Physical Mechanisms of Exchange Coupling Effects in

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Physical Mechanisms of Exchange Coupling Effects in Nanoparticulate Diluted Magnetic Oxides Obtained by Laser Pyrolysis. Victor Eugen Kuncser, Gabriel Alexandru Schinteie, Andrei Cristian Kuncser, Aurel Leca, Monica Gina Scarisoreanu, Ion Morjan, and George Filoti J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.7b01500 • Publication Date (Web): 05 Apr 2017 Downloaded from http://pubs.acs.org on April 11, 2017

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The Journal of Physical Chemistry

Physical Mechanisms of Exchange Coupling Effects in Nanoparticulate Diluted Magnetic Oxides Obtained by Laser Pyrolysis

Victor Eugen Kuncser1, Gabriel Alexandru Schinteie1*, Andrei Cristian Kuncser1,2, Aurel Leca1,3, Monica Scarisoreanu4, Ion Morjan4 and George Filoti1 1

National Institute of Materials Physics, PO Box MG-7, 077125 Bucharest-Magurele, Romania

2

Faculty of Physics, University of Bucharest, 077125, Bucharest-Magurele, Romania

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University Politehnica Bucharest, Faculty of Electronics, Telecommunications and Information

Technology, Bucharest, Romania 4

National Institute for Lasers, Plasma and Radiation Physics, Bucharest-Magurele, Romania

*E-mail: [email protected]. Phone: +40 21 369 01 85, Fax: +40 21 369 01 77

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ACS Paragon Plus Environment


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Abstract TiO2 nanoparticles, un-doped and doped with Fe, have been prepared by laser pyrolysis and further investigated with respect to morphological, structural and magnetic aspects by transmission electron microscopy,

diffractometry,

Mössbauer

spectroscopy and

magnetometry.

The

obtained

nanoparticles, consisting of mainly anatase phase, agglomerate in clusters of tenths of units and present a large size distribution in the range from 5 to 40 nm. The anatase to rutile weight ratio (about 9) and the morphology of particles is similar in all analyzed samples (doped by up to 12 at. % Fe). Only Fe3+ ions in high spin configuration were observed mainly at the surface of TiO2 nanoparticles, either distributed or forming fine clusters of Fe oxide. Both a paramagnetic phase and a superparamagnetic one with blocking temperature lower than 50 K are superposed over a long range ferromagnetic phase specific to diluted magnetic oxide systems. The influence of doping Fe ions on the magnetic behavior of each phase is discussed in detail. Evidences for interface exchange couplings (with unidirectional anisotropy in specific conditions) between the long range ferromagnetic phase and the fine clusters (antiferromagnetic in nature), which become frozen below the blocking temperature of 50 K are provided. The specificity of the processing route and the physical mechanisms responsible for the observed relevant magnetic features which can be tailored for suitable applications are discussed.

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

Introduction The laser pyrolysis was proven to be an adequate procedure to obtain nanocomposites with

controlled size and composition of the components.1-6 Recently, successful attempts to produce titania (TiO2) nanoparticles by the above mentioned method were reported,7 followed by the description of their magnetic and local interactions properties.8 Among the last specific applications of titania (excepting fillers in paints, pigments for inks, cosmetics, etc.) there are the catalytic ones, including the photocatalytic properties allowing the decomposition of organic pollutants or photogeneration of hydrogen. Such properties are nevertheless enhanced for the nanosized forms of titania due to the higher specific surface and activity. Diluted magnetic oxides including TiO2, pure or in which Ti ions are substituted by magnetic ions (usually of 3d type) form a new class of smart materials with interrelated optical, magnetic and magneto-transport properties of potential interest in spintronics and sensoristics.9,

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There are valuable reports on magnetic properties of semiconductors or oxides in

nanosized forms without doping with magnetic elements, the specific magnetism being known as a “d0” type. It is accredited that the magnetic interactions in this class of materials are controlled by defects, vacancies or surface centers.11 In case of un-doped TiO2 nanoparticles with both most known structures of anatase and rutile, the d0 magnetism might not exclude also the presence of the magnetic Ti3+ ions, as an indirect effect of the intrinsic defects and oxygen vacancies. In the case of oxides or semiconductors doped with transition-metal (TM) ions, the mechanisms responsible for the magnetic behavior are even more complex, consisting usually in superposed effects due to isolated or clustered TM magnetic moments (localized magnetism) as well as to the long range magnetism persisting up to high temperatures (and known as room temperature ferromagnetism –RTFM-), specific also to some of the un-doped nanosized oxides. This long-range magnetism specific to the diluted magnetic 3



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semiconductors (DMS) or oxides (DMO) remains controversial by both manifestation and origin. There are different models proposed for explaining RTFM in doped TiO2 (carrier mediated exchange, magnetic polarons, etc.12), but no definite conclusions, for the simple reason that it seems that the mechanism involved appeared to be strongly related to the preparation conditions of the samples. Although this type of magnetism seems to be not directly related to the presence of the TM ions (e.g. as suggested by the presence of the d0 magnetism), it is still important to know the indirect effect of the doping element on the involved long range magnetism in nanosized TiO2 depending on the sample morphology and processing routes. Concerning the up to date reports, the magnetic data were discussed in the case of rutile TiO2 nano-clusters,

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in terms of excess of oxygen treatments and the presence of trivalent Titanium and

oxygen vacancies in pair’s neighborhood. A report on the anatase TiO213 emphasized the role of oxygen deficient treatment atmosphere and the existence of a ferromagnetic behavior up to 880 K. The presence of defects and surface clusters in 10 nm thick thin films underlined the interface mechanism involved in the magnetic properties of the systems.14 It was reported that defects or vacancies ( induced by annealing in vacuum the Fe doped TiO2 nanoparticles) increase the magnetization.

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This feature

was confirmed by other papers for different morphologies, for example in nanoribbons16 or in nanorods, 17 among some other. According to one of our previous studies on DMO systems,

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the weak RTFM observed in

hydrothermally obtained anatase TiO2 nanoparticles doped with a very low amount of Fe (less than 1 at. %) does not depend on Fe doping. Most of the Fe ions are isolated or form polynuclear clusters with antiferromagnetic coupling. The results reported in the present paper are based on the strategy to increase the density and size of the anti-ferromagnetic Fe clusters in order to induce an observable exchange coupling between the long-range ferromagnetic phase and the anti-ferromagnetic 4


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nanoclusters. This may be obtained by increasing the amount of Fe in the TiO2 nanosized system, which can be conveniently achieved by an appropriate processing method (e.g. laser pyrolysis in this case). In consequence, this paper will compare, discuss and emphasize specific magnetic aspects and parameters related to un-doped and, respectively, Fe-doped TiO2 compounds (up to 12 at. % Fe) obtained by laser pyrolysis. Because of the exchange coupling induced in samples of higher level doping, increased coercive forces and exchange bias field are reported for the first time, offering new perspective for the applications of the DMO and DMS systems in spintronics and related fields.

 Experimental The laser synthesis technique is based on the resonance between the emission line of a CW CO2 laser (λ = 10.6 μm) and the infrared absorption band of a gas precursor. The synthesis of Fe–doped TiO2 nanopowders was performed via a modified version of the pyrolysis set-up in which TiCl4, Fe(CO)5 and air, as gas phase precursors are simultaneously allowed to emerge into the flow reactor where they are orthogonally crossed by the laser beam (400 W nominal power). Since the reactants TiCl4 and Fe(CO)5 do not possess IR absorption bands at the CO2 laser radiation, ethylene served as an energy transfer agent and as carrier for the TiCl4 vapor precursor. The aggregation of the hot, freshly nucleated particles is rapidly stopped, by rapid cooling/freezing outside the reaction zone. It was shown that for pure TiO2 the transformation from anatase to rutile takes place in air at 500 C. However, it is worth to note that, as demonstrated by the selected analytical characterization methods (see the X-ray Diffraction and Transmission Electron Microscopy analysis below), the temperature reached during the synthesis procedure is lower than that required for the complete transformation of anatase to rutile phase.

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Transmission Electron Microscopy (TEM) data, both conventional and as Selected Area Diffraction (SAED) have been obtained via an ARM-200F Analytical Electron Microscope with an aberration corrected field emission gun. The device is also equipped with STEM (scanning transmission electron microscopy), EDX (Energy Dispersive X-ray Spectroscopy) and GIF (Gatan Image Filter) with EELS (Electron Energy Loss Spectroscopy) capability units. The samples consisted of powder deposited on 3 mm diameter copper grids, covered with a thin layer of carbon. The magnetic measurements have been performed on a Superconducting Quantum Interference Device (SQUID) MPMS 7T by Quantum Design working under the sensitive reciprocal space option. Zero field cooled–field cooled (ZFC-FC) measurements as well as hysteresis loops were obtained at different temperatures between 5 K and 300 K under applied fields of up to 2 T. The Mössbauer spectra (MS) were acquired using a constant acceleration spectrometer with symmetrical waveform. A

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Co (1.0 GBq) source has been used. Temperature dependent Mössbauer

measurements (4.2–295 K) were performed in a He close cycle cryostat. The commercially available NORMOS program was considered for fitting the collected spectra.

 Results and Discussions Three samples, having the corresponding codes S0, S6 and S12 (the numbers approach well the Fe content in at. % as resulted from the EDX data), have been processed by laser pyrolysis and carefully investigated. The X-ray Diffraction (XRD) patterns of all samples as well as the TEM images (conventional and SAED) of sample S12 are shown in figure 1.

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Figure 1. The XRD spectra of samples S0 to S12 (left side) and conventional TEM image of sample S12 with selected area electron diffraction evidencing the corresponding plane of anatase TiO2 (a=0.378 nm)

According to the XRD data, the main formed compound is TiO2 in its both stable structures of anatase and rutile (as indexed in the above figure by A, and respectively, R). Expectedly, the iron in Fe doped samples is, either dispersed or agglomerated in very fine (non-crystalline) clusters whereas traces of carbon from ignition of ethylene and of Fe precursor are amorphous like and therefore, are not evidenced by the XRD patterns. According to the XRD data, the relative amount of anatase is much higher than that of rutile, their relative ratio decreasing slightly from about 9.0 in S0 to 8.7 in S12. Irrespective of the Fe content the lattice parameters are a=0.379(1) nm and c=0.952(4) nm in anatase and a=0.459(1) nm and c=0.296(1) nm in rutile). The structural coherence length (related to the 7



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nanoparticle size) is of 18(2) nm for anatase and 20(1) nm for rutile. EDX data point for an additional amount of Carbon in all samples (e.g of about 18(1) at. % in S0, 20(1) at. % in S6 and 12(1) at. % in S12). The conventional TEM image of sample S12 is also shown in figure 1, being representative for all three samples. The samples consist of mainly spherical nanoparticles with sizes ranging between 5 and 40 nm (see figure 1). However, the nanoparticles agglomerate in large clusters of tenths of units sometime embedded in carbon allotropes. The nanoparticles with crystalline structure belong to TiO2 (mainly anatase phase) as directly proven also by the Electron Diffraction pattern presented in the inset of the same figure. The Fe containing samples have been analysed by using the STEM mode of the microscope, coupled with EELS spectroscopy as exemplified in the figure 2 for sample S6. The analysis has been focused on a nanoparticle of around 20 nm, where the image was unaffected by the presence of allotrope carbon or carbon located on the copper grid. The nanoparticle was studied via Spectrum Imaging, an EELS spectrum being automatically acquired in every point (with respect to a specific resolution) of the area framed in the open green square of the upper-left figure. The bottom of figure 2 is showing the spectrum image (left side) together with two EELS patterns extracted from points (a) and (b) in the spectrum image (right side), as corresponding to arbitrary locations at the middle and the border of the nanoparticle. Finally, the EELS-spectrum imagining map of the investigated nanoparticle is shown in the upper-right figure, where red pixels corresponds to Ti enriched and green pixels to Fe enriched elementary domains, respectively. Quantitatively, a relative atomic ratio Ti:O:Fe of 1:0.58(8):0.03(1) is evidenced in point (a) and a relative atomic ratio of 1:0.54(7):0.07(1) in point (b). This result suggests that the concentration of Fe is significantly higher at the border of the nanoparticle than in the middle, what is also suggestively confirmed by the elemental map in the upper-right side of 8


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figure 1. Hence the electron microscopy data support well the average Fe content of about 5 at. % (6 at. % provided by EDS) but also give evidence for Fe ions distributed mainly on the surface of TiO 2 nanoparticles as previously reported.12

Figure 2. STEM Dark-Field image of TiO2 nanopatricles in sample S6 with the selected frame for the spectrum image (upper-left figure). The spectrum image (with the involved sampling for EELS acquisition) for the 20 nm TiO2 nanoparticle with EELS patterns acquired in points (a) and (b) corresponding to the middle and the border of the nanoparticle (down figures). The EELS-spectrum imagining map of the investigated nanoparticle with red pixels corresponding to Ti and green pixels to Fe (upper-right figure).

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One has to consider that the actual result might also provide an indirect indication of the fact that, excepting the Fe ions distributed at the particle surface, additional segregation of Fe in larger clusters of Fe oxide (a few nm in size) can take place at the boundary between the TiO2 nanoparticles. However, while the electron microscopy is a spatial selective technique, additional investigation tools are required in order to give direct evidence for such behaviour (e.g. Mössbauer spectroscopy or magnetometry data). For the sake of a more understandable perception of Fe related magnetic properties in these systems, the

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Fe Mössbauer spectra acquired at 5 K and 50 K on samples S6 and S12, respectively,

are shown in figure 3. At the lowest temperature, excepting the main central paramagnetic pattern, there is also a magnetically split, rather broad, component which is clearly evidenced in case of sample S12 (figure 3c). The best fit in case of sample S12 was obtained via a central doublet whose hyperfine parameters (isomer shift, IS=0.49 mm/s and quadrupole splitting, QS=0.70 mm/s) are specific to paramagnetic Fe3+ ions1, 4, 12 and two broad sextets with hyperfine magnetic fields, Bhf , of 44.5 T and 37.6 T, respectively and with IS values of 0.49 and 0.45 mm/s, respectively (one unit error of the last digit has to be considered anywhere). It is worth mentioning that the very large linewidth of the two fitting sextets (over 1 mm/s) gives indications for overall distributed Fe configurations possible to be alternatively fitted by a hyperfine field distribution. However, the two nonequivalent sextets used to fit conveniently the distributed configurations of the magnetically ordered iron are assigned by the specific values of the hyperfine parameters also to Fe3+ ions.1, 4, 12 The sextets reflect that a part of the total iron ions are close enough to induce interactions among them, leading to small clusters with dispersed sizes and strengths of the intra-cluster coupling (poorly crystallized Fe2O3 species). At 50 K, the low temperature sextets collapse into an additional doublet with IS=0.48 mm/s and QS=0.55 mm/s. This relaxation behavior of the low temperature magnetic pattern provides strong support for its 10


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assignation to Fe oxide clusters with a blocking temperature lower than 50 K. An average size of a few nm can be roughly estimated for the Fe oxide clusters by using a Neel-Brown type relaxation law with anisotropy constant specific to the maghemite, a blocking temperature of 40(10) K, a time relaxation constant of order of 10-10 s and a time window of 5·10-9 s, specific to

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Fe Mössbauer

spectroscopy.4

Figure 3. The Mossbauer spectra of iron doped sample: a) S6 at 5 K, b) S6 at 50 K, c) S12 at 5 K, d) S12 at 50 K.

Concerning the 5 K Mössbauer spectrum of sample S6 (figure 3a), in spite of a poor evidence of the sextet component, the fit is consistently improved by using in addition to the central paramagnetic doublet (IS=0.48 mm/s, QS=0.66 mm/s), the two sextets considered also in case of sample S12 at 5 K 11



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(this time Bhf is 49.0 and 38.1 T, respectively and IS is 0.45 and 0.35 mm/s, respectively). The same collapsing behavior of the sextet pattern into a central doublet is observed at 50 K, as specific to Fe based clusters with a blocking temperature of less than 50 K. The mentioned hyperfine parameters give clear support for only Fe3+ ions also in sample S6, with the mention that the relative amount of clustering Fe3+ ions is higher in S12 than in S6, as directly resulting from the relative spectral area of the sextets appearing in the 5 K spectra of figure 3 (e.g. 61% in S12 and 41% in S6). In conclusion, Mössbauer spectroscopy data point to the presence of only Fe3+ ions in high spin state, which can be either highly dispersed down to the atomic level (being reflected by the central paramagnetic doublet at 5 K) or forming magnetic nanoclusters (a few nm in size) with blocking temperatures below 50 K (being reflected by the broad magnetic sextet feature at 5 K, collapsing into a central doublet at 50 K).

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Figure 4. ZFC-FC behaviour of the investigated samples: a) S0, b) S6, c) S12. The cooling and measuring field was 100 Oe. The next step toward the understanding of the overall magnetic behavior of the samples was provided by magnetometry data. Zero Field Cooled (ZFC)-Field Cooled (FC) curves (with cooling and measuring field of 100 Oe) are presented in figure 4. As general observations, these present a finite magnetization at high temperatures (in 100 Oe measuring field), a diverging behavior at low temperatures with the FC curves over the ZFC ones and a specific increase at very low temperature. All these aspects support the long range magnetic order at high temperatures (e.g. the RTFM phase), the existence of weak antiferromagnetic couplings among the magnetic entities and the presence of paramagnetic like centers in the samples (more evident in the Fe doped), respectively. Magnetic hysteresis loops have been obtained at increasing temperatures and in fields lower than 20 kOe, for all samples, after cooling down the systems in the absence of a magnetic field. For samples S0 and S12, the same hysteresis loops have been successively acquired also after cooling the sample in a field of 20 kOe down to 5 K. Representative loops at 3 different temperatures are shown in the upper insets of figure 5 from a) to d) as described in the figure caption. The magnetization of sample S12 in high fields is almost one order of magnitude higher than in case of sample S0 with a relatively enhanced diamagnetic component above 100 K and a significantly reduced paramagnetic component below 50 K. While the diamagnetic susceptibility does not depend on temperature, it was estimated from the negative slope of the loop collected at 300 K (where the paramagnetic contribution becomes negligible) for each sample. Further on, the slopes of the M(H) curves at lower temperatures have been estimated and compensated by the diamagnetic contribution in order to derive the pure paramagnetic susceptibility. Then, after subtracting both the diamagnetic and paramagnetic contributions at each temperature (following the linear M(H) dependence in enough high 13



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fields up to 15 kOe), saturated hysteresis loops with finite coercive fields (more visible in the lower insets of figure 5) are obtained and attributed to only those phases presenting magnetic order.

Figure 5. Hysteresis loops belonging to the magnetic ordered phases of samples S12 (left side) and S0 (right side), obtained after diamagnetic and paramagnetic corrections. Graphs a) and b) correspond to sample S12 after cooling in zero field and in 20 kOe, respectively, whereas graphs c) and d) correspond to sample S0, cooled in the same conditions. In the upper inset of each figure are shown the raw data whereas in the lower one, the central part of the loops is presented with a better field resolution.

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The temperature evolution of the saturation magnetization of these phases is presented in figure 6 (a). It has to be observed that in all cases the magnetization is rather constant from 300 K down to above 50 K and exhibits an increase at low temperatures (below 50 K). Although this increase is very pronounced in the Fe containing samples, it might be still assumed also in sample S0 (at a much lower temperature).

Figure 6. Temperature evolutions of: magnetization at saturation (a), coercive field (b) and exchange bias field (c) of the analyzed samples cooled in different conditions. The relative errors, including the remanence of the superconducting magnet, of the values displayed on the graphs are less than 5 %.

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Having also in mind the Mössbauer data which point for the presence of Fe oxide nanoparticles with blocking temperatures below 50 K, one has to assign the strong increase of the magnetization at low temperatures to the Fe oxide nanoparticles which become more and more magnetically frozen. The almost constant magnetization above 100 K has to be assigned to a magnetic phase exhibiting longrange interactions (leading to the so called RTFM). By extension, we attribute the weak increase of the saturation magnetization in sample S0, also to some magnetic nanoclusters due to gathering magnetic defects (related to either oxygen vacancies or magnetic Ti3+, other magnetic entities than in samples containing doping Fe). It is worth mentioning that the magnetization of the RTFM phase is increasing with the total Fe content in the sample, manifestly faster than linearly. If the inverse of the pure paramagnetric susceptibilities (extracted from the hysteresyis loops as mentioned before) are plotted versus temperature (there is an almost linear dependence at temperatures higher than 50 K, as corresponding to the equation 1/χ = (1/C)·T ), a direct information is furnished about the inverse of the Curie constant C=Nµ2/3kB, where kB is the Boltzmann constant, µ is the magnetic moment of the magnetic center and N is the density of magnetic centers (if susceptibilities are given on emu/g/Oe, N can be taken as number of magnetic centers per gram of sample). In the present case the magnetic centers responsible for the paramagnetic contribution are the Fe3+ ions and therefore µ=5µB. Accordingly, from the slope of the inverse susceptibility versus T providing values for 1/C, the number of effective magnetic centers per gram of sample, N, is estimated at 0.7(1)*1019 in sample S6 and 2.0(2)*1019 in sample S12. On the other hand, taking into account the atomic percentage of elements in each sample (provided by EDX), a relative amount of Fe of 5(1) wt % is straightforward estimated in case of sample S6 and of 10(1) wt % in case of sample S12. Hence, the total number of Fe3+ ions per gram of sample, N*, is about 50(5)*1019 in sample S6 and 100(5)*1019 in sample S12, that is more than 50 times higher than the number of effective magnetic centers carrying 5µB. This result (consisting in a 16


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huge discrepancy between the number of Fe3+ ions and the number of effective magnetic centers carrying 5µB each) can be explained only by the formation of clusters of antiferromagnetically coupled Fe3+ ions. Although clusters with an average odd number of about 50 antiferomagnetically coupled Fe ions may be assumed from the above estimation, according to the Mössbauer data, a bimodal size distribution of the clusters has to be considered as follows: (i) 1/3/5…Fe3+ ions giving rise to the paramagnetic doublet in the 5 K spectra of the Fe containing samples and (ii) clusters consisting of a few hundreds of antiferromagnetic coupled Fe3+ ions with low net magnetic moment and superparamagnetic behavior above 50 K, as accounted by the collapsing sextet in the low temperature Mössbauer spectra. This peculiar behavior of the doping magnetic ions seems to be characteristic to DMS/DMO systems.12, 18

A last, but important, observation is related to the field asymmetry of the loops collected at low temperatures after field cooling samples S0 and S12 (see lower insets of figure 5 b) and d). It is well known, that such asymmetries are related to interfacial exchange interaction (including a unidirectional component) between a ferromagnetic phase and a related phase involving antiferromagnetic couplings (e.g. anti-ferromagnetic, ferromagnetic or spin glass type).19-23 As macroscopic effects, the unidirectional anisotropy (usually induced by field cooling the sample below an exchange bias blocking temperature24) involves both an increase of the coercive field and the shift of the hysteresis loop by the so called exchange bias field (usually negative). The temperature evolution of these 2 parameters (coercive field Hc and exchange bias field Eb) are presented in figure 6b and c in both zero field cooled and field cooled procedures. Concerning the coercive fields in Fe containing samples cooled in zero field, they respect the general increasing trend at decreasing temperatures down to 100 K.25, 26 However, there is always an unexpected decrease at definitely lower temperatures (below 50 K), most probable due to the additional contribution of the Fe oxide nanoparticles (carrying a relatively low net magnetic 17



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moment) with coercive forces weaker than that of the long-range ferromagnetic phase coupled to the blocked nanoparticles. This unexpected decrease is no longer visible in sample S0, where at the lowest measuring temperatures the coercive field presents just a saturation-like behavior. Referring to magnitude, the coercive fields of the RTFM phase are of order of 100 Oe in both samples S0 and S6 and five to six times larger in sample S12, showing another indirect role of the doping Fe ions on the exchange integrals of the long range magnetic ordered phase. Hysteresis loops collected at increasing temperatures after cooling the sample in magnetic field, provide evidence for both increased coercive fields (clearly visible for sample S12) and negative exchange bias fields (of significant magnitude mainly below 50 K), as compared to the zero field cooled samples (see figure 5c and d compared to a and b as well as figure 6b and c). Hence, there is a direct proof for inducing unidirectional anisotropy in DMO systems by the interfacial coupling of the long range ferromagnetic phase to other specific phases (most probable cluster-like) involving antiferromagnetic couplings. The exchange bias field in sample S12-FC reaches a maximum value of about 240 Oe at 25 K, decreasing slightly at about 160 Oe at 5 K, where the contribution of the symmetric loop related to the Fe oxide nanoparticles in the magnetic frozen state has to be maximal. The sensibility of both the exchange bias field and coercive field in sample S12 to the magnetic relaxation of the Fe oxide nanoparticles sustains the importance of this additional phase with anti-ferromagnetic order pinning the spins of the long range ferromagnetic phase. By extrapolation, in case of sample S0 (without Fe) it is naturally to consider that the magnetic centers, related this time to different types of defects located at the nanoparticle surface, might couple in nanoclusters with intrinsic spin disorder and hence with a fractional amount of anti-ferromagnetic couplings, which can pin the long range magnetic phase. However, the exchange coupling is by far much enhanced in the Fe doped samples.

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 Conclusions The processing route (laser pyrolysis in this case) may be a key factor in promoting the exchange coupling and unidirectional anisotropy in DMO systems. The simultaneous presence of anti- and ferromagnetic phases (the last one persistent even at room temperature) both in un-doped and doped (with iron) TiO2 nano-materials pleads for the specific role of morpho-structural aspects influencing the type of magnetic defects and of surface distributed iron entities. The iron appears only in the highest ionic valence (Fe3+) and spin (S =5/2) state. The calculated number of Fe3+ ions forming the clusters, which is 50 times larger than the number of effective magnetic centers (carrying 5 µB as typical to Fe3+), together with the Mössbauer data, give evidence for a bimodal size distribution of the Fe based clusters: (i) very fine ones consisting of single/a few Fe3+ ions with paramagnetic behavior already at 5 K and (ii) larger ones consisting in a few hundreds of Fe3+ ions with superparamagnetic behavior above 50 K. An antiferromagnetic coupling is present between the Fe3+ ions or other local magnetic centers (due to magnetic defects) aggregating in clusters. The interfacial coupling (unidirectional in nature in certain situations), studied for the first time in these DMO systems, can be tuned by controlling the density and size of Fe oxide nanoparticles at the surface of larger TiO2 nanoparticles (e.g. via the processing route). This feature opens new possibilities for applications, as for example the magnetic actuation of different DMO systems with catalytic activity.27

 Aknowlegments The financial support through the Core Program PN16-480102 and the projects PNIII95/PED/2017 is highly acknowledged. Some of the authors (MS and IM) acknowledge the financial support in the frame of the Project PN 16 47 - LAPLAS IV.

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

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Figure 1. The XRD spectra of samples S0 to S12 (left side) and conventional TEM image of sample S12 with selected area electron diffraction evidencing the corresponding plane of anatase TiO2 (a=0.378 nm) 1388x637mm (96 x 96 DPI)


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Figure 2. STEM Dark-Field image of TiO2 nanopatricles in sample S6 with the selected frame for the spectrum image (upper-left figure). The spectrum image (with the involved sampling for EELS acquisition) for the 20 nm TiO2 nanoparticle with EELS patterns acquired in points (a) and (b) corresponding to the middle and the border of the nanoparticle (down figures). The EELS-spectrum imagining map of the investigated nanoparticle with red pixels corresponding to Ti and green pixels to Fe (upper-right figure). 100x97mm (300 x 300 DPI)



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Figure 3. The Mossbauer spectra of iron doped sample: a) S6 at 5 K, b) S6 at 50 K, c) S12 at 5 K, d) S12 at 50 K. 297x420mm (300 x 300 DPI)


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Figure 4. ZFC-FC behaviour of the investigated samples: a) S0, b) S6, c) S12. The cooling and measuring field was 100 Oe. 199x139mm (300 x 300 DPI)



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Figure 5. Hysteresis loops belonging to the magnetic ordered phases of samples S12 (left side) and S0 (right side), obtained after diamagnetic and paramagnetic corrections. Graphs a) and b) correspond to sample S12 after cooling in zero field and in 20 kOe, respectively, whereas graphs c) and d) correspond to sample S0, cooled in the same conditions. In the upper inset of each figure are shown the raw data whereas in the lower one, the central part of the loops is presented with a better field resolution. 199x139mm (300 x 300 DPI)


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Figure 6. Temperature evolutions of: magnetization at saturation (a), coercive field (b) and exchange bias field (c) of the analyzed samples cooled in different conditions. The relative errors, including the remanence of the superconducting magnet, of the values displayed on the graphs are less than 5 %. 286x199mm (300 x 300 DPI)



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TOC 203x152mm (300 x 300 DPI)


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Physical Mechanisms of Exchange Coupling Effects in

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