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Study of ferromagnetism in Co-doped rutile powders and float-zone grown single crystals
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Journal of Magnetism and Magnetic Materials 320 (2008) 887–894 www.elsevier.com/locate/jmmm
Study of ferromagnetism in Co-doped rutile powders and float-zone grown single crystals S.M. Koohpayeh, D. Fort, A.I. Bevan, A.J. Williams, J.S. Abell Department of Metallurgy and Materials, University of Birmingham, Edgbaston, Birmingham, B15 2TT, UK Received 17 July 2007 Available online 17 September 2007
Abstract Co-doped rutile samples in the form of both powders and bulk single crystals have been studied with particular emphasis on the dependence of their magnetic, compositional and structural properties upon the type of atmosphere used during their preparation. Both powders and single crystals were characterized using X-ray diffractometry and vibrating sample magnetometry, while the crystals were also studied using the X-ray Laue technique, scanning electron microscopy and energy dispersive X-ray analysis. The results indicate that an oxygen deficient environment during the preparation of Co-doped TiO2 powders is crucial for the observation of room temperature ferromagnetism, while preparation in oxygen rich conditions destroys the ferromagnetism due to the formation of the paramagnetic second phase CoTiO3. Floating zone growth of crystals under oxygen also led to the formation of material containing second phase CoTiO3 that was paramagnetic at room temperature, while crystals grown under argon were ferromagnetic and contained Co-rich inclusions. Crown Copyright r 2007 Published by Elsevier B.V. All rights reserved. Keywords: Floating zone crystal growth; Ferromagnetism; Co-doped rutile
1. Introduction In the quest to make electronic devices smaller, faster and cheaper, a promising new field of research is emerging in spintronics [1], which is a more powerful branch of electronics where the spin property of electrons, in addition to their charge, is utilized to achieve greater efficiency and functionality. In this context, dilute magnetic semiconductors (DMS), and in particular transition metal-doped oxide materials with a Curie point (Tc) at or above room temperature, have recently attracted a great deal of attention, leading to the discovery of room-temperature ferromagnetism in cobalt-doped TiO2 thin films by Matsumoto et al. [2]. Although stoichiometric TiO2 itself, in either the rutile or anatase form, is not a magnetic material, these thin films were reported to be ferromagnetic even above 400 K with a magnetic moment of 0.32 mB/Co (Bohr magnetons per cobalt atom). Corresponding author. Tel.: +44 121 414 5168; fax +44 121 414 5232.
E-mail address: [email protected] (J.S. Abell).
Since this initial discovery, Co-doped TiO2 has been studied extensively using a wide variety of thin film growth methods by several groups [3–24], but ongoing controversy concerns the mechanism giving rise to ferromagnetism; in particular the question of whether the cobalt atoms really incorporate into the TiO2 structure (either rutile or anatase), by substituting for Ti atoms, or form Co clusters or inclusions has yet to be answered conclusively. In this regard, some studies support the theory that cobalt is well incorporated (not clustered) in Co-doped anatase [3–6], while other work, such as Kim et al. [7], claims that Co clusters in the TiO2 structure are responsible for the high temperature magnetic behaviour in Co-doped anatase thin films. The formation of such Co clusters or inclusions, which would indicate that any ferromagnetism is due to a magnetic secondary phase, rather than being an intrinsic property of the host material, has also been supported in other reports [8–13]. A possible factor influencing the Co distribution in Codoped TiO2 could be the presence of oxygen vacancies, as indicated by Kim et al. [12], who investigated the influence of the oxygen partial pressure ðPO2 Þ during preparation on
0304-8853/$ - see front matter Crown Copyright r 2007 Published by Elsevier B.V. All rights reserved. doi:10.1016/j.jmmm.2007.09.006
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the Co distribution in epitaxial anatase films grown on SrTiO3 substrates using pulsed laser deposition (PLD). Using high resolution XTEM, they found an increasing tendency of Co to cluster, together with an increasing magnetic moment per Co atom, with decreasing PO2 ; these results were also supported by Murakami [14] and Stampe [11]. As well as enhancing the ferromagnetism [15–16], the presence of oxygen vacancies is further reported to affect both the electronic and optical properties, as reported by Weng [16]. An interesting recent development has been the claim by Yoon et al. [25] that ferromagnetism at temperatures up to 880 K can be induced in TiO2d films grown by PLD even without the introduction of magnetic ions (i.e. undoped). The origin of the magnetism in these TiO2d films (for which no value for d was given) was attributed to the presence of anion defects (i.e. on oxygen sites) introduced via the substrate mismatch and by processing in an oxygen deficient atmosphere (similar behaviour in HfO2 thin films deposited by PLD has been attributed to defect doping [26]). The reported room temperature magnetic moments in Codoped TiO2 thin films range from 0.1 mB/Co [7] to values as high as 2.0 mB/Co [13]. These somewhat contradictory results suggest that the magnetic properties of Co-doped TiO2 films could be dependent upon the method and conditions of sample preparation [17]. Although the origin of the ferromagnetism observed in Co-doped TiO2 films is not clear at present, if the magnetism is intrinsic in nature, one may expect to see similar behaviour in bulk samples, where any effects on the microstructure and magnetic interactions due to lattice mismatch between substrate and thin film should be eliminated. Preparation of bulk CoxTi1xO2 samples has been performed by a few research groups, e.g. polycrystalline powders have been prepared by solid state reaction [27] or sol–gel [28–29] methods, while single crystalline rutile substrates have been implanted with cobalt [30]. However, as is the case with thin films, not only does the main mechanism for the observed ferromagnetism remain unresolved for bulk samples, but the magnitude of reported room temperature magnetic moments are also inconsistent [27–28]. The main aim of the present study, therefore, was to prepare and characterize powders and, for the first time, bulk single crystals of Co-doped rutile (grown by the floating zone (FZ) technique) to discover what effect varying preparation conditions (e.g. using oxidizing or inert environments) have upon structural and magnetic properties. It should be noted that controlled solidification from the melt represents a near equilibrium growth process, in contrast to many thin film growth mechanisms which are non-equilibrium in nature. 2. Experimental procedures 2.1. Powder and feed rod preparation Two starting material combinations were used to make Co-doped TiO2 powders and feed rods (for FZ
crystal growth), namely mixes of CoO+TiO2 powders and Co+TiO2 powders (in both cases the TiO2 was in the form of anatase). The starting powders were mixed in the desired ratio and the Co content was varied from 2 to 8 atomic % with respect to titanium. The CoO+TiO2 mixtures were calcined at 1100 1C for 8 h under oxygen, while the Co+TiO2 mixes were calcined under vacuum at 106 mbar or oxygen for 8 h at 1100 1C. After grinding and phase identification using XRD, calcined powders were used for both magnetic measurements and feed rod preparation. To make starting rods for FZ crystal growth, powder was packed and subsequently sealed into a rubber tube while evacuating air using a vacuum pump. The powder was then compacted into a rod, typically 6 mm in diameter and 70 mm long, using a hydraulic press under an isostatic pressure of 700 bar. After removal from the rubber tube, the pressed rods were sintered at 1100 1C for 8 h in oxygen for the powders calcined in oxygen, or in vacuum for powders calcined in vacuum. (Higher sintering temperatures were found not appropriate due to the evaporation of phases including cobalt which led to departures from the stoichiometric composition). 2.2. Floating zone crystal growth Single crystals were grown from the sintered feed rods using a four-mirror optical FZ furnace (Crystal System Inc. FZ–T–10000–H–VI–VP) having four 1 kW halogen lamps as the heating source. Pure rutile single crystals, oriented with the ‘c’ axis along their lengths, were used as seed crystals for all growths. The molten zone was moved upwards during all float zoning runs with the seed crystal being at the bottom and the feed rod above it. Growths were performed at 5 mm/h with rotation rates of 15 rpm for the growing crystal (lower shaft) and 0 rpm for the feed rod (upper shaft) in either an oxygen or argon environment. 2.3. Characterization Phase identification and the determination of crystal structure of all samples (both synthesized powders and single crystals) were performed using X-ray powder diffraction techniques, while X-ray Laue diffraction was utilized to check the crystalline quality and orientations of the crystals prepared. The compositional homogeneities of as-grown single crystals were analyzed by energy dispersive X-ray (EDX) technique using scanning electron microscopy (SEM). Magnetization measurements were carried out using a Lake Shore Cryotronics, Inc. vibrating sample magnetometer (VSM) on both calcined powders and oriented single crystals (approximately 3 3 1.5 mm along the ‘a’ and ‘c’ directions, respectively) which were cut from the as-grown crystals.
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3. Results and discussion 3.1. Co-doped TiO2 powders One apparent difference between powders after calcining was that powders synthesized in oxygen were greenish in colour, while powders synthesized in vacuum were greyer. Phase identification using X-ray diffraction indicated that the combination of CoO+TiO2 (anatase) synthesized in oxygen was converted during calcining mainly to the rutile structure with some weak second phase reflections consistent with CoTiO3, the intensities of which increased with increasing Co content from 2% to 8%. This is illustrated in Fig. 1a, which shows a typical XRD pattern taken from the powder with a Co content of 4%. These results are consistent with the phase relations in the TiO2–CoO pseudo-binary system [31]. Starting material combinations of Co+TiO2 (anatase) powder (%Co ¼ 2–8) were also calcined in oxygen at 1100 1C for 8 h and XRD from the resulting powders indicated the rutile structure with extra peaks consistent with CoTiO3, i.e. similar to powders prepared using CoO:TiO2 mixtures synthesized in oxygen. However, XRD patterns taken from powders of Co+TiO2 (anatase) mixture synthesized in vacuum showed only reflections for the rutile structure (Fig. 1b) with no clear evidence for any second phases (e.g. Co, CoO, Co3O4, CoTiO3 or CoTi2O5), although the presence of some small amounts of second phase, below the detection limit of the XRD technique, is, of course, possible. 3.2. Magnetic properties of Co-doped rutile powders The magnetic properties of the as-synthesized powders described in Section 3.1 were measured at room temperature. Magnetization measurements performed on initially
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CoO:TiO2 mixtures synthesized in oxygen (which X-ray diffractometry indicated as having the rutile structure with some second phase CoTiO3) showed that the powders were paramagnetic (Figs. 2a–c). As is evident from this figure, increasing the Co content from 2% to 8% (which should increase the amount of second phase CoTiO3 in the powders) increased the magnetization values. This is consistent with the fact that CoTiO3 is a paramagnetic material at room temperature (CoTiO3 is antiferromagnetic at low temperature with a Neel temperature of 35.6 K), while rutile is a diamagnet. In contrast, magnetization vs. field plots measured from initially Co:TiO2 mixtures synthesized in vacuum (which led to the formation of the rutile structure with no obvious evidence of other phases on the XRD patterns shown in Fig. 1b) showed room temperature ferromagnetism for all powders (Figs. 2d–f) with magnetization values two orders of magnitude higher than those measured from powders synthesized in oxygen. While increasing the Co content from 2% (Fig. 2d) to 8% (Fig. 2f) increased the magnetic moment per kilogram for these Co:TiO2 mixtures synthesized in vacuum, no significant difference was observed in the calculated magnetization values per Co atom for Co contents between 2% and 8%, which were all in the range 1.8–1.9 mB/Co atom. Since these magnetic moment values per Co atom for vacuum-prepared powders are close to those of pure cobalt (1.7 mB), the obvious conclusion is that the enhanced magnetization in these powders could be attributed to the presence of Co particles which were too small to be detected by the X-ray diffraction technique. This observation of room temperature ferromagnetism in Co-doped TiO2 powders prepared in a reduced oxygen environment is consistent with the work by Cho et al. [28], where the presence of Co metal clusters was observed in
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Fig. 1. XRD patterns taken from (a) CoO+TiO2 (anatase) powder synthesized in oxygen and (b) Co+TiO2 (anatase) powder synthesized in vacuum. Both powders were synthesized at 1100 1C for 8 h, and the Co content was 4%.
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Fig. 2. Magnetization measurements performed on CoO:TiO2 powders synthesized in oxygen for Co contents of (a) 2%, (b) 4% and (c) 8%, and Co:TiO2 powders synthesized in vacuum with (d) 2%Co, (e) 4%Co and (f) 8%Co.
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vacuum-annealed powders (synthesized by the sol-gel method) using a high resolution TEM. The topic of the exact origin of the ferromagnetism seen in vacuum prepared powders is returned to in a later section on single crystals. Intensity (arb. units)
Crystal growths were only attempted with Co contents of 4% and 8%, rather than 2%, since detection of the formation of any second phases would be more reliable at higher cobalt contents. Growths using feed rods sintered in oxygen with a Co content of 8% were not successful, however, due to the formation of a mushy (liquid+solid) region below the molten zone, which led to the zone falling out. By reducing the Co content to 4%, though, it proved possible to perform a successful growth due to the formation of a more solid region below the melt zone and, consequently, a more stable molten zone. This more solid region below the molten zone at lower Co contents is consistent with the reported phase relations [31] which predict that solidification of a 4% Co composition should start with a lesser amount of liquid compared to 8% Co. Growth using feed rods sintered in vacuum also exhibited a stable molten zone for the 4%Co composition. A gas pressure of 5 bar (either oxygen using the feed rod sintered in oxygen, or argon using the feed rod sintered in vacuum) was applied during the crystal growths. These relatively high pressures were intended to both sharpen the temperature gradient along the growing crystal [32], thereby attaining a more solid region below the molten zone, and to limit evaporation of any volatile phases (as had occurred during sintering at 1600 1C under ambient pressure). In both cases, crystals of approximately 5 mm in diameter and 30 mm in length were grown, which initially appeared to be black in colour, although cutting and polishing revealed that while the crystal grown in argon was black, the crystal grown in oxygen was actually dark red. X-ray Laue diffraction photographs taken along the lengths and cross-sections of the crystals indicated that both exhibited an equally high crystalline quality and that the crystals had adopted the orientation of the seed crystal (i.e. with the ‘c’ direction along their length). X-ray diffraction patterns taken from sections of the asgrown crystals after crushing indicated that the crystal grown in oxygen had the rutile structure with second phase CoTiO3 (Fig. 3a) similar to the starting powder. However, the crystal grown in argon showed the rutile structure with very weak and broad extra peaks consistent with the 002 and 101 hcp reflections for pure Co at 2y values of 44.71 and 47.51, respectively (Fig. 3b). This contrasts with the starting powder for this crystal in which no extra reflections beyond those expected for the rutile structure were evident. The broadening seen in these extra peaks could indicate strain and/or that they were not from pure Co, but contained some Ti or O.
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Fig. 3. The XRD patterns taken from the crushed Co-doped rutile (%Co ¼ 4) single crystals grown under 5 bar pressure of (a) oxygen, and (b) argon. The crystal grown in oxygen had the rutile structure with second phase CoTiO3, while the crystal grown in argon showed the rutile structure with very weak and broad extra peaks consistent for pure Co.
For further characterization, back-scattered SEM micrographs were taken from polished cross and longitudinal sections of the crystal grown in oxygen (Figs. 4a–b) and argon (Figs. 4c–d). The formation of second phases in both crystals is clearly shown in these back-scattered images, although the inclusions are differently shaped. EDX chemical mapping images taken from the crystal grown in oxygen (Fig. 5a) revealed that Ti and O were uniformly distributed throughout the whole area, while the secondary phase particles also contain Co. Conversely, the EDX maps taken from the crystal grown in argon (Fig. 5b) showed a significant drop in Ti and O signal intensities in the precipitates indicating that the secondary phase is either Co or very rich in Co. While the EDX technique is not particularly suited to examining light elements such as oxygen, analyses
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Fig. 4. SEM back scattered micrographs of cross and longitudinal sections cut from 4%Co:rutile single crystals grown in oxygen (a and b) and argon (c and d). Growth in oxygen led to the formation of CoTiO3 second phase, while the crystal grown in argon had Co rich particles elongated along the growth direction (‘c’).
performed on both crystal types (i.e. grown in oxygen and argon) indicated that the matrix was very close to TiO2 with no clear evidence of Co incorporation, while all second phases consistently included Co. Although the technique was found not to be sufficiently accurate to identify the exact Co–Ti–O second phase in the sample grown in oxygen, results are consistent with the XRD data that indicated the second phase to be CoTiO3. For the crystal grown in argon, the EDX results indicated that second phases were Co rich with the average atomic percentage of more than 90% (with the atomic percentage of Ti varying between 3.72% and 10%), which again agrees with the XRD data from the as-crushed crystal that showed broad peaks consistent with hcp Co. Since these Co-rich inclusions consistently have a square shape in cross section and are elongated along the ‘c’ direction of the matrix (growth direction) it can be surmised that they are either single crystal or polycrystalline with a preferential crystalline orientation. Concerning the ongoing controversy of whether Co atoms actually incorporate into the TiO2 structure or form clusters or inclusions, it would appear likely from the evidence of the above results for bulk crystals, that if significant Co incorporation does occur, it is a nonequilibrium effect peculiar to thin films. 3.4. Magnetic properties of Co-doped rutile single crystals The magnetic response from the 4%Co:TiO2 single crystal samples was measured as a function of magnetic
field along the ‘c’, ‘a’ and 451 from ‘c’ towards ‘a’ directions for the sample grown in argon, and along ‘c’ and ‘a’ for the sample grown in oxygen, as shown in Fig. 6. These measurements clearly indicate that room temperature ferromagnetism was observed only in the crystal grown in argon, while crystals grown in oxygen showed paramagnetic behaviour. The room temperature paramagnetic behaviour of crystals grown in oxygen can be explained by the formation of paramagnetic second phase CoTiO3, as described earlier for the powder samples. This weakly ferromagnetic argon-grown crystal had a saturation magnetization of about 0.24 A m2 kg1, with the easy magnetization direction along ‘c’ and the hard direction of magnetization along ‘a’. This magnetization value equates to 0.083 mB/Co atom if the crystal composition is taken to be the same as the starting powder (i.e. 4%Co:TiO2), although there are some doubts as to whether this is the case (see below). A single FZ experiment under argon was also performed with a CoO:TiO2 (%Co ¼ 4) mixture initially calcined in oxygen, which gave a very weakly ferromagnetic sample with a saturation magnetization about 0.05 A m2 kg1 along the ‘c’ direction. The various routes used to prepare the crystals from starting mixes to final FZ growth are summarized in Fig. 7. The differences of the magnetic properties of the two main crystal types, as shown in Fig. 6, can be explained from the results of the XRD and SEM microstructural studies. In this regard, the room temperature ferromagnetism observed in crystals grown in argon is almost certainly due to the existence of ferromagnetic Co-rich clusters.
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Fig. 6. M–H curves measured at room temperature for 4%Co:rutile single crystals. Room temperature ferromagnetism with magnetization 0.24 A m2 kg1 (0.083 mB/Co atom) was observed in the crystal grown in argon, while a weak paramagnetic behaviour was obtained from the crystal grown in oxygen.
Fig. 5. EDX chemical mapping images of cross-sections cut from 4%Co:TiO2 single crystals grown in (a) oxygen and (b) argon.
However, one interesting feature is that the saturation magnetization of these crystals at 0.24 A m2 kg1 is significantly lower than the magnetization of the feed rod/starting material (i.e. the 4%Co:TiO2 powders prepared in vacuum), which had values of about 5.61 A m2 kg1. As noted previously, this latter value equates to 1.9 mB/Co atom, which is close to the magnetic moment of pure cobalt, 1.7 mB/Co atom. To ensure that the form of the magnetization samples (powder for the starting material against bulk single crystal) was not a factor in these differences, a further magnetization measurement was performed on the single crystal grown in argon after powdering, but the measured saturation magnetization value (0.204 A m2 kg1) was not significantly different to the value obtained from the crystal before powdering (0.24 A m2 kg1). Explanations for the significant reduction in magnetization values for 4%Co:TiO2 after FZ crystal growth include that it resulted from a reduced amount of Co in the final
crystal compared to the start material and/or a reduction in the magnetization of the Co atoms present. A reduction in Co content in the as-grown crystals could occur as a result of evaporation of Co during FZ, or from segregation of Co during the zoning process. Although no overt evidence for Co volatilization losses (e.g. condensation of Co onto the furnace tube) were detected during the crystal growth, some loss of Co is likely; losses of about 1.5–2.0%Co were reported [33] during melting and crystal growth of CoTi crystals, for example, which has a significantly lower melting point than TiO2. Segregation of components during FZ is a common occurrence, especially when one component has a low solubility in the matrix, as Co appears to have in TiO2. Thus during FZ of Co-doped TiO2, it is not unrealistic to presume that Co would be ‘‘pushed’’ along by the molten zone leaving a reduced amount of Co in the as-grown crystal. Some strength to this argument was given by a magnetization measurement taken from the final part of the rod traversed by the molten zone during the crystal growth process (i.e. the section directly above the as-grown crystal) which had magnetization values an order of magnitude higher than the crystal at 2.3 A m2 kg1. Besides a reduction of the Co content as a result of the FZ process, a lowering of magnetization values of crystals may also result from changes in the magnetic/electronic structure of the Co atoms themselves. In this regard, several studies have indicated that Co between Ti layers (as the local environment) has very small magnetic moments (down to 0.07 mB [34]), while in Co100xTix thin films, increasing the Ti concentration significantly decreased the magnetic moment of Co [35–37]. Although the X-ray data from the crushed crystal indicated that the Co-rich inclusions still retained a hexagonal crystal structure, the solubility limit for Ti dissolved in Co has a
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Starting mixtures CoO+TiO2
Calcining atmosphere
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Argon
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Fig. 7. Preparation routes and magnetic properties of Co-doped rutile powders and single crystals.
maximum of 15% at 1190 1C. Overall, it is thought probable that a combination of circumstances (i.e. some reduction in Co content together with a lowering of the magnetization of Co atoms due to their local environment) explains these results, although further work is needed to discover which is the dominant effect. The weak dependence of the magnetic properties of the single crystal on a preferred direction (i.e. easy and hard magnetization directions along ‘c’ and ‘a’ directions, respectively) shows that this material exhibits some magnetic anisotropy. Since the matrix, rutile, is a diamagnetic material, this apparent anisotropy must come from the Co-rich second phases and probably arises either from the shape anisotropy of the individual precipitates (due to magnetostatic effects, i.e. the demagnetizing field was less when measuring magnetization along the long axis of these needle shaped inclusions) and/or from the intrinsic magnetocrystalline anisotropy of Co (if these second phase inclusions are single crystals or preferentially aligned). However, since these inclusions are small in size, shape anisotropy effects could be more dominant than magnetocrystalline anisotropy. To end our discussion, it is worth commenting upon the very recent work on the magnetic properties of undoped TiO2d films grown by PLD [25] which indicated that these oxygen deficient anatase films show ferromagnetism at temperatures up to 880 K even without the introduction of magnetic ions. This magnetism was attributed to the presence of defects on oxygen sites (as a result of interfacial defects created at the film-substrate interface and from processing in an oxygen deficient atmosphere during PLD growth) leading to the creation of Ti3+ and Ti2+ ions (with unpaired 3d electrons) by stripping electrons from the s band. Thus questions are now being asked [41] as to what degree transition metal doping plays an essential role in producing ferromagnetism in thin films of non-magnetic oxides. Earlier work in the authors’ laboratory on FZ crystal growth of TiO2 has shown that oxygen deficient (i.e. blue colour) bulk rutile single crystals can be grown in the image furnace by using a higher than normal molten zone temperature [38]. When the magnetic properties of these crystals were measured, however, they gave no indication of ferromagnetism at room temperature. While this work on bulk crystals cannot be directly compared with the thin film work of Yoon et al. (since the TiO2 crystal was in the
form of rutile while the films were anatase), it would appear to confirm that oxygen deficiency by itself does not significantly affect the magnetic properties of all oxygen deficient TiO2-based crystals. However, it should be noted that the Ti–O phase diagram shows several TiO2d phases [39], which contain various types of defects such as Tin+ interstitials, oxygen vacancies and planar defects (e.g. crystallographic shear planes) although the question of which type of defect is dominant in which region of oxygen deficiency, is still subject to debate [40]. It is not inconceivable, therefore, that in bulk crystals one of these oxygen deficient phases could exhibit enhanced magnetic properties even without the strain caused by lattice mismatch in Yoon et al.’s thin film work. 4. Conclusions It can be concluded that an oxygen deficient environment during the preparation of Co-doped TiO2 powders using a solid-state reaction was crucial for the observation of room temperature ferromagnetism. Preparation in oxygen-rich conditions destroyed the ferromagnetism and led to the formation of the paramagnetic second phase CoTiO3. Crystal growth of Co-doped rutile using the FZ melting technique led to the formation of cobalt-based second phases within crystals; crystals grown in oxygen were paramagnetic at room temperature and contained second phase CoTiO3, while crystals grown in argon were ferromagnetic at room temperature and contained Co-rich second phases. No evidence of significant solution of Co into the rutile matrix was found in either crystal type. Acknowledgement This work was supported by the UK Engineering and Physical Sciences Research Council. References [1] S.A. Wolf, D.D. Awschalom, R.A. Buhrman, J.M. Daughton, S.V. Molnar, M.L. Roukes, A.Y. Chtchelkanova, D.M. Treger, Science 294 (2001) 1488. [2] Y. Matsumoto, M. Murakami, T. Shono, T. Hasegawa, T. Fukumura, M. Kawasaki, P. Ahmet, T. Chikyow, S.Y. Koshihara, H. Koinuma, Science 291 (2001) 854.
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[3] Y. Matsumoto, R. Takahashi, M. Murakami, T. Koida, X.J. Fan, T. Hasegawa, T. Fukumura, M. Kawasaki, S.Y. Koshihara, H. Koinuma, Jpn. J. Appl. Phys. 40 (2001) L1204. [4] L.F. Liu, J.F. Kang, Y. Wang, H. Tang, L.G. Kong, X. Zhang, R.Q. Han, Solid State Commun. 139 (2006) 263. [5] K.A. Griffin, A.B. Pakhomov, C.M. Wang, S.M. Heald, K.M. Krishnan, Phys. Rev. Lett. 94 (2005) 157204. [6] W.K. Park, R.J. Ortega-Hertogs, J.S. Moodera, A. Punnoose, M.S. Seehra, J. Appl. Phys. 91 (2002) 8093. [7] J.Y. Kim, J.H. Park, B.G. Park, H.J. Noh, S.J. Oh, J.S. Yang, D.H. Kim, S.D. Bu, T.W. Noh, H.J. Lin, H.H. Hsieh, C.T. Chen, Phys. Rev. Lett. 90 (2003) 017401. [8] S.R. Shinde, S.B. Ogale, S.D. Sarma, J.R. Simpson, H.D. Drew, S.E. Lofland, C. Lanci, J.P. Buban, N.D. Browning, V.N. Kulkarni, J. Higgins, R.P. Sharma, R.L. Greene, T. Venkatesan, Phys. Rev. B 67 (2003) 115211. [9] G.C. Han, Y.H. Wu, M. Tay, Z.B. Guo, K.B. Li, C.T. Chong, J. Magn. Magn. Mater. 272–276 (2004) e1537. [10] B.S. Jeong, Y.W. Heo, D.P. Norton, J.G. Kelly, R. Rairigh, A.F. Hebard, J.D. Budai, Y.D. Park, Appl. Phys. Lett. 84 (2004) 2608. [11] P.A. Stampe, R.J. Kennedy, Y. Xin, J.S. Parker, J. Appl. Phys. 92 (2002) 7114. [12] D.H. Kim, J.S. Yang, K.W. Lee, S.D. Bu, T.W. Noh, S.J. Oh, Y.W. Kim, J.S. Chung, H. Tanaka, H.Y. Lee, T. Kawai, Appl. Phys. Lett. 81 (2002) 2421. [13] G.C. Han, P. Luo, Z.B. Guo, F.U. Nahar, M. Tay, Y.H. Wu, S.J. Wang, Thin Solid Films 505 (2006) 137. [14] M. Murakami, Y. Matsumoto, M. Nagano, T. Hasegawa, M. Kawasaki, H. Koinuma, Appl. Surf. Sci. 223 (2004) 245. [15] S.A. Chambers, S. Thevuthasan, R.F.C. Farrow, R.F. Marks, J.U. Thiele, L. Folks, M.G. Samant, A.J. Kellock, N. Ruzycki, D.L. Ederer, U. Diebold, Appl. Phys. Lett. 79 (2001) 3467. [16] H. Weng, X. Yang, J. Dong, H. Mizuseki, M. Kawasaki, Y. Kawazoe, Phys. Rev. B 69 (2004) 125219. [17] R. Janisch, P. Gopal, N.A. Spaldin, J. Phys. Condens. Matter 17 (2005) R657. [18] N. Popovici, E. Jimenez, R.C. da Silva, W.R. Branford, L.F. Cohen, O. Conde, J. Non-Cryst. Solids 352 (2006) 1486. [19] L.F. Liu, J.F. Kang, Y. Wang, H. Tang, L.G. Kong, L. Sun, X. Zhang, R.Q. Han, J. Magn. Magn. Mater. 308 (2007) 85. [20] N.H. Hong, J. Sakai, W. Prellier, A. Ruyter, J. Appl. Phys. 38 (2005) 816.
[21] M. Murakami, Y. Matsumoto, T. Hasegawa, P. Ahmet, K. Nakajima, T. Chikyow, H. Ofuchi, I. Nakai, H. Koinuma, J. Appl. Phys. 95 (2004) 5330. [22] Y.H. Jeong, S.J. Han, J.H. Park, Y.H. Lee, J. Magn. Magn. Mater. 272–276 (2004) 1976. [23] G.C. Han, Y.H. Wu, M. Tay, K.B. Li, Z.B. Guo, T.C. Chong, J. Magn. Magn. Mater. 268 (2004) 159. [24] S.A. Chambers, C.M. Wang, S. Thevuthasan, T. Droubay, D.E. McCready, A.S. Lea, V. Shutthanandan, C.F. Windisch Jr., Thin Solid Films 418 (2002) 197. [25] S.D. Yoon, Y. Chen, A. Yang, T.L. Goodrich, X. Zuo, D.A. Arena, K. Ziemer, C. Vittoria, V.G. Harris, J. Phys. Condens. Matter 18 (2006) L355. [26] M. Venkatesan, C.B. Fitzgerald, J.M.D. Coey, Nature 430 (2004) 630. [27] Y.G. Joh, H.D. Kim, B.Y. Kim, S.I. Woo, S.H. Moon, J.H. Cho, E.C. Kim, D.H. Kim, J. Korean Phys. Soc. 44 (2004) 360. [28] J.H. Cho, B.Y. Kim, H.D. Kim, S.I. Woo, S.H. Moon, J.P. Kim, C.R. Cho, Y.G. Joh, E.C. Kim, D.H. Kim, Phys. Stat. Sol. 241 (2004) 1537. [29] G.S. Hui, W.X. Wei, K.X. Ming, Z.X. Yun, X. Li, Z.Y. Lu, Y.X. Lin, Z.Y. Xuan, Chin. Phys. Lett. 22 (2005) 1772. [30] N. Akdogan, B.Z. Rameev, L. Dorosinsky, H. Sozeri, R.I. Khaibullin, B. Aktas, L.R. Tagirov, A. Westphalen, H. Zabel, J. Phys. Condens. Matter 17 (2005) L359. [31] B. Brezny, A. Muan, J. Inorg. Nucl. Chem. 31 (1969) 649. [32] S.M. Koohpayeh, Ph.D. Thesis, University of Birmingham, UK, July 2007. [33] L. Zhang, M.L. Jenkins, G. Taylor, J. Cryst. Growth 289 (2006) 308. [34] A. Jezierski, Phys. Stat. Sol. 3 (2006) 225. [35] G. Suran, J. Sztern, J.A. Aboaf, T.R. McGuire, IEEE Trans. Magn. Mag-17 (1981) 3065. [36] K.T. Wu, Y.D. Yao, W.C. Chang, C.Y. Chu, J. Magn. Magn. Mater. 140–144 (1995) 687. [37] K. Qunadjela, G. Suran, F. Machizaud, Phys. Rev. B 40 (1989) 578. [38] S.M. Koohpayeh, D. Fort, J.S. Abell, J. Cryst. Growth 282 (2005) 190. [39] G.V. Samsonov, The Oxide Handbook, IFI/Plenum Press, New York, 1982. [40] U. Diebold, Surf. Sci. Rep. 48 (2003) 53. [41] N.H. Hong, J. Sakai, F. Gervais, J. Magn. Magn. Mater. 316 (2007) 214.
Journal of Magnetism and Magnetic Materials 320 (2008) 887–894 www.elsevier.com/locate/jmmm
Study of ferromagnetism in Co-doped rutile powders and float-zone grown single crystals S.M. Koohpayeh, D. Fort, A.I. Bevan, A.J. Williams, J.S. Abell Department of Metallurgy and Materials, University of Birmingham, Edgbaston, Birmingham, B15 2TT, UK Received 17 July 2007 Available online 17 September 2007
Abstract Co-doped rutile samples in the form of both powders and bulk single crystals have been studied with particular emphasis on the dependence of their magnetic, compositional and structural properties upon the type of atmosphere used during their preparation. Both powders and single crystals were characterized using X-ray diffractometry and vibrating sample magnetometry, while the crystals were also studied using the X-ray Laue technique, scanning electron microscopy and energy dispersive X-ray analysis. The results indicate that an oxygen deficient environment during the preparation of Co-doped TiO2 powders is crucial for the observation of room temperature ferromagnetism, while preparation in oxygen rich conditions destroys the ferromagnetism due to the formation of the paramagnetic second phase CoTiO3. Floating zone growth of crystals under oxygen also led to the formation of material containing second phase CoTiO3 that was paramagnetic at room temperature, while crystals grown under argon were ferromagnetic and contained Co-rich inclusions. Crown Copyright r 2007 Published by Elsevier B.V. All rights reserved. Keywords: Floating zone crystal growth; Ferromagnetism; Co-doped rutile
1. Introduction In the quest to make electronic devices smaller, faster and cheaper, a promising new field of research is emerging in spintronics [1], which is a more powerful branch of electronics where the spin property of electrons, in addition to their charge, is utilized to achieve greater efficiency and functionality. In this context, dilute magnetic semiconductors (DMS), and in particular transition metal-doped oxide materials with a Curie point (Tc) at or above room temperature, have recently attracted a great deal of attention, leading to the discovery of room-temperature ferromagnetism in cobalt-doped TiO2 thin films by Matsumoto et al. [2]. Although stoichiometric TiO2 itself, in either the rutile or anatase form, is not a magnetic material, these thin films were reported to be ferromagnetic even above 400 K with a magnetic moment of 0.32 mB/Co (Bohr magnetons per cobalt atom). Corresponding author. Tel.: +44 121 414 5168; fax +44 121 414 5232.
E-mail address: [email protected] (J.S. Abell).
Since this initial discovery, Co-doped TiO2 has been studied extensively using a wide variety of thin film growth methods by several groups [3–24], but ongoing controversy concerns the mechanism giving rise to ferromagnetism; in particular the question of whether the cobalt atoms really incorporate into the TiO2 structure (either rutile or anatase), by substituting for Ti atoms, or form Co clusters or inclusions has yet to be answered conclusively. In this regard, some studies support the theory that cobalt is well incorporated (not clustered) in Co-doped anatase [3–6], while other work, such as Kim et al. [7], claims that Co clusters in the TiO2 structure are responsible for the high temperature magnetic behaviour in Co-doped anatase thin films. The formation of such Co clusters or inclusions, which would indicate that any ferromagnetism is due to a magnetic secondary phase, rather than being an intrinsic property of the host material, has also been supported in other reports [8–13]. A possible factor influencing the Co distribution in Codoped TiO2 could be the presence of oxygen vacancies, as indicated by Kim et al. [12], who investigated the influence of the oxygen partial pressure ðPO2 Þ during preparation on
0304-8853/$ - see front matter Crown Copyright r 2007 Published by Elsevier B.V. All rights reserved. doi:10.1016/j.jmmm.2007.09.006
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the Co distribution in epitaxial anatase films grown on SrTiO3 substrates using pulsed laser deposition (PLD). Using high resolution XTEM, they found an increasing tendency of Co to cluster, together with an increasing magnetic moment per Co atom, with decreasing PO2 ; these results were also supported by Murakami [14] and Stampe [11]. As well as enhancing the ferromagnetism [15–16], the presence of oxygen vacancies is further reported to affect both the electronic and optical properties, as reported by Weng [16]. An interesting recent development has been the claim by Yoon et al. [25] that ferromagnetism at temperatures up to 880 K can be induced in TiO2d films grown by PLD even without the introduction of magnetic ions (i.e. undoped). The origin of the magnetism in these TiO2d films (for which no value for d was given) was attributed to the presence of anion defects (i.e. on oxygen sites) introduced via the substrate mismatch and by processing in an oxygen deficient atmosphere (similar behaviour in HfO2 thin films deposited by PLD has been attributed to defect doping [26]). The reported room temperature magnetic moments in Codoped TiO2 thin films range from 0.1 mB/Co [7] to values as high as 2.0 mB/Co [13]. These somewhat contradictory results suggest that the magnetic properties of Co-doped TiO2 films could be dependent upon the method and conditions of sample preparation [17]. Although the origin of the ferromagnetism observed in Co-doped TiO2 films is not clear at present, if the magnetism is intrinsic in nature, one may expect to see similar behaviour in bulk samples, where any effects on the microstructure and magnetic interactions due to lattice mismatch between substrate and thin film should be eliminated. Preparation of bulk CoxTi1xO2 samples has been performed by a few research groups, e.g. polycrystalline powders have been prepared by solid state reaction [27] or sol–gel [28–29] methods, while single crystalline rutile substrates have been implanted with cobalt [30]. However, as is the case with thin films, not only does the main mechanism for the observed ferromagnetism remain unresolved for bulk samples, but the magnitude of reported room temperature magnetic moments are also inconsistent [27–28]. The main aim of the present study, therefore, was to prepare and characterize powders and, for the first time, bulk single crystals of Co-doped rutile (grown by the floating zone (FZ) technique) to discover what effect varying preparation conditions (e.g. using oxidizing or inert environments) have upon structural and magnetic properties. It should be noted that controlled solidification from the melt represents a near equilibrium growth process, in contrast to many thin film growth mechanisms which are non-equilibrium in nature. 2. Experimental procedures 2.1. Powder and feed rod preparation Two starting material combinations were used to make Co-doped TiO2 powders and feed rods (for FZ
crystal growth), namely mixes of CoO+TiO2 powders and Co+TiO2 powders (in both cases the TiO2 was in the form of anatase). The starting powders were mixed in the desired ratio and the Co content was varied from 2 to 8 atomic % with respect to titanium. The CoO+TiO2 mixtures were calcined at 1100 1C for 8 h under oxygen, while the Co+TiO2 mixes were calcined under vacuum at 106 mbar or oxygen for 8 h at 1100 1C. After grinding and phase identification using XRD, calcined powders were used for both magnetic measurements and feed rod preparation. To make starting rods for FZ crystal growth, powder was packed and subsequently sealed into a rubber tube while evacuating air using a vacuum pump. The powder was then compacted into a rod, typically 6 mm in diameter and 70 mm long, using a hydraulic press under an isostatic pressure of 700 bar. After removal from the rubber tube, the pressed rods were sintered at 1100 1C for 8 h in oxygen for the powders calcined in oxygen, or in vacuum for powders calcined in vacuum. (Higher sintering temperatures were found not appropriate due to the evaporation of phases including cobalt which led to departures from the stoichiometric composition). 2.2. Floating zone crystal growth Single crystals were grown from the sintered feed rods using a four-mirror optical FZ furnace (Crystal System Inc. FZ–T–10000–H–VI–VP) having four 1 kW halogen lamps as the heating source. Pure rutile single crystals, oriented with the ‘c’ axis along their lengths, were used as seed crystals for all growths. The molten zone was moved upwards during all float zoning runs with the seed crystal being at the bottom and the feed rod above it. Growths were performed at 5 mm/h with rotation rates of 15 rpm for the growing crystal (lower shaft) and 0 rpm for the feed rod (upper shaft) in either an oxygen or argon environment. 2.3. Characterization Phase identification and the determination of crystal structure of all samples (both synthesized powders and single crystals) were performed using X-ray powder diffraction techniques, while X-ray Laue diffraction was utilized to check the crystalline quality and orientations of the crystals prepared. The compositional homogeneities of as-grown single crystals were analyzed by energy dispersive X-ray (EDX) technique using scanning electron microscopy (SEM). Magnetization measurements were carried out using a Lake Shore Cryotronics, Inc. vibrating sample magnetometer (VSM) on both calcined powders and oriented single crystals (approximately 3 3 1.5 mm along the ‘a’ and ‘c’ directions, respectively) which were cut from the as-grown crystals.
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3. Results and discussion 3.1. Co-doped TiO2 powders One apparent difference between powders after calcining was that powders synthesized in oxygen were greenish in colour, while powders synthesized in vacuum were greyer. Phase identification using X-ray diffraction indicated that the combination of CoO+TiO2 (anatase) synthesized in oxygen was converted during calcining mainly to the rutile structure with some weak second phase reflections consistent with CoTiO3, the intensities of which increased with increasing Co content from 2% to 8%. This is illustrated in Fig. 1a, which shows a typical XRD pattern taken from the powder with a Co content of 4%. These results are consistent with the phase relations in the TiO2–CoO pseudo-binary system [31]. Starting material combinations of Co+TiO2 (anatase) powder (%Co ¼ 2–8) were also calcined in oxygen at 1100 1C for 8 h and XRD from the resulting powders indicated the rutile structure with extra peaks consistent with CoTiO3, i.e. similar to powders prepared using CoO:TiO2 mixtures synthesized in oxygen. However, XRD patterns taken from powders of Co+TiO2 (anatase) mixture synthesized in vacuum showed only reflections for the rutile structure (Fig. 1b) with no clear evidence for any second phases (e.g. Co, CoO, Co3O4, CoTiO3 or CoTi2O5), although the presence of some small amounts of second phase, below the detection limit of the XRD technique, is, of course, possible. 3.2. Magnetic properties of Co-doped rutile powders The magnetic properties of the as-synthesized powders described in Section 3.1 were measured at room temperature. Magnetization measurements performed on initially
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CoO:TiO2 mixtures synthesized in oxygen (which X-ray diffractometry indicated as having the rutile structure with some second phase CoTiO3) showed that the powders were paramagnetic (Figs. 2a–c). As is evident from this figure, increasing the Co content from 2% to 8% (which should increase the amount of second phase CoTiO3 in the powders) increased the magnetization values. This is consistent with the fact that CoTiO3 is a paramagnetic material at room temperature (CoTiO3 is antiferromagnetic at low temperature with a Neel temperature of 35.6 K), while rutile is a diamagnet. In contrast, magnetization vs. field plots measured from initially Co:TiO2 mixtures synthesized in vacuum (which led to the formation of the rutile structure with no obvious evidence of other phases on the XRD patterns shown in Fig. 1b) showed room temperature ferromagnetism for all powders (Figs. 2d–f) with magnetization values two orders of magnitude higher than those measured from powders synthesized in oxygen. While increasing the Co content from 2% (Fig. 2d) to 8% (Fig. 2f) increased the magnetic moment per kilogram for these Co:TiO2 mixtures synthesized in vacuum, no significant difference was observed in the calculated magnetization values per Co atom for Co contents between 2% and 8%, which were all in the range 1.8–1.9 mB/Co atom. Since these magnetic moment values per Co atom for vacuum-prepared powders are close to those of pure cobalt (1.7 mB), the obvious conclusion is that the enhanced magnetization in these powders could be attributed to the presence of Co particles which were too small to be detected by the X-ray diffraction technique. This observation of room temperature ferromagnetism in Co-doped TiO2 powders prepared in a reduced oxygen environment is consistent with the work by Cho et al. [28], where the presence of Co metal clusters was observed in
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Fig. 2. Magnetization measurements performed on CoO:TiO2 powders synthesized in oxygen for Co contents of (a) 2%, (b) 4% and (c) 8%, and Co:TiO2 powders synthesized in vacuum with (d) 2%Co, (e) 4%Co and (f) 8%Co.
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vacuum-annealed powders (synthesized by the sol-gel method) using a high resolution TEM. The topic of the exact origin of the ferromagnetism seen in vacuum prepared powders is returned to in a later section on single crystals. Intensity (arb. units)
Crystal growths were only attempted with Co contents of 4% and 8%, rather than 2%, since detection of the formation of any second phases would be more reliable at higher cobalt contents. Growths using feed rods sintered in oxygen with a Co content of 8% were not successful, however, due to the formation of a mushy (liquid+solid) region below the molten zone, which led to the zone falling out. By reducing the Co content to 4%, though, it proved possible to perform a successful growth due to the formation of a more solid region below the melt zone and, consequently, a more stable molten zone. This more solid region below the molten zone at lower Co contents is consistent with the reported phase relations [31] which predict that solidification of a 4% Co composition should start with a lesser amount of liquid compared to 8% Co. Growth using feed rods sintered in vacuum also exhibited a stable molten zone for the 4%Co composition. A gas pressure of 5 bar (either oxygen using the feed rod sintered in oxygen, or argon using the feed rod sintered in vacuum) was applied during the crystal growths. These relatively high pressures were intended to both sharpen the temperature gradient along the growing crystal [32], thereby attaining a more solid region below the molten zone, and to limit evaporation of any volatile phases (as had occurred during sintering at 1600 1C under ambient pressure). In both cases, crystals of approximately 5 mm in diameter and 30 mm in length were grown, which initially appeared to be black in colour, although cutting and polishing revealed that while the crystal grown in argon was black, the crystal grown in oxygen was actually dark red. X-ray Laue diffraction photographs taken along the lengths and cross-sections of the crystals indicated that both exhibited an equally high crystalline quality and that the crystals had adopted the orientation of the seed crystal (i.e. with the ‘c’ direction along their length). X-ray diffraction patterns taken from sections of the asgrown crystals after crushing indicated that the crystal grown in oxygen had the rutile structure with second phase CoTiO3 (Fig. 3a) similar to the starting powder. However, the crystal grown in argon showed the rutile structure with very weak and broad extra peaks consistent with the 002 and 101 hcp reflections for pure Co at 2y values of 44.71 and 47.51, respectively (Fig. 3b). This contrasts with the starting powder for this crystal in which no extra reflections beyond those expected for the rutile structure were evident. The broadening seen in these extra peaks could indicate strain and/or that they were not from pure Co, but contained some Ti or O.
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Fig. 3. The XRD patterns taken from the crushed Co-doped rutile (%Co ¼ 4) single crystals grown under 5 bar pressure of (a) oxygen, and (b) argon. The crystal grown in oxygen had the rutile structure with second phase CoTiO3, while the crystal grown in argon showed the rutile structure with very weak and broad extra peaks consistent for pure Co.
For further characterization, back-scattered SEM micrographs were taken from polished cross and longitudinal sections of the crystal grown in oxygen (Figs. 4a–b) and argon (Figs. 4c–d). The formation of second phases in both crystals is clearly shown in these back-scattered images, although the inclusions are differently shaped. EDX chemical mapping images taken from the crystal grown in oxygen (Fig. 5a) revealed that Ti and O were uniformly distributed throughout the whole area, while the secondary phase particles also contain Co. Conversely, the EDX maps taken from the crystal grown in argon (Fig. 5b) showed a significant drop in Ti and O signal intensities in the precipitates indicating that the secondary phase is either Co or very rich in Co. While the EDX technique is not particularly suited to examining light elements such as oxygen, analyses
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Fig. 4. SEM back scattered micrographs of cross and longitudinal sections cut from 4%Co:rutile single crystals grown in oxygen (a and b) and argon (c and d). Growth in oxygen led to the formation of CoTiO3 second phase, while the crystal grown in argon had Co rich particles elongated along the growth direction (‘c’).
performed on both crystal types (i.e. grown in oxygen and argon) indicated that the matrix was very close to TiO2 with no clear evidence of Co incorporation, while all second phases consistently included Co. Although the technique was found not to be sufficiently accurate to identify the exact Co–Ti–O second phase in the sample grown in oxygen, results are consistent with the XRD data that indicated the second phase to be CoTiO3. For the crystal grown in argon, the EDX results indicated that second phases were Co rich with the average atomic percentage of more than 90% (with the atomic percentage of Ti varying between 3.72% and 10%), which again agrees with the XRD data from the as-crushed crystal that showed broad peaks consistent with hcp Co. Since these Co-rich inclusions consistently have a square shape in cross section and are elongated along the ‘c’ direction of the matrix (growth direction) it can be surmised that they are either single crystal or polycrystalline with a preferential crystalline orientation. Concerning the ongoing controversy of whether Co atoms actually incorporate into the TiO2 structure or form clusters or inclusions, it would appear likely from the evidence of the above results for bulk crystals, that if significant Co incorporation does occur, it is a nonequilibrium effect peculiar to thin films. 3.4. Magnetic properties of Co-doped rutile single crystals The magnetic response from the 4%Co:TiO2 single crystal samples was measured as a function of magnetic
field along the ‘c’, ‘a’ and 451 from ‘c’ towards ‘a’ directions for the sample grown in argon, and along ‘c’ and ‘a’ for the sample grown in oxygen, as shown in Fig. 6. These measurements clearly indicate that room temperature ferromagnetism was observed only in the crystal grown in argon, while crystals grown in oxygen showed paramagnetic behaviour. The room temperature paramagnetic behaviour of crystals grown in oxygen can be explained by the formation of paramagnetic second phase CoTiO3, as described earlier for the powder samples. This weakly ferromagnetic argon-grown crystal had a saturation magnetization of about 0.24 A m2 kg1, with the easy magnetization direction along ‘c’ and the hard direction of magnetization along ‘a’. This magnetization value equates to 0.083 mB/Co atom if the crystal composition is taken to be the same as the starting powder (i.e. 4%Co:TiO2), although there are some doubts as to whether this is the case (see below). A single FZ experiment under argon was also performed with a CoO:TiO2 (%Co ¼ 4) mixture initially calcined in oxygen, which gave a very weakly ferromagnetic sample with a saturation magnetization about 0.05 A m2 kg1 along the ‘c’ direction. The various routes used to prepare the crystals from starting mixes to final FZ growth are summarized in Fig. 7. The differences of the magnetic properties of the two main crystal types, as shown in Fig. 6, can be explained from the results of the XRD and SEM microstructural studies. In this regard, the room temperature ferromagnetism observed in crystals grown in argon is almost certainly due to the existence of ferromagnetic Co-rich clusters.
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Fig. 6. M–H curves measured at room temperature for 4%Co:rutile single crystals. Room temperature ferromagnetism with magnetization 0.24 A m2 kg1 (0.083 mB/Co atom) was observed in the crystal grown in argon, while a weak paramagnetic behaviour was obtained from the crystal grown in oxygen.
Fig. 5. EDX chemical mapping images of cross-sections cut from 4%Co:TiO2 single crystals grown in (a) oxygen and (b) argon.
However, one interesting feature is that the saturation magnetization of these crystals at 0.24 A m2 kg1 is significantly lower than the magnetization of the feed rod/starting material (i.e. the 4%Co:TiO2 powders prepared in vacuum), which had values of about 5.61 A m2 kg1. As noted previously, this latter value equates to 1.9 mB/Co atom, which is close to the magnetic moment of pure cobalt, 1.7 mB/Co atom. To ensure that the form of the magnetization samples (powder for the starting material against bulk single crystal) was not a factor in these differences, a further magnetization measurement was performed on the single crystal grown in argon after powdering, but the measured saturation magnetization value (0.204 A m2 kg1) was not significantly different to the value obtained from the crystal before powdering (0.24 A m2 kg1). Explanations for the significant reduction in magnetization values for 4%Co:TiO2 after FZ crystal growth include that it resulted from a reduced amount of Co in the final
crystal compared to the start material and/or a reduction in the magnetization of the Co atoms present. A reduction in Co content in the as-grown crystals could occur as a result of evaporation of Co during FZ, or from segregation of Co during the zoning process. Although no overt evidence for Co volatilization losses (e.g. condensation of Co onto the furnace tube) were detected during the crystal growth, some loss of Co is likely; losses of about 1.5–2.0%Co were reported [33] during melting and crystal growth of CoTi crystals, for example, which has a significantly lower melting point than TiO2. Segregation of components during FZ is a common occurrence, especially when one component has a low solubility in the matrix, as Co appears to have in TiO2. Thus during FZ of Co-doped TiO2, it is not unrealistic to presume that Co would be ‘‘pushed’’ along by the molten zone leaving a reduced amount of Co in the as-grown crystal. Some strength to this argument was given by a magnetization measurement taken from the final part of the rod traversed by the molten zone during the crystal growth process (i.e. the section directly above the as-grown crystal) which had magnetization values an order of magnitude higher than the crystal at 2.3 A m2 kg1. Besides a reduction of the Co content as a result of the FZ process, a lowering of magnetization values of crystals may also result from changes in the magnetic/electronic structure of the Co atoms themselves. In this regard, several studies have indicated that Co between Ti layers (as the local environment) has very small magnetic moments (down to 0.07 mB [34]), while in Co100xTix thin films, increasing the Ti concentration significantly decreased the magnetic moment of Co [35–37]. Although the X-ray data from the crushed crystal indicated that the Co-rich inclusions still retained a hexagonal crystal structure, the solubility limit for Ti dissolved in Co has a
ARTICLE IN PRESS S.M. Koohpayeh et al. / Journal of Magnetism and Magnetic Materials 320 (2008) 887–894
Starting mixtures CoO+TiO2
Calcining atmosphere
Magnetic properties of calcined powders used to make feed rods
Oxygen
Paramagnetic
Growth atmosphere
Co+ TiO2 Co+ TiO2
Vacuum
Ferromagnetic
893
Magnetic properties of FZ grown crystals
Oxygen
Paramagnetic
Argon
Very weakly ferromagnetic
Argon
Weakly ferromagnetic
Fig. 7. Preparation routes and magnetic properties of Co-doped rutile powders and single crystals.
maximum of 15% at 1190 1C. Overall, it is thought probable that a combination of circumstances (i.e. some reduction in Co content together with a lowering of the magnetization of Co atoms due to their local environment) explains these results, although further work is needed to discover which is the dominant effect. The weak dependence of the magnetic properties of the single crystal on a preferred direction (i.e. easy and hard magnetization directions along ‘c’ and ‘a’ directions, respectively) shows that this material exhibits some magnetic anisotropy. Since the matrix, rutile, is a diamagnetic material, this apparent anisotropy must come from the Co-rich second phases and probably arises either from the shape anisotropy of the individual precipitates (due to magnetostatic effects, i.e. the demagnetizing field was less when measuring magnetization along the long axis of these needle shaped inclusions) and/or from the intrinsic magnetocrystalline anisotropy of Co (if these second phase inclusions are single crystals or preferentially aligned). However, since these inclusions are small in size, shape anisotropy effects could be more dominant than magnetocrystalline anisotropy. To end our discussion, it is worth commenting upon the very recent work on the magnetic properties of undoped TiO2d films grown by PLD [25] which indicated that these oxygen deficient anatase films show ferromagnetism at temperatures up to 880 K even without the introduction of magnetic ions. This magnetism was attributed to the presence of defects on oxygen sites (as a result of interfacial defects created at the film-substrate interface and from processing in an oxygen deficient atmosphere during PLD growth) leading to the creation of Ti3+ and Ti2+ ions (with unpaired 3d electrons) by stripping electrons from the s band. Thus questions are now being asked [41] as to what degree transition metal doping plays an essential role in producing ferromagnetism in thin films of non-magnetic oxides. Earlier work in the authors’ laboratory on FZ crystal growth of TiO2 has shown that oxygen deficient (i.e. blue colour) bulk rutile single crystals can be grown in the image furnace by using a higher than normal molten zone temperature [38]. When the magnetic properties of these crystals were measured, however, they gave no indication of ferromagnetism at room temperature. While this work on bulk crystals cannot be directly compared with the thin film work of Yoon et al. (since the TiO2 crystal was in the
form of rutile while the films were anatase), it would appear to confirm that oxygen deficiency by itself does not significantly affect the magnetic properties of all oxygen deficient TiO2-based crystals. However, it should be noted that the Ti–O phase diagram shows several TiO2d phases [39], which contain various types of defects such as Tin+ interstitials, oxygen vacancies and planar defects (e.g. crystallographic shear planes) although the question of which type of defect is dominant in which region of oxygen deficiency, is still subject to debate [40]. It is not inconceivable, therefore, that in bulk crystals one of these oxygen deficient phases could exhibit enhanced magnetic properties even without the strain caused by lattice mismatch in Yoon et al.’s thin film work. 4. Conclusions It can be concluded that an oxygen deficient environment during the preparation of Co-doped TiO2 powders using a solid-state reaction was crucial for the observation of room temperature ferromagnetism. Preparation in oxygen-rich conditions destroyed the ferromagnetism and led to the formation of the paramagnetic second phase CoTiO3. Crystal growth of Co-doped rutile using the FZ melting technique led to the formation of cobalt-based second phases within crystals; crystals grown in oxygen were paramagnetic at room temperature and contained second phase CoTiO3, while crystals grown in argon were ferromagnetic at room temperature and contained Co-rich second phases. No evidence of significant solution of Co into the rutile matrix was found in either crystal type. Acknowledgement This work was supported by the UK Engineering and Physical Sciences Research Council. References [1] S.A. Wolf, D.D. Awschalom, R.A. Buhrman, J.M. Daughton, S.V. Molnar, M.L. Roukes, A.Y. Chtchelkanova, D.M. Treger, Science 294 (2001) 1488. [2] Y. Matsumoto, M. Murakami, T. Shono, T. Hasegawa, T. Fukumura, M. Kawasaki, P. Ahmet, T. Chikyow, S.Y. Koshihara, H. Koinuma, Science 291 (2001) 854.
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