagonal ferrites; Ferroxplana; Fibres; Sol—gel; Grain growth; Ferromagnetic resonance
1. Introduction This investigation into the characterisation and microwave properties of aligned Co Z, calcium2 doped Co Z, Co Y and Co W ferrite fibres is part 2 2 2 of an ongoing programme to study the development of ceramic fibres, manufactured by aqueous
* Corresponding author.
sol—gel routes. The inorganic fibres research group at Warwick have previously reported both random [1] and aligned hexagonal ferrite fibres [2] made from an aqueous sol—gel route which show improvements in microstructure over conventional materials fired at equivalent temperatures. The many advantages of sol—gel processes for manufacturing fibrous materials compared to bulk polycrystalline powders and the potential enhancement of their magnetic properties have been discussed in
0304-8853/98/$19.00 ( 1998 Elsevier Science B.V. All rights reserved. PII: S 0 3 0 4 - 8 8 5 3 ( 9 8 ) 0 0 0 9 8 - 5
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our previous publications. The characterisation and morphology of these ferrite fibres is reported here, as well as a comparison of the microwave permeabilities of the fibrous hexagonal ferrites with literature data on the same phases in powder form, where available. Among the new class of planar hexagonal ferrites discovered between 1952 and 1956 by Philips [3,4] were Y ferrite (Ba M Fe O ), W ferrite 2 2 12 22 (BaM Fe O ) and Z ferrite (Ba M Fe O ). 2 16 27 3 2 24 41 Crystallographically these comprise basic units of hexagonal barium M ferrite and cubic spinel ferrites in various combinations, and they retain a hexagonal structure, usually with the direction of magnetisation parallel to the c-axis. However, if the metal M"cobalt(II), then Co Y, Co W and Co Z 2 2 2 are formed, known as ferroxplana ferrites, so called because their preferred direction of magnetisation is at an angle to the c-axis [4]. Co Y has a cone of 2 magnetisation at an angle to the c-axis below !58°C [5], above which temperature it has a preferred plane of magnetisation perpendicular to the c-axis until the Curie point is reached at 340°C [6]. Likewise Co Z prefers the basal plane at room 2 temperature following a transition to this from a cone of magnetisation at !53°C, but it then undergoes a further change to magnetisation parallel to the c-axis from 207°C to the Curie point at 400°C [5]. Co W has a cone of magnetisation at 2 70° from the c-axis at room temperature and remains so up to 180°C, at which point this angle steadily decreases until the direction of magnetisation is parallel to the c-axis at 280°C, where it then stays up to the Curie point at 490°C [7,8]. This means that at room temperature these materials are magnetically soft, because although a large amount of energy is needed to move out of this plane or cone, the magnetic vector can easily rotate within the preferred plane or cone [9]. Therefore these materials are of little use as permanent magnets despite their spontaneous magnetisation and high thermal stability, but they have a much higher permeability and ferromagnetic resonance up to the GHz region compared to the 300 MHz ceiling encountered with the spinel ferrites [10], and this brings them into the microwave region useful for inductor cores and uhf communications [9]. As the crystals of these materials are
naturally isotropic within the basal plane, the permeability of compacted polycrystalline components can be raised even further by the alignment of individual crystals in a preferred direction, without a lowering of the limiting frequency or the loss maximum [11]. These properties make the ferroxplana ferrites potentially ideal for use in non-reciprocal microwave devices, provided the resonance line-width can be reduced to single crystal values and that insertion losses can be reduced, by exploiting the effect of gyromagnetic resonance that appears when the material is subjected to a low power microwave field perpendicular to the axis of magnetisation. The production of these materials as fibres in an aligned form will enable the product to be incorporated into a composite matrix, opening up a new area of novel materials and applications exploiting the useful magnetic and microwave properties of these ferrites. Common commercial methods of producing the ferroxplana ferrites are similar to those for all the hexagonal ferrites [12], with the resulting grain diameter exceeding the critical domain size of 1 lm [13,14]. Low porosity and submicron size are each desirable properties in microwave ferrite components, but the commercial methods used to obtain both, such as hot pressing, cannot be used in the production of a fibrous material [15]. There have been many studies to reduce both grain size and porosity of spinel ferrite powders using nonstoichiometric additives, but very little involving the ferroxplana ferrites. Therefore, in an attempt either to promote the formation of Z ferrite at a lower temperature or to improve the microstructure of the ceramic material a calcium oxide doped Co Z fibre was also produced. 2 2. Experimental 2.1. Sample preparation The spinning solution was made from an acid peptised halogen stabilised iron(III) hydroxide sol. Stoichiometric amounts of the barium and cobalt salts were dissolved into solution with an organic liganding agent, before blending with the iron sol to give the required Co Z composition. Spinnability 2 was bestowed by the addition of a small amount of
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polyethylene oxide spinning aid and after preparation the overall Fe/anion ratio, including the salt addition, was 1.45 : 1. The fibres were produced by a modified proprietary blow spinning process [16] in which the spinning solution is extruded through a row of holes on either side of which impinge parallel jets of attenuating air. The fibres were gelled by mixing in a stream of hot secondary air and then collected on a rotating drum as an aligned blanket. After collection the fibres were removed from the drum and stored in a circulating oven at 110°C. The gel fibres were then heat treated in a muffle furnace, firstly being pre-fired to 400°C at 100°C/h to remove water and organic compounds, and then further heat treated at 200°C/h up to 1000°C in a re-crystallised alumina vessel for three hours. The fibres were further fired to between 1100 and 1300°C on a platinum foil sheet for three hours to form Co Z ferrite. Platinum was found to be neces2 sary for the vessel material when firing above 1100°C to prevent cobalt migration into the alumina boat and the subsequent loss of stoichiometry. The calcium-doped Co Z fibres were produced 2 as above, but with a calcium salt also added to the iron(III) sol along with the other salts in the ratio of 1 : 10 of Ca : Ba, which was equivalent to CaO"0.67% by weight of the final ceramic. These fibres were also spun and collected as an aligned tow blanket, and then stored and fired under the same regime. Co Y and Co W ferrites were also prepared by 2 2 similar methods as previously reported [1,2], and they too were both collected as an aligned tow blanket. The Co Y and Co W gel fibres were then 2 2 first heated to 400°C and then fired to 1000°C in a recrystallised alumina vessel, and the Co W gel 2 fibres were also heated on a platinum sheet at temperatures above 1000°C, as above. All the fibres were then characterised and compared, and their ferromagnetic resonance (FMR) frequencies measured. 2.2. Characterisation 2.2.1. Photon correlation spectroscopy (PCS) Particle size measurement of the sol above the 3 nm diameter range was measured on a Malvern
315
Instruments Lo-C autosizer and series 7032 multi-8 correlator, using a 4 mW diode laser, 670 nm wavelength. 2.2.2. Assessment of fibre alignment The fibres were collected as a blanket on a high speed rotor, in a manner similar to that used for ‘Safimax’ alumina fibres [17]. There were differences compared to ‘Safimax’, which was an aligned blanket with 90% of the fibres within $10° and all within $20°. 1. A small proportion of fibres crossing the general alignment was estimated using an optical microscope at 40]magnification. The number of aligned fibres in a field was counted, together with the few crossing the alignment in the same field. Counts were made in up to 5 separate fields summing to several hundred generally aligned fibres and up to about 50 crossing the alignment. 2. The direction of the generally aligned fibres was analysed by traversing the electron micrographs with a protractor normal to the axis of alignment and measuring the deviations of at least 100 individual fibres. Two sets of data were taken from opposite sides of each micrograph. 2.2.3. X-ray photoelectron spectroscopy (XPS) The XPS analysis was performed using a Kratos XSAM 800 spectrometer fitted with a dual anode (Mg/Al) X-ray source and a multichannel detector. The spectrometer was calibrated using the Ag3d5@2 line at 397.9 eV and the AgMVV line at 901.5 eV. Al K radiation (1486.6 eV) was the exciting source a (120 W) and spectra were collected in the high resolution mode (1.2 eV) and Fixed Analyser Transmission (FAT). The Kratos DS800 software was used for data acquisition and analysis. 2.2.4. X-ray fluorescence spectrometry (XRF) The elemental composition of the samples was measured on a Philips PW2400 sequential X-ray spectrometer fitted with a rhodium target end window X-ray tube, and Philips X-40 analytical software. The samples were analysed in the form of a fused bead, where 1 g of sample was fused with 10 g of lithium tetraborate flux at 1250°C for 12 min and then cast to form a glass bead.
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2.2.5. X-ray powder diffraction (XRD) measurement X-ray powder diffraction patterns of the samples treated at various temperatures were recorded in the region of 2h"10—80° with a scanning speed of 0.25°/min on a Philips PW1710 diffractometer using Cu K radiation with a nickel filter. Cell a parameters were calculated and further refined using linear regression procedures applied to the measured peak positions of all major reflections up to 2h"90° with the Philips APD 1700 software. This software was also used to calculate the average size of the crystallites in a sample using the Scherrer equation. 2.2.6. Scanning electron microscopy (SEM) Scanning electron micrographs and analysis of the morphology of the samples was carried out on a Cambridge Instruments Stereoscan 90 SEM operating at 15 keV. Conducting samples were prepared by gold sputtering fibre specimens. 2.2.7. Surface area and porosity measurements Surface areas and pore size distributions of the fibres were performed on a Micrometrics ASAP 2000 using N as the adsorption gas. Samples were 2 degassed at 300°C for 6 h prior to analysis. 2.2.8. Microwave measurements The complex permeability (k*) spectra of the ceramic ferrite fibres were determined using established techniques. The fibres were milled to a fine powder and dispersed in paraffin wax to a volume loading of 30% ferrite. The wax/ferrite mixture was die-pressed to form a toroidal specimen which was inserted in a coaxial cell. Measurements of transmission and reflection were made between 0.5 and 18 GHz using a Hewlet Packard HP8510C vector network analyser coupled to an HP8515A S-parameter test-set. The values of k* for the ferrite/wax composite were calculated from the measured S-parameters using the transmission line method of Nicholson and Ross [18]. The intrinsic ferrite properties were isolated using the Lichtenecker effective medium expressions: DkD "» DkD #(1!» ) DkD , %&& & &%33*5% & 8!9 tan d "tan d #tan d , %&& &%33*5% 8!9
where DkD , DkD and DkD are the modulus of %&& &%33*5% 8!9 permeability of the composite, the ferrite and wax respectively. » is the volume fraction of ferrite and & tan d is the loss tangent, defined by the relation tan d"kA/k@, where k@ is the real permeability and kA is the imaginary permeability.
3. Results and Discussion 3.1. Doped iron(III) sol and spinning sol characterisation and stability The stoichiometric mixture required to produce Co Z contains a relatively high proportion of metal 2 salts as dopants, which not only contribute nothing to gel formation but actually destabilise the sol itself. The stability and the resulting size of the sol particles are also sensitive to the preparative techniques and conditions employed, and PCS enabled us to measure and control the properties of the sol to a certain extent. The PCS data indicated that the average particle size of the Co Z precursor sol was 2 6.8 nm, with a polydispersity of 0.690 and a weight average of 4.8]104 amu. The volume distribution of the sol particles has a direct effect on the spinning process, and even a small number of large particles can severely impede or even nullify the spinnability of the sol. Therefore volume distribution is considered a more relevant measure than the Z average, and the volume average was found to be 8.1 nm, with an upper limit of 40 nm. The Ca-doped Co Z precursor sol had a slightly 2 improved volume average of 6.2 nm and an upper limit of 35 nm, but this appeared to be an insignificant difference that had no observable effect during the spinning process. The actual spinning sol was also studied, after the addition of the spinning aid and concentration to the required viscocity. This showed a doubling in average particle size and weight average, but despite this the sol remained spinnable due to only a slight increase in the upper limit of particle size. The sizes of the Co Z and 2 Ca-doped Co Z were similar to those of the Co Y 2 2 and Co W sols, as reported previously [1,2], and 2 all are compared in Table 1. It must be considered
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317
Table 1 Characterisation of the iron(III) sol with the addition of stoichiometric amounts of metal salts for Co Z, 0.67% CaO-doped Co Z, 2 2 Co Y, Co W and the Co Z spinning sol 2 2 2 Formula of sol
Z Av. (nm)
Vol. Av. (nm)
Vol. Tail (nm)
Polydispersity
Co Z sol 2 Ca-doped Co Z sol 2 Co Z Spinning sol 2 Co Y sol 2 Co W sol 2
6.8 7.1 11.2 8.4 6.8
8.1 6.2 14.1 7.5 7.5
40 35 45 40 35
0.690 0.578 0.741 0.524 0.679
Wt. Av.]104 amu 4.8 5.2 10 7.4 4.4
that the technique is unable to detect particles below the 3 nm threshold, and therefore these measurements may be greater than the actual true figures. After addition of the spinning aid both the Co Z and Ca-doped Co Z sols were spun and 2 2 collected in aligned form, as were the Co Y and 2 Co W fibres, and all were then calcined to form the 2 ceramic ferrite fibres. 3.2. Alignment of the fibres The proprietary ‘Safimax’ fibre process [17] was modified to allow the collection of more friable gel fibres. The collection rotor was operated well below, rather than slightly above the fibre generating velocity, and an open diverging rather than a converging air duct was used. As a result 5—15% of the ferrite fibres cross the general alignment because of looping on to the rotor. The remaining 85—95% are well aligned as can be seen from Figs. 1 and 2, which show the Co Z and Co Y aligned ferrite 2 2 fibres respectively. Both the protractor measurement on micrographs or direct measurement at 40]magnification are viewing deviations set into the fibre on a 1—2 mm scale. These cannot be removed by subsequent tensioning and can affect the packing into composites. The Co Z and Ca-doped Co Z fibres behaved 2 2 identically during the spinning procedure, and the addition of calcium conferred no advantage to the alignment or quality of the gel fibres. The fibre diameter was between 5 and 10 lm in both cases, and the gel fibres were smooth and parallel sided. 97.7% of the fibres were found to be within$20° of the axis of alignment, and 75.6% were within $10°.
Fig. 1. Photograph demonstrating the alignment of 0.67% CaO-doped Co Z fibres fired to 1200°C. 2
Fig. 2. Photograph demonstrating the alignment of Co Y fibres 2 fired to 1000°C.
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The Co Y gel fibres were identical to those re2 ported previously [1], and were equally smooth and even-sided with a diameter of 4—8 lm. They also possessed excellent alignment, similar to that of the Co Z fibres, of 96.1% within $10° of the 2 axis of alignment and 75.6% within $20°, as can be seen in Fig. 2. The Co W gel fibres were as 2 described previously [1] with a diameter of 3—5 lm and a comparable degree of alignment. 95—98% of the fibres reported here are within $20° of the axis of alignment and 75—85% within $10°, and this compares reasonably well with ‘Safimax’ alumina [17], a development product, where all the fibres are within $20 and 90% within $10°. The present fibres are sufficiently well aligned at present to allow a programme of composite evaluation to start. The problems of looping and the slight waviness in the generally aligned fibres will be removed through further optimisation of the spinning process, such as reducing the differential between the fibre velocity and the slower rotational speed of the collection drum. 3.3. Fibre characterisation and morphology Because the ferroxplana compounds have such complex chemical compositions, a deviation in either the stoichiometries or oxidation states of the components can have an adverse affect on their magnetic properties. These were confirmed for the Co Z and Ca-doped Co Z fibres, after both were 2 2 calcined at 1000°C to remove any halide or organic contaminants remaining from the spinning-sol precursor. The XPS analysis of the fibres showed the oxidation state of the surface iron to be Fe(III) with a binding energy of 709.7 eV for the main Fe 2p peak, and the XRF elemental analysis for the oxides BaO, Fe O and Co O confirmed their com2 3 3 4 position to be Ba Co Fe O . The Co Y and 3 2 24 41 2 Co W fibres were likewise characterised by the 2 same techniques and shown to be of the correct stoichiometry, agreeing with previously reported results [1,2]. The XRD patterns for the Co Z fibres taken 2 between 400°C and 1250°C are shown in Figs. 3 and 4. At 400°C hematite had formed as the only identified crystalline phase from the amorphous background and this remained the sole component
Fig. 3. XRD patterns of Co Z fibres fired to (a) 400°C, (b) 600°C 2 and (c) 800°C for 3 h.
at 600°C (Fig. 3a and b). However, at 800°C the hematite peaks had either been replaced or hidden by the appearance of BaFe O (BaM ferrite) and 12 19 CoFe O (cobalt spinel, S ferrite) as the major 2 4 phases (Fig. 3c). Due to the presence of halides in the sample below 1000°C these XRD patterns were taken using a glass sample holder, and this is responsible for the large background hump between 2h"15—40°. The formation of these two phases between 600°C and 1000°C, coincidental with the loss of hematite, paralleled the formation processes observed for the Co Y and Co W fibres previously 2 2 [1,2], and as with these fibres, there was no discernible surface microstructure up to these temperatures. Nitrogen adsorption measurements taken at 800°C gave the surface area of the fibres as 7.9 m2 g~1 and a porosity of 0.07 cm3 g~1, with an average pore size of at least 40 nm.
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Fig. 4. XRD patterns of Co Z fibres fired to (a) 1000°C, (b) 2 1100°C, (c) 1200°C and (d) 1250°C for 3 h.
The XRD patterns taken at 1000°C or above used a polished silicon wafer to mount the sample, and therefore contained no such background hump. By 1000°C the spinel had been lost due to the formation of Co Y, and the two ferrites BaM 2 and Co Y appeared to co-exist in equal propor2 tions (Fig. 4a). This is an unavoidable step in the production of Co Z [19], which is stoichiometri2 cally made up from one unit of each of BaM and Co Y; the crystallisation of Co Z has never been 2 2 reported without undergoing either the formation or mixing of these two phases first, and then only when there is a further increase in temperature to at least 1200°C. The relationships between the components of this system are notoriously complex [20] and the precise mechanism is unknown for the
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formation of Co Z, although it has been suggested 2 that it proceeds via a topotactic reaction between the interfaces of the phases involved [21]. The fibres still had smooth parallel sides at 1000°C, although a rapid increase in grain size had already been observed above 800°C, the microstructure consisting of what appeared to be a mixture of needles 0.5—3 lm long and large hexagonal plates which were submicron in thickness but between 2 and 5 lm in diameter (Fig. 5). This observed mixture of two distinct crystalline mixtures was not seen during the formation of previous ferrite fibres, and it is quite possible that the ‘needles’ were in fact just hexagonal plates viewed edge-on. By this temperature the surface area had decreased to 1 m2 g~1 with a further loss of porosity, indicating a high degree of sintering concomitant with the formation of the Co Y phase. 2 These two phases persisted to over 1100°C (Fig. 4b) with a further increase in both the size and number of the hexagonal plates but the material retained a fibrous nature despite this. The formation of Co Z at 1200°C (Fig. 4c) seemed to coincide 2 with the loss of BaM, but not Co Y, which co2 existed with Co Z in approximately equal 2 amounts. Further more it remained a major phase above 1200°C until the formation of single-phase Co Z, and although this is concurrent with the 2 published data it remains unclear why the two precursor ferrites should not decompose together
Fig. 5. SEM micrograph of Co Z fibres fired to 1000°C for 3 h. 2
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equally to form Co Z. Accompanying the creation 2 of Co Z was a rapid and exaggerated form of 2 crystal growth, which compromised the integrity of the fibres due to the formation of large hexagonal plates 1—2 lm thick and up to 20 lm in diameter, several times the actual diameter of the fibres (Fig. 6). Amazingly, although the material had been mechanically weakened, the elongated growth occurred along the fibre axis and they remained as discrete, separate fibres of between 4 and 7 lm in diameter despite resembling a stack of building bricks about to topple over. This is similar to the rapid grain growth seen in the formation of Co W 2 fibres previously [2], and was correspondingly accompanied by a drastic lowering in surface area and surface porosity to negligible values. This exaggerated grain growth is another common feature of the crystallisation of Co Z, sometimes resulting in 2 a grain size as large as hundreds of microns [22] along with a corresponding loss of porosity, but it has not been noted to occur at temperatures as relatively low as these. It has been reported to happen after 10 h at 1260°C and only 10 min at 1310°C, but no growth at all has been previously observed even after heating for 24 h at 1220°C with conventionally prepared ceramic specimens [23]. Although low porosity is a desirable property in microwave ferrites, so is a small grain size. Furthermore, the rapid growth process may also enclose pores within the structure of the material between old grain boundaries, possibly explaining why
Fig. 6. SEM micrograph of Co Z fibres fired to 1200°C for 3 h. 2
densities less than the theoretical maximum are obtained even with fully sintered Co Z [23]. Therefore 2 a method of preventing or limiting the exaggerated platy growth process would be extremely advantageous in the manufacture of these materials. From the XRD data at 1250°C the fibres appeared to consist of only Co Z (Fig. 4d), a relative2 ly low temperature for single phase Co Z to occur, 2 but it was thought unlikely that 100% pure Co Z 2 had been formed, as the system is so complex and there are usually either traces of Co Y from the 2 formation process, or Co Z begins to decompose to 2 Co W [21]. Due to the sharing of many peak 2 positions the signals of small amounts of these compounds are extremely difficult to distinguish from the major phases in the very convoluted and similar XRD spectra of the hexagonal ferrites. The average crystallite size of the Co Z ferrite fibre was 2 calculated to be 0.2 lm using the Scherrer equation on the 100% peak at 2h"32.72°. There was no evidence of the Co Z fibres decomposing to give 2 Co W either at temperatures up to 1300°C or 2 through a prolonged heat treatment of 24 h, as has been previously reported to happen with firing times of as little as four hours at 1200°C [21]. There was also no worsening in the morphology of the fibres during the transition from a mixed to a single phase material, although on firing to 1300°C some individual fibres fused together into a larger crystalline mass. The 0.67% CaO-doped Co Z fibres behaved 2 exactly as the undoped Co Z fibres up to 1100°C 2 (Fig. 7a), consisting of equal amounts of BaM and Co Y ferrite and having a similar morphology to 2 the Co Z at this temperature. However, at 1150°C 2 the Co Z phase had already begun to form with the 2 corresponding loss of the BaM phase, with the XRD data indicating roughly equal amounts of Co Z and Co Y (Fig. 7b), just as with the Co Z 2 2 2 fibres at 1200°C. However, the microstructure of the Ca-doped fibres was drastically different, with the runaway platy growth process seemingly controlled resulting in more equiaxial looking grains, and maintaining the fibrous characteristics of the fibres to a higher degree (Fig. 8). The hexagonal grains were 1—3 lm across the hexagonal plane and around 1 lm thick, and the fibres were decidedly more robust and handleable as a result, even as
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Fig. 7. XRD patterns of 0.67% CaO-doped Co Z fibres fired to 2 (a) 1100°C, (b) 1150°C, (c) 1175°C and (d) 1200°C for 3 h.
a thin web, although they were still not as mechanically strong as fibres fired to lower temperatures. By 1175°C, Co Z seemed to be the major of the 2 two components, and the presence of Co Y was not 2 indicated at all by the XRD pattern at 1200°C, indicating that Co Z was the sole phase at this 2 temperature (Fig. 7c and d), although again caution must be used as it is unlikely that pure single phase Co Z ferrite was actually obtained. Unfortu2 nately at this stage the exaggerated grain growth had accompanied the full transformation to Co Z, 2 resulting in the now familiar large, plate-like crystals seen also at this temperature in the undoped Co Z fibres (Fig. 9). Nevertheless, this is a very low 2 temperature for Co Z to form as the major com2 ponent, and the average crystallite size derived from the 100% peak at 2h"32.69° was 97 nm, less than half the size of the undoped Co Z fibres at 2 1250°C. An addition of calcium oxide and silica to spinel ferrites can cause an increase in the resistivity and a corresponding decrease of eddy current loss due to the additives segregating at grain boundaries [24] and forming an insulating layer upon cooling. An addition of CaO to manganese-zinc spinel ferrites has also been reported to decrease grain growth rate and sintering rate by the same segregation process [25]. It is suggested that the rate of growth in the plane perpendicular to the c-axis of the Co Z is being slowed by the segregation of 2
Fig. 8. SEM micrograph of 0.67% CaO-doped Co Z fibres 2 fired to 1150°C for 3 h.
Fig. 9. SEM micrograph of 0.67% CaO-doped Co Z fibres 2 fired to 1200°C for 3 h.
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calcium at the grain boundaries, allowing the growth along the c-axis to ‘catch-up’, thus preventing the exaggerated grain growth stage and resulting in the more equiaxed grains observed. The increase in temperature to 1200°C causes the dissolution of the additive in to the crystal lattice resulting in the onslaught of the exaggerated growth seen at this point. The aligned Co Y fibres were found to be identi2 cal in microstructure and composition to the non-aligned fibres reported previously [1] at the corresponding temperatures. The XRD data indicated that they formed a mixture of a-barium ferrite, cobalt ferrite and M ferrite between 600°C and 800°C, which then proceeded to yield a single phase XRD pattern for Co Y at 1000°C. The crys2 tallite size was calculated to be 56 nm from the Scherrer Equation using the 100% peak at 2h"41.19°. The fibres demonstrated themselves to be made of a densely packed mass of randomly oriented thin hexagonal plates between 1 and 3 lm in diameter and 0.1—0.5 lm in thickness, and were noticeably stronger and more handleable than any of the other fibres which had been fired to produce the Co Z or Co W phases. 2 2 The Co W fibres investigated also behaved like 2 those reported previously [2], forming a mixture of M ferrite and cobalt spinel ferrite at 1000°C which then produced at 1200°C what appeared to be single phase Co W from the XRD data. The forma2 tion of the Co W was also accompanied by the 2 exaggerated grain growth process, forming fibres consisting of lozenge-shaped hexagonal plates, some over 10 lm long, with a resulting loss of mechanical strength, as reported before. Despite this the calculated crystallite size was found to be 60 nm, much smaller than that for the Z ferrite fibres, using the 100% peak at 2h"34.54°.
magnetisation within the ab plane as well as the stiffness out of the ab plane. The corresponding anisotropy fields are H and H respectively, and ( h these are connected to the anisotropy constants K , 1 K (H ) and K (H ). The resonance condition is 2 h 3 ( given by 2pf "cJ(H H ) 3%4 h ( and literature values for H and H will therefore h ( give an indication of the FMR frequency expected for the Z, Y and W hexaferrites studied here. Typically H is much larger than H . For Co Z, h ( 2 H "13 000 Oe and H values have been reported h ( ranging from between 16 and 112 Oe [26]. These compute to f between 1.3 and 3.4 GHz. For 3%4 Co Y, H "28 000 Oe and H &150 Oe [27] re2 h ( sulting in f "5.7 GHz. H of Co W has been 3%4 h 2 reported as being 21 200 Oe [28], but no H values ( are given for these samples. The natural FMR frequency is, therefore, unclear. The measured spectra of k@ and kA for the Co Z 2 fibres fired to 1250°C are shown in Fig. 10. The calculated f was in excellent accord with our 3%4 measurements, and the values of both of k@ and kA at resonance agree closely with those previously reported for conventionally prepared specimens of Co Z [26,29], which for this ferrite all report a res2 onance frequency of approximately 1.3 GHz. This suggested that the sample was single phase Z ferrite, as was indicated by the XRD data discussed earlier.
3.4. Microwave measurements The natural ferromagnetic resonance (FMR) frequency is defined as the point at which kA is a maximum and which coincides with the half peak value of k@, where k@ is the real permeability and kA is the imaginary permeability. In materials with planar anisotropy the natural FMR frequency will be governed by the rotational stiffness of the
Fig. 10. FMR permeability components k@ (real permeability) and kA (imaginary permeability) for Co Z fibres fired to 1250°C. 2
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The spectra for the Ca-doped Co Z fibres fired to 2 1200°C are shown in Fig. 11. Whereas there was a peak in the kA spectrum at 1.3 GHz which suggested the presence of Co Z ferrite, the permeability 2 levels were much lower than for either standard pure Co Z ferrite samples or the Co Z fibres at 2 2 1250°C. It is possible that the reduction in the permeability at low GHz frequencies could be either an intrinsic effect due to the addition of calcium, or caused by the lower reaction-firing temperature used for this sample resulting in a mixedphase ferrite containing a small proportion of Co Y. 2 The Y ferrites are known to form at a lower temperature [13] than the W or Z ferrites, and this would seem consistent with the FMR spectra for the Co Y fibres fired to 1000°C, shown in Fig. 12. 2 A broad resonance in kA was seen, peaking between 6 and 7 GHz, and there was also a possible small resonance at 1.3 GHz, and these results appear to agree with those of bulk polycrystalline samples reported previously [30,31]. As mentioned above, there is no report of the natural FMR frequency of Co W ferrite. However, 2 the k@ and kA spectra of the Co W fibres fired to 2 1200°C (Fig. 13) were in good agreement with values reported in the frequency range 0.03—2 GHz [19]. This reference, which reports measurements made on isotropic polycrystalline samples, shows k@
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Fig. 12. FMR permeability components k@ (real permeability) and kA (imaginary permeability) for Co Y fibres fired to 2 1000°C.
Fig. 13. FMR permeability components k@ (real permeability) and kA (imaginary permeability) for Co W fibres fired to 2 1200°C.
Fig. 11. FMR permeability components k@ (real permeability) and kA (imaginary permeability) for 0.67% CaO-doped Co Z 2 fibres fired to 1200°C.
to be fairly constant at about 3 GHz and kA starting low and rising with increasing frequency, reaching a maximum at 1—2 GHz. There is no resonance reported, whereas a slight peak in kA was seen in the FMR spectra of the Co W fibres. This suggested 2 the minor presence of a Co Z phase, despite there 2 being no evidence of this in the XRD pattern [2], but this is consistent with previous work on the phases present and the difficulties encountered when attempting to produce pure W or Z ferrite phases [32].
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4. Conclusions Stable iron(III) sols were successfully produced with cobalt and barium doping, giving precursors which are stoichiometrically equivalent to Ba Co Fe O , Ca Ba Co Fe O , 3 2 24 41 0.3 3 2 24 41 Ba Co Fe O and BaCo Fe O . These sols 2 2 12 22 2 16 27 were all found to have a similar particle size to our previously reported ferroxplana-analogue doped precursor sols. Gel fibres of all the ferrites were successfully spun and collected as an aligned tow blanket. The alignment of these fibres was of a degree comparable to that found in commercially produced materials, which was maintained after subsequent heat treatment to form the ceramic ferrite product. The Co Z formulation gel fibres were found to 2 form pure Co Z ferrite at a temperature of 1250°C, 2 and the FMR spectra agreed that the material was indeed single phase, an achievement considering the complex range of similar phases which can co-exist at these temperatures. The seemingly unavoidable exaggerated platy grain growth process was still observed however, compromising the structural integrity of individual fibres. The addition of 0.67% CaO improved both the morphology and formation temperature of the Co Z fibres significantly. The Z phase was a major 2 component of the material at only 1150°C, and a much improved microstructure was observed with no exaggerated grain growth, resulting in smaller, more equiaxed grains even up to 1175°C, by which point Co Z was the dominant phase. It is 2 suggested that was is due to the calcium segregating at the boundary edges and slowing the rate of grain growth in the plane perpendicular to the c-axis. This allowed the growth to proceed in a slower and less two dimensional manner, until the dissolution of calcium into the lattice occurred and the runaway grain growth proceeded as usual. At 1200°C the exaggerated platy growth had occurred and Co Z appeared to be the only phase present by the 2 XRD pattern at this temperature, but FMR data indicated that the microwave properties were not as good as with the undoped Co Z fibres, which 2 could be either an intrinsic property of the calcium doped ferrite or due to the retention of some Co Y 2 as a minor phase.
The aligned Co Y and Co W fibres showed 2 2 identical compositions and morphologies to those reported earlier [1,2], with the Co Y fibres demon2 strating the better microstructure and hence mechanical properties at formation of the ferrite phase, and the Co W fibres demonstrating the exag2 gerated platy growth phase associated with the formation of Co W and Co Z ferrites. The FMR 2 2 spectra of Co Y and Co W fibres were also meas2 2 ured and both agreed with previous results published on polycrystalline forms, although there is very little published data to refer to regarding the high frequency properties of Co W ferrite. The 2 FMR spectra also indicated that the Co W fibres 2 may contain some Co Z phase as an impurity, 2 which is unresolved in the XRD spectra, but this is not unusual for samples of Co W and it is present 2 only as a very minor component.
Acknowledgements Our thanks to D. Croci for surface area measurements and R.C. Reynolds for the XPS and XRD characterisation (both at the Centre for Catalytic Systems and Materials Engineering, Department of Engineering, University of Warwick) and R. Burton for the XRF analysis (Materials Research Institute, Sheffield Hallam University). The data contained herein is published with the permission of the controller of Her Brittanic Majesty’s Stationery Office.
References [1] R.C. Pullar, M.D. Taylor, A.K. Bhattacharya, J. Mater. Sci. 32 (1997) 365. [2] R.C. Pullar, M.D. Taylor, A.K. Bhattacharya, J. Mater. Sci. 32 (1997) 873. [3] J.J. Went, G.W. Rathenau, E.W. Gorter, G.W. Van Oosterhaut, Phil. Tech. Rev. 13 (1952) 194. [4] H. Jonker, H.P. Wijn, P.B. Braun, Phil. Tech. Rev. 18 (1956) 145. [5] J. Smit, H.P.J. Wijn, Ferrites, Philips Technical Library, Eindhoven, 1959, pp. 204—207. [6] J. Smit, H.P.J. Wijn, Ferrites, Philips Technical Library, Eindhoven, 1959, p. 197. [7] D. Samaras, A. Collomb, S. Hadjivasilou, C. Achilleos, J. Tsoukalas, J. Pannetier, J. Rodriguez, J. Magn. Magn. Mater. 79 (1989) 193.
R.C. Pullar et al. / Journal of Magnetism and Magnetic Materials 186 (1998) 313—325 [8] A. Collomb, B. Lambert-Andron, J.X. Boucherle, D. Samaras, Phys. Stat. Sol. A 96 (1986) 385. [9] C. Heck, Magnetic Materials and their Applications, Butterworths, London, 1974, pp. 511—517. [10] I. Gordon, R.L. Harvey, R.A. Braden, J. Amer. Cer. Soc. 45 (1962) 297. [11] A.L. Stijts, H.P.J. Wijn, Phil. Tech. Rev. 19 (1958) 209. [12] H. Stablein, in: E.P. Wohlfarth (Ed.), Ferromagnetic Materials, vol. 3, North-Holland Physics Publishing, Amsterdam, 1982, pp. 462—535. [13] L.M. Catelliz, K.M. Kim, P.S. Boucher, J. Can. Ceram. Soc. 38 (1969) 57. [14] S. Ram, J.C. Joubert, J. Magn. Mag. Mater. 99 (1991) 133. [15] A.L. Stuijts, in: R.M. Fulrath, J.A. Pask (Eds.), Ceramic Microstructures, Wiley, New York, 1968, pp. 460—467. [16] M.J. Morton, J.D. Birchall, J.E. Cassidy, (ICI), UK Pat. 1360200 (1974). [17] M.H. Stacey, M.D. Taylor, (ICI), Eur. Pat. 318203 (1987). [18] A.M. Nicholson, G. Ross, IEEE Trans. Instrum. Meas. 19 (1970) 377. [19] E. Neckenburger, H. Severin, J.K. Vogel, G. Winkler, Z. Angew. Phys. 18 (1964) 65.
325
[20] M.A. Vinik, Russ. J. Inorg. Chem. 10 (1965) 1164. [21] S.I. Kuznetsova, E.P. Naiden, T.N. Stepanova, Inorg. Mat. 24 (1988) 856. [22] E.M.C. Huijser-Gerits, G.D. Rieck, D.L. Vogel, J. Appl. Crystallogr. 3 (1970) 243. [23] E.M.C. Huijser-Gerits, G.D. Rieck, J. Appl. Crystallogr. 9 (1976) 18. [24] T. Akashi, Japan J. Appl. Phys 30 (1961) 708. [25] M. Paulus, Phys. Stat. Sol. 2 (1962) 1181 and 1325. [26] J. Smit, H.P.J. Wijn, Ferrites, Philips Technical Library, Eindhoven, 1959, pp. 278—282. [27] J. Smit, H.P.J. Wijn, Ferrites, Philips Technical Library, Eindhoven, 1959, p. 206. [28] R.A. Braden, J. Gordon, R.L. Harvey, IEEE Trans. Magn. 2 (1966) 43. [29] M. Sugimoto, in: E.P. Wohlfarth (Ed.), Ferromagnetic Materials, vol. 3, North-Holland Physics Publishing, Amsterdam, 1982, pp. 428—433. [30] H.S. Belson, C.J. Kriessman, J. Appl. Phys. 18 (1959) 175S. [31] C.R. Buffler, J. Appl. Phys. 33 (1962) 1360. [32] M.A. Vinnik, A.I. Agranovskaya, N.N. Semenova, Russ. J. Inorg. Chem. 12 (1967) 18.