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Synthesis and characterization of Fe–Pt based multishell magnetic nanoparticles
Journal of Alloys and Compounds 574 (2013) 477–485
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Synthesis and characterization of Fe–Pt based multishell magnetic nanoparticles O. Pana a, C. Leostean a,⇑, M.L. Soran a, M. Stefan a, S. Macavei a, S. Gutoiu a, V. Pop b, O. Chauvet c a
National Institute for R&D of Isotopic and Molecular Technology, 65–103 Donath St., 400295 Cluj-Napoca, Romania ‘‘Babesß-Bolyai’’ University, Faculty of Physics, 1 Mihail Kogalniceanu St., 400084 Cluj-Napoca, Romania c Institut des Matériaux Jean Rouxel (IMN)-UMR 6502, Université de Nantes, CNRS, 2 rue de la Houssière, BP 32229, 44322 Nantes Cedex 3, France b
a r t i c l e
i n f o
Article history: Received 7 March 2013 Received in revised form 20 May 2013 Accepted 22 May 2013 Available online 31 May 2013 Keywords: Core–shell Nanoparticles Inverse micelles Iron/platinum Magnetization
a b s t r a c t Iron/platinum based core–shell nanoparticles were obtained by the inverse micelles method in two stages. Due to its specificity, this method produces an intermediate amorphous iron oxide layer. In this way the nanoparticles architecture is Fe@Fe-oxide@Pt. By thermal treatment in inert atmosphere, the amorphous shell crystallize and an additional magnetically ordered FePt alloy shell is formed at the interface between the Fe oxide and the outer Pt shell (Fe@Fe3O4/Fe2O3@FePt@Pt). The properties of these composites nanoparticles are investigated by TEM, HRTEM, X-ray diffraction (XRD), X-ray Photoelectron spectroscopy (XPS) and superconducting quantum interference device (SQUID) magnetization measurements. The thermally treated nanoparticles show one order of magnitude higher coercivities than the initial Fe@Fe-oxide@Pt nanoparticles while preserving almost the same saturation magnetization. The complex magnetic behavior of these both multilayer coupled magnetic materials as a function of temperature and applied magnetic field is also discussed. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Increased interest in magnetic nanoparticles was observed in the last decade by virtue of their various potential applications in fields ranging from ultrahigh-density recording and catalytic chemistry to biology and medicine [1–4]. Concerning various applications in biological systems it worth to mention sensing devices [5] cancer treatment by hyperthermia and drug delivery [6–10], and imaging in vivo [11,12]. The use of magnetic nanoparticles in biology and medicine demands the development of multifunctional nanoparticles having magnetic, optical and stability properties, as well as good functionalization capabilities. A convenient structure for multifunctional nanoplatforms is a core/shell type configuration of nanoparticles. Usually, as a specific architecture, the magnetic core is coated by an outer shell made either from inorganic and/or organic materials. This type of particles has drawn the attention of physicist, chemists and material scientists since they allow the tailoring of the combined shell and core properties [13]. Recently, there has been an increased interest in studying platinum nanoparticles and platinum based magnetic nanocomposites because of their ability to combine the simplicity of bioconjugation
⇑ Corresponding author. Tel.: +40 264 584037; fax: +40 264 420042. E-mail address: [email protected] (C. Leostean). 0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.05.153
and catalytic properties with high magnetic moments making them suitable for various applications [14–17]. Iron/platinum based core–shell nanoparticles (Fe@Pt) offer an attractive potential for adjustment of various properties. For instance the outer Pt shell can be directly provided with various functions or combined with an additional controllable gold shell then used for NIR adjustable laser ablation of targeted cells. An additional silica shell will provide, among others, the possibility to attach florescent markers or to be used in drug delivery [8,18]. Regarding other medical applications, the iron cores, having high saturation magnetizations, low remanences, and a superparamagnetic behavior, can be directly involved in hyperthermia, for cell destruction, as well as in NMR imaging. In order to improve the efficiency of the processes, these biomedical applications also require, if possible, an adjustable increased coercivity. In case of superparamagnetic nanoparticles there is a small, but non zero, coercive field. It can be increased by realizing core–shell structures formed by exchange coupled hard/soft magnetic cores and/or shells. This could be achieved by adjusting the sizes of various coupled magnetic phases with small but controlled values. In the present work one intends that, starting from the core– shell Fe@Fe-oxide@Pt nanoparticles and by using a proper thermal treatment, an additional magnetically ordered FePt alloy shell to be realized at the interface between the Fe inner oxide and the outer Pt shell (Fe@Fe3O4/Fe2O3@FePt@Pt). This way one expects to produce a magnetically coupled core multi-shell system with both in-
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creased coercivity and high saturation magnetization having high thermal energy transfer capability. Fe-oxide notation will be used with respect to amorphous iron oxide phase existing inside the as prepared sample. For crystallized oxides we will use the nominal formula. The properties of these composites nanoparticles are investigated by TEM, HRTEM, X-ray diffraction (XRD), X-ray Photoelectron spectroscopy (XPS) and magnetization measurements. 2. Experimental Core–shell Fe@Fe-oxide@Pt (FP) nanoparticles were obtained by the inverse micelles method by using FeSO4 and K2PtCl6 as starting reagents. Two stages were used: (i) synthesis of Fe0 and (ii) the realization of Pt0 within the ‘‘core–shell’’ structure. In the first stage there were prepared the Fe0 nanoparticles by the reduction of Fe2+ in the presence of NaBH4. It was done by mixing and stirring the FeSO4, the celtyltrimethylammonium bromide (CTAB) and NaBH4, solved in butanol and octane mixture under Ar atmosphere, at room temperature for one hour. In the second stage, the solution of K2PtCl6 was added to the system together with a new quantity of CTAB and NaBH4 in butanol and octane mixture. The reaction was done by stirring under Ar atmosphere for about 5 h. Here, the NaBH4 was acting as the reducing agent and CTAB as surfactant. The following molar ratios were used: surfactant/K2PtCl6 – 12.5%, and FeSO4/K2PtCl6 – 1.8%. In order to verify if an intermediate magnetic FePt alloy layer can be formed between the Fe core and the Pt shell the FP nanoparticles were thermally annealed in Ar atmosphere at a temperature of 650 °C for half an hour. This way the final product should be a multiple core–shell composite of Fe@Fe3O4/Fe2O3@FePtA@Pt type (abbreviated as FP-T). Here FePtA designates an ordered FePt alloy phase of either Fe3P, FePt (L10) or FePt3 form and Fe3O4/Fe2O3 represents one and/or other of the crystalline phases of iron oxide. The structural characterization of samples was made by X-ray diffraction (XRD). The data were recorded at room temperature by using a Bruker D8 Advance X-ray diffractometer set-up, at 40 kV and 40 mA with vertical powder Bragg–Brentano geometry. Ni filtered Cu Ka radiation, k = 1.54178 Å was used for diffraction. The morphology of the composite nanoparticles was determined by transmission electron microscopy (TEM) and high resolution TEM (HRTEM). The TEM measurements were performed with JEM JEOL 1010 microscope, with an accelerating voltage of 20 kV. The powder was dispersed in ethanol by using an ultrasonic bath then deposited on carbon film copper grid of 400 meshes. The HRTEM images were collected with Hitachi H9000NAR transmission electron microscope. X-ray Photoelectron spectroscopy (XPS) associated with Ar ion etching was used for qualitative and quantitative compositional analysis of nanocomposites, using a SPECS custom built system. Al anode was used for the X-rays source excitation (ht = 1486.6 eV). The magnetic behavior was recorded on dried samples by using a superconducting interference device (SQUID) magnetometer system, MPMS XL7 from Quantum Design.
3. Results and discussion In Fig. 1a it is presented a TEM image of an ensemble of core– shell FP nanoparticles while Fig. 1b presents the HRTEM image of the same nanoparticles. One can see the tendency of nanoparticles to form agglomerations as a result of the used surfactant (CTAB). The size distribution of nanoparticles determined from TEM and HRTEM images is shown in Fig. 2. It was realized by considering approximately 450 nanoparticles. The large sized particles as seen in Fig. 1, were not counted here. Continuous line in this graph represents the best fit performed using the ‘‘lognormal’’ distribution function. The calculated fit parameters are D0 = 3.14 nm and the dispersion r = 0.22 respectively. The qualitative compositional analysis of Fe@Pt sample was performed by energy-dispersive Xray spectroscopy (EDX) associated to HRTEM equipment. The HRTEM image of FP-T sample is shown in Fig. 3a. In Fig. 3b is presented the EDX spectrum corresponding to the small ensemble of agglomerated nanoparticles, marked by a circle (Fig. 3a), where both Fe and Pt various emission lines can be observed. The XRD patterns of FP nanoparticles together with both Fe and Pt line indexes are shown in Fig. 4. Indexing was made considering the bcc and ccp cubic structures for iron and platinum respectively. In order to avoid fluorescence effects due to Fe component, the structural determinations by X-ray diffraction were performed by using a Co Ka1 anode (1.78892 Å) instead of a Cu Ka1 one. Pt diffraction peak (1 1 1) was used to calculate the Pt lattice constant and to evaluate the average sizes of nanocrystallites by applying
Fig. 1. TEM (a) and HRTEM (b) images of core–shell nanoparticle ensembles of FP sample.
Fig. 2. Size distribution of the FP core–shell nanoparticles. A mean diameter of 3.14 nm is obtained by the best fit using the lognormal distribution function.
Scherrer method. The following values were determined for lattice constant and mean crystallite size: aPt = 3.9 Å and DPt = 5.7 nm respectively. The peak centered at 2h = 54.4° contains partially
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(a) I (a.u.)
150
Pt M 1 Pt M 2 Pt M
Fe K 1
Fe K 1
100
Cu K 1
Cu K 1
(b)
Pt L 1 Pt L 2 Pt L 1 Pt L 2
50
0
10 nm
250
500
750
1000
1250
E (eV)
Fig. 3. HRTEM image of FP-T nanoparticles (a) and the EDX spectrum corresponding to a small ensemble of agglomerated nanoparticles (b).
Fig. 4. XRD diffraction pattern of the FP core–shell nanoparticles. The corresponding diffraction line indexes for pure iron and platinum are also indicated.
overlapping diffraction lines of Pt (2 0 0) and Fe (1 1 0) and was deconvoluted by using pseudo Voigt line profiles. In this case the calculated lattice constants for iron and platinum were aPt = 3.88 Å and aFe = 2.79 Å. The corresponding nanocrystallite sizes are DPt = 5.6 nm and DFe = 4.6 nm respectively. These values are a little higher than the mean size determined from HRTEM. The difference arises from large sized nanoparticles (D > 20 nm) as seen in Fig. 1a. The large nanoparticles, even few in number, give a significant contribution to the diffraction XRD line narrowing. For both diffraction peak analyzed, the lattice constants are slightly different from standard ones namely a = 3.92 Å for Pt and a = 2.86 Å for Fe. As observed in case of Fe@Fe-oxide@Au core shell nanoparticles, prepared also by the reverse micelles method [19] the intermediate shell of oxidized Fe should be mostly amorphous, thus is ‘‘unseen’’ by XRD. The XRD patterns of thermally treated FP-T sample are presented in Fig. 5. By annealing of the FP sample some of the iron atoms from the oxide shell will diffuse into the neighboring Pt shell with the formation, at the interface, of an ordered FePt alloy. In the same time, as a result of thermal treatment, the amorphous oxide shell will be crystallized into an ordered oxides shell. This can be seen in Fig. 5 where FePt alloy (L10 phase), Fe3O4, and Fe2O3 hematite diffraction patterns are present. In this case the iron core contribution is too small to be seen. Due to the thermal diffusion of iron into platinum shell, accompanied by the formation of the FePt alloy, the remaining oxide in the vicinity of Pt becomes depleted in Fe, thus hematite rather than Fe3O4 is most expected here. The lines are narrower than in case of FP sample indicating that, as a result of coalescence, some larger nanoparticles are produced in the system (20–30 nm). Due to the multiple peaks seen in Fig. 5 and to the noisy background an accurate determination by Scher-
Fig. 5. XRD diffraction pattern of the thermally treated FP-T core–multishell nanoparticles. The corresponding diffraction line indexes for FePt alloy L10, Fe3O4 and Fe2O3 hematite are also indicated.
rer equation is uncertain. The large nanoparticles, even if they are few in numbers, give an important contribution to the diffraction patterns since they contain an enough large fraction of the total number of atoms. For further compositional analysis of FP and FP-T nanocomposites XPS associated with Ar ions etching was used. The samples were subject of consecutive Ar ions etchings until by doing any additional sputtering the XPS spectra remained unchanged in shape and intensity. At this stage they reflect the real composition of the core–shell nanoparticles. Fe 2p, O 1s, C 1s, K 2p, Cl 4s as well as Pt 4d and Pt 4f core-levels lines were identified and can be seen in XPS survey spectrum of FP sample presented in Fig. 6. The existence of potassium and chlorine lines indicates that non-reacted
Fig. 6. XPS survey spectrum of FP sample.
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small quantities of K2PtCl6 are still there while C 1s core-level line is due to the remaining CTAB surfactant. For quantitative determinations the individual core-level lines were deconvoluted and fitted. As an example, the XPS spectra of Fe 2p as well as of Pt 4d core-level doublets from FP sample, together with the corresponding deconvolutions and fitted curves are shown in Fig. 7. As expected, besides the Fe(0) peaks two additional doublets corresponding to oxidized iron were seen at higher binding energies [20]. Depending on the synthesis conditions the oxide shell is more or less thick. In our case the best fit for the Fe 2p lines was obtained considering Fe2+ and Fe3+ oxidation states. The Fe 2p main lines are positioned at: 708.4 eV (2p3/2), 721.6 eV (2p1/2) for Fe(0) (denoted A), 709.5 eV (2p3/2), 722.9 eV (2p1/2) for Fe2+ (denoted B) and 711.9 eV (2p3/2), 725.3 eV (2p1/2) for Fe3+ (denoted C). Two pairs of iron shake-up satellite features at 714.9, 728.5 eV and 718.1, 732.2 eV are shown. The first one incorporates shake – up features associated to the Fe(0) and Fe2+ main lines and the second is associated to Fe3+. In this way we wanted to avoid placing too many lines in the spectrum. The Fe 2p doublet separation was constrained between 13.2–13.4 eV. Pt 4s core-level line which is superimposed to the Fe 2p spectrum is evidenced at 723.0 eV in Fig. 7a. In case of Pt 4d core-level spectrum, as shown in Fig. 7b, three doublet peaks can be seen: 312.9 eV (4d5/2), 330.0 eV (4d3/2) associated to the main platinum shell (label A), 309.1 eV (4d5/2), 326.5 eV (4d3/2) for platinum surface states (label B) and, at higher binding energies, 316.3 eV (4d5/2), 333.4 eV (4d3/2) are the more positive features belonging Pt ions surrounded by 6 negative chlorine from the K2PtCl6 residual precursor (label C). The quantity of platinum states associated to residual K2PtCl6 is in agreement with the potassium amount as determined from K 2p core-level doublet lines. The Pt 4d doublet separation was set between 17.0 and 17.1 eV.
(a)
(b)
For p, d and f core-levels lines the restrictions used for the fit core-level XPS spectra refer to the usual relation between areas of spin–orbit coupling doublet components. By using CASA software data base the integral intensities were normalized by dividing the raw XPS deconvoluted intensities to the corresponding real sensitivity, transmission and electronic mean free path factors. It was observed that, due to the core–shell structuring of nanoparticles, by increasing the etching time the XPS intensities of the Pt outer shell lines decrease while the line intensities of inner core Fe increase until the XPS signal, as coming from a very large number of sectioned nanoparticles, became stabilized [21]. As it was mentioned, XRD does not show any specific diffraction patterns of crystalline oxides thus the oxidized iron should to be mostly into an amorphous phase. It results that the real architecture of FP nanoparticles is, as expected, Fe@Fe-oxide@Pt. Molar and weight concentrations were evaluated dividing the integral intensities by the attenuation lengths (escape depths), kAL, specific for each component of the composite system [22]. They reflect the mean depth from which emerged photoelectrons are collected into the energy analyzer. The values of kAL for a given component depend on the material density, molar weight, and mean Z number of the compound as well as on the kinetic energy of the emerged photoelectrons [23]. The values of the attenuation lengths, calculated upon Ref. [23] are as following: 0.82 and 0.9 nm for Fe(0) and Fe3O4 respectively as one refers at Fe 2p core-level, 1.83 nm regarding K2PtCl6 at K 2p core-level line, 2.14 nm for CTAB at C 1s core-level line, and 0.85 nm for Pt 4d line. The weight content of different components from samples, xi, was evaluated by using the following relation:
li kIii Ii i li ki
Xi ¼ P
ð1Þ
Here Ii represents the integral intensity of the XPS core-level peak for a specific element of a molecule used for calculation and li the corresponding molar mass. The resulted molar and weight contents of Fe(0), Fe oxide, Pt and small quantities of as mentioned non-reacted components are presented in Table 1. The radial dimensions were calculated by considering a mean diameter of 3.14 nm as determined from size distribution of nanoparticles resulted from TEM. Taking into consideration both mean size (Fig. 2) and the composition of FP nanoparticles (Table 1) the mean volumes of the three components of the system were calculated as: 5.24, 7.87, and 3.1 nm3 for Fe core, amorphous iron oxide sell, and Pt shell respectively. The small value resulted for Pt outer shell thickness represents an average between small and large particles having different platinum coverage levels. This inhomogeneity in platinum coverage is a result of the reversed micelles method used for the preparation. One specificity of the method is that the larger the water micelle formed by surfactants greater the size of the resulted nanoparticles. Besides, the micellar surfactants have a decisive role in the platinum reduction. They act as growth-directing adsorbents on metal surfaces. As the size of the water droplet increases the micellar concentration decreases thus reducing the rate of metal reduction inside the confined hydrophilic volume. Due to higher concentration of hydrophilic terminations in small micelles, the small particles will be
Table 1 The calculated molar and weight content values of the FP sample as resulted from XPS.
Fig. 7. XPS spectra of Fe 2p (a) as well as of Pt 4d (b) core-level doublets from FP sample, together with the corresponding deconvolutions and fitted curves.
FP sample
Fe(0)
Feoxide
Pt
K2PtCl6
CTAB
yi (mol.%) xi (wt.%) Shell thickness (nm)
56.7 23.9 1.08 (radius)
13.5 23.7 0.38
28.0 38.6 0.11
3.3 12.3 –
0.5 1.5 –
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favored during the reduction process and much thick and complete outer Pt shell will result [24–26]. For larger nanoparticles a rather incomplete or cracked external shell is most probable. As it can be seen in the inset of Fig. 1b the interplanar distance evidenced by0 HRTEM in case of uniform small particles has the value of 2.2 Å A which corresponds to (1 1 1) crystalline plane family of Pt. Well defined Pt crystallites are also indicated by the XRD pattern from Fig. 4. The same XPS lines were also observed in case of FP-T sample. For instance, the corresponding Fe 2p and Pt 4d features of FP-T are presented in Fig. 8. The oxidized iron features are positioned at 709.2 eV (2p3/2) and 724.4 eV (2p1/2) for Fe2+ and at 711.1 eV (2p3/2) and 724.4 eV (2p1/2) for Fe3+ respectively. To avoid having too many peaks in the deconvolution, no separate peak doublets were set for each of the Fe3O4 and Fe2O3 components. The area of Fe3+ feature was left to evolve free with respect to the Fe2+ one in order to incorporate both Fe3O4 and Fe2O3 contributions. A more intense Fe3+ peak as compared to the Fe2+ one resulted accordingly. The Fe(0) from the inner core and Fe atoms incorporated into the alloy phase have binding energies of 708.4, 707.6 eV for 2p3/2 and 721.7, 720.8 eV for 2p1/2 respectively. Again, to avoid the use of too many components in the deconvolution the only one shake-up satellite peak was considered for every doublet component. Each satellite incorporates shake-up contributions from different principal lines. The Pt 4d core-level XPS spectrum of FP-T sample is presented in Fig. 8b. Here A and B doublet features denote the platinum atoms from FePt (L10) alloy and platinum outer shell, respectively. Additionally, an intense Ar 2s feature and a plasmonic peak are also observed. Due to the uncertainty regarding the type of potassium, chlorine and carbon end products resulted from the annealing of FP nanoparticles in the presence of highly catalytic nanostructured platinum as well as due to the presence of adsorbed/absorbed
(a)
(b)
Fig. 8. XPS spectra of Fe 2p (a) and Pt 4d (b) core-levels from FP-T sample. The corresponding deconvolutions and fitted curves are also presented.
residual Ar (from Ar atmosphere inside the kiln) a quantitative analysis, similar to that determined in case of FP sample, was no longer possible for FP-T nanoparticles. Nevertheless, considerations relative to the Fe(0), Fe3O4/Fe2O3, FePt (L10), and Pt from FP-T multi-shell sample can be made in correlation with the composition of FP nanoparticles. Thus, from the ratio of A and B peaks in Fig. 8b, it results that 33% from total Pt atoms were included into FePt alloy phase. Also it is reasonable to make the assertion that the Fe(0) core remains with the same volume as in the initial FP sample. An estimation regarding the relative content of the two iron oxides is rather difficult to be made. The magnetic properties of the FP nanoparticles were examined and compared with properties of the thermally treated FT-T nanoparticles. The magnetizations of FP compound as a function of the applied field H, at room temperature (RT) and at 4.2 K, are presented in Fig. 9a. The nanocomposite has with a quite small coercive field of 29 Oe at RT and, due to the blocking process, the coercive field increases to 282 Oe at 4.2 K. In case of nanoparticle ensembles, small coercive fields represents an indication that superparamagnetic behavior may be present in the system. The saturation magnetization of the iron core (Ms), normalized to the specific Fe weight content as resulted from XPS analysis (Table 1), is about 179 emu/g(Fe), which represents a typical value for iron nanoparticles [27]. The magnetic behavior changes after thermal annealing, as seen in Fig. 9b. When the temperature decreases from 298 K to 4.2 K, the coercive field increases and the saturation magnetization remain at almost the same level as for the FP sample. It is a prime indication that some of the magnetic components are magnetically coupled. The first derivative of demagnetization/magnetization curves of FP sample compared with the same first derivative FP-T sample at RT and 4.2 K are presented in Fig. 10. Since they are symmetric we
(a)
(b)
Fig. 9. Magnetization vs. applied magnetic field, taken at room temperature and 4.2 K, corresponding to: FP nanoparticles (a) and thermally treated FP-T nanoparticles (b). The coercive fields are shown in the insets.
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will refer only to the derivative of demagnetizing curves. The first derivative of demagnetizing curve shows the local slope of the M– H dependences and thus the number of peaks or bumps present in the dM/dH vs. H curve reveals the number of magnetic phases within the sample. The horizontal shift of these peaks from their nominal coercive field positions (for pure phases) is directly related to the strengths of the magnetic phases coupling. When hard and soft magnetic phased are exchange coupled the peaks or shoulders associated with hard magnetic phases are shifted from their nominal coercivities towards zero field while the peaks of the soft phase components are moved to higher fields with respect to their usual (uncoupled) coercive field positions. In case of fully coupled systems the two peaks can merge into a single broader one [28]. In case of FP-T sample the presence of the inflection points, at both RT and 4.2 K, it is an indication that several magnetically coupled phases are present into the system. The assignment of different contributions is based on the intrinsic coercivities expected from each component as (I) coupled ferromagnetic (FM) Fe(0), (II) coupled ferrimagnetic Fe3O4, (III) coupled FM FePt, and (IV) non-coupled FM FePt. They are indicated by arrows in Fig. 10. The last attributed position is based on the usual values of coercive field [29]. The so called DM plots were also calculated as in Ref. [30]. DM plots represent the difference between the ascending and descending parts of the hysteresis loop for B > 0. Like the first derivative of the demagnetization (magnetization) curves the derivative of these plots, dM/dB, reveal, with a more accuracy, mixed coercivity fractions. One maximum (or bump) means one coercivity fraction in the loop. The superimposed plots of FP and FP-T samples, obtained from RT and 4.2 K hysteresis loops, are presented In Fig. 11. One can see that the magnetic couplings appear both at RT and 4.2 K for the thermally treated sample. Fig. 11a shows an additional inflection point of the FP-T sample that is connected
(a)
(b)
Fig. 11. DM plot for the FP and FP-T nanoparticles obtained from room temperature (a) and 4.2 K (b) magnetic hysteresis measurements.
to the coupled week FM Fe2O3. Below 260 K hematite Fe2O3 it’s an antiferromagnet, hence this feature is not longer observed in Fig. 11b. In case of the FP sample in 4.2 K DM plot a weak coupling is observed at fields related to the coupled ferrimagnetic Fe3O4. This feature it is prove of the sensitivity of the DM plot method. One has to mention that XRD measurements didn’t revealed any crystallized form with regard to the Fe oxides inside the FP sample. The temperature dependences of the mean coercive field HC, defined as HC ¼ ðHleft Hright Þ=2, for both FP and FP-T samples are presented in Fig. 12. One can see that down to about 50 K the FP-T sample has much larger coercivities than initial FP sample. Bellow 50 K, due to both blocking process and energy density anisotropy increase, the coercive field of FP sample starts to rise. In case of FP-T sample the coercive field starts rising at about 110 K then, after reaching a maximum around 25 K, it begins to decrease. This maximum of the coercive field can be related to a max-
(b)
Fig. 10. Comparison of the magnetization derivative maxima corresponding to the FP-T nanoparticles at room temperature (a) and 4.2 K (b).
Fig. 12. Coercive field temperature dependence corresponding to the FP nanoparticles and to the thermally treated FP-T nanoparticles.
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imum of the effective anisotropy density of the composite system. It can be also an indication that some of the magnetic components of FP-T sample are exchange coupled. At more detailed analysis one can see that HC has slightly different values in case of the magnetization curve as compared to the demagnetization one. In case of FP sample at 298 K HC is 24 Oe for the magnetization process and 34 Oe for the demagnetization curve. At 4.2 K the corresponding values are 272 and 293 Oe respectively. It is an indication that small quantities of crystallized ordered iron oxide are present even in the initial FP sample thus a small exchange is possible to occur here (as already evidenced by DM plot). In case of FP-T sample, as compared to the initial sample, a more intense exchange coupling appears here. At RT and 4.2 K the demagnetization–magnetization left–right coercive fields are 131–111 Oe and 534–491 Oe, respectively. The values of HE and HC are determined from the two fields (Hleft and Hright) of the hysteresis loop where the magnetization goes through zero. The following relation was used for the exchange field: HE ¼ ðHleft Hright Þ=2. Thus, a hysteresis loop which is shifted to the left of H = 0 is defined as having a positive exchange field [31]. The evolution of the exchange field, HE for the whole temperature range together with the temperature dependence of the field position of Fe(0) maximum (I) in the derivative of the demagnetization curve, as seen in Fig. 10, is presented in Fig. 13. Iron maximum position displacement to the left side with respect to H = 0 (derivative of the demagnetization curve) it is a measure of the coupling strength so that larger the displacement more intense the coupling with the neighbor magnetite shell. As one can see the Fe(0) maximum displacement follow the general tendency of the exchange field HE behavior as function of T. Both fields have a minimum at around 260 K which corresponds to the Morin transition of Fe2O3 component. Above this temperature the bulk hematite is a week ferromagnet (WF) while below it becomes an antiferromagnet (AF). For T < 260 K the hematite nanoparticles it may or may not exhibit a spin-flop Morin transition between WF and AF states, depending on its stoichiometry, crystallinity, surface condition, and microstructure. Generally particle size does not correlate directly with the presence or absence of a Morin transition but instead, that the Morin transition is controlled by the precise lattice parameters. The annealing of samples at temperatures above 550 °C, as it is our case, resulted in a reduction of defects and an increase of TM up to 260 K [32–34]. Once the Fe2O3 AFM phase is set, an exchange bias may be established with its ferromagnetic and ferrimagnetic neighbors, the outer FePt shell and the inner Fe3O4 one, respectively. As a consequence the overall exchange field starts to increase and, through the oxides layer (hematite and magnetite), a coupling, going from
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the outer FePt shell down to the inner iron core, can be realized. In the vicinity of 120 K the Fe3O4 component undergoes the Verwey transition accompanied by a peak in the temperature dependence of the exchange field HE. At about 130 K the Verwey transition is preceded by the so-called magnetic isotropic point (Ti) [35]. One has to remember that the iron ions, in magnetite, are distributed between two sublattices. The sublattice A has Fe3+ occupying tetrahedral positions and sublattice B with equally distributed Fe2+ and Fe3+ in octahedral positions. The magnetic moments of Fe3+ ions from A and B sublattices are antiferromagnetically coupled. Within B sublattice Fe2+ and Fe3+ magnetic moments are ferromagnetically coupled [36]. It should be mentioned that below the Verwey transition temperature, the magnetite changes its crystal structure from inverse cubic spinel to monoclinic symmetry. The process is accompanied reorientation of the easy magnetization axis along c direction. The interplay between crystal structure, magnetism and conductivity associated to the Verwey transition are still under debate [37]. Since the dimensions of the FM components are at the nanometer scale they behave like monodomanins and are subject of thermally activated magnetization fluctuations. It appears that during the Verwey transition thermal fluctuations are reduced and the exchange field increases. At about 34.7 K both fields presented in Fig. 13 reach a local maximum. The meaning of these features can be understood in connection with the distribution of the energy barriers as determined from field cooled (FC) and zero field cooled (ZFC) temperature dependences of magnetizations. The magnetization versus temperature in ZFC–FC regimes were measured in 100 Oe magnetic field. In Fig. 14 there are represented the temperature dependence of the heights of the energy barriers for FP sample (a) and FP-T sample (b). They were deduced in accordance with Refs.
(a)
(b)
Fig. 13. The evolution of the exchange field, HE for the whole temperature range together with the temperature dependence of the field position of Fe(0) maximum (I) in the derivative plot of the demagnetization curve, as seen in Fig. 10.
Fig. 14. The energy barriers distributions that separates the two orientations of the nanoparticles magnetization corresponding to the FP (a) and thermally treated FP-T nanoparticles (b). They were calculated from FC and ZFC M = f(T) dependences shown in the insets.
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[19,38]. The FC and ZFC temperature dependences for the two samples are shown in the corresponding insets of Fig. 14a and b. One can see that a maximum appears in the ZFC curve from the inset of Fig. 14a at a temperature of about 110 K. It represents the mean blocking temperature TB of the magnetic nanoparticle system. Considering the small size of iron cores, the existence of a small coercive field, as indicated above in Fig. 9a, and the presence of a mean blocking temperature TB one can confirm that FP core–shell has a superparamagnetic character [19,21,22,38,39]. The temperature dependences energy barrier heights are connected to the Keff V product as:
K eff V ¼ kBT ln
s s0
ð2Þ
Here Keff represents the effective anisotropy energy density, V is the magnetically ordered volume of the nanoparticles and s and s0 are the observed and characteristic magnetization switching times respectively. The graph maximum from Fig. 14a, positioned at 17 K, corresponding to the energy barriers distribution of Fe(0) cores from FP sample [24] can be used to determine the effective anisotropy energy density constant K Fe eff [38]. Assuming an iron core volume of 5.24 nm3 as determined from XPS analysis a value of 6 3 K Fe eff 1:1 10 J=m results. This value is consistent with values determined for small sized Fe nanoparticles [24]. The energy barrier distribution is completely different in case of FP-T sample where two maxima are present into the graph. The intense maximum at 113 K is connected to Fe3O4 component and corresponds to the Verwey transition preceded at around 130 K by the inversion point of magnetite crystalline anisotropy [35]. The maximum at Tm = 44K corresponds to energy barriers distribution of iron cores, shifted, in FP-T sample, to higher T due to an increased K Fe eff value. The shift of the Fe(0) contribution can be attributed to the exchange existing between the iron cores and the newly formed magnetite intermediate shell. It represents the sum of several contributions: Fe K Fe eff ¼ K ð0Þ þ
6 ðK ðSÞ þ K ðSÞex Þ þ K ex : DFe
ð3Þ
Here K Fe ð0Þ is the axial intrinsic iron anisotropy energy density and Kex represents the exchange contribution. K(S) and K(S)ex are the intrinsic surface anisotropy and the exchange contribution to the surface anisotropy, respectively. Due to the roughness and inhomogeneities at the Fe(0)/Fe3O4 interface the exchange couplings are quite complex [37]. However, one can estimate that the coupling of surface Fe magnetic moments with Fe3+ occupying tetrahedral positions of the sublattice A or octahedral sites of sublattice B will have an exchange bias character thus contributing to K(S)ex. The coupling with Fe2+ states of B sublattice involves a direct exchange coupling and contributes to Kex. Also superexchange involving O 2p spin paired orbitals is possible and probably contributes to Kex. A quantitative estimation of various contributions seems to be difficult to be made at the present stage. At a detailed analysis, the maximum seen in Fig. 13 at T 0m ¼ 34:7 K (marked by an arrow) represents the same blocking process of the iron cores as observed at Tm = 44 K in energy barriers distribution from Fig. 14b. Due to the temperatures steps in recording different hysteresis loops its real position should be probably in-between the two consecutive points, namely 34:7 K T 0m 49:7 K. Further, considering the maximum associated with Fe(0) core from Fig. 14b, this temperature range can be reduced to 34:7 K T 0m T m . On the other hand, the graphs presented in Fig. 14 were deduced from thermo remnant magnetizations following [38,40], and as a consequence Eq. (2), which stands for H = 0, applies here. On the other hand the temperature dependence
of the exchange field, HE, was calculated from hysteresis loops taken at different temperatures, and thus Eq. (2) is no longer valid. Instead the following relation applies here [40]:
a H s K eff V 1 ¼ kB T ln HC s0
ð4Þ
Here HC is the field at which the given barrier vanishes. The exponent a takes values between 1.5 and22 [40]. For the a = 2 case and, if 0 < H/HC < 2, the factor 1 HHC act towards reducing the height of the energy barrier, thus decreasing the temperatures at which magnetization thermal fluctuations are blocked. At zero magnetization HE is the local field acting on Fe(0) cores while the field which cancels the magnetization is HC + HE, as recorded in the demagnetization curve. As a consequence the temperature associated to the maximum for which most of the iron cores are blocked is shifted to lower values as in Fig. 13. By considering Eqs. (2) and (3) the ratio between T 0m =T m became
2 T 0m HE ¼ 1 Tm HC þ HE
ð5Þ
Considering the corresponding values of HE and HC one obtains for T 0m ¼ 41:2 K which is a temperature within the expected range. The incompletely resolved maximum at low temperatures as seen in Fig. 14b can be attributed to the blocking of FePt thin shell. This assertion is sustained by calculating the blocking temperature of FePt alloy. As resulting from XPS analysis of spectrum shown in Fig. 8b it results that about one third of the initial platinum metal turns into FePt L10 alloy and thus its estimated mean volume is 6 3 about 1 nm3. Considering K FePt eff 6—7 10 J=m [29] one obtains FePt FePt a value of T B ¼ 16:1—18:7 K depending on T eff value and the value chosen for ln ss0 namely 25–27. At temperatures below these values the exchange field HE has a sharp increase due to the blocking of FePt alloy shell. Concerning the magnetite, below the Verwey transition temperature and, due to the neighbor couplings (high in3 O4 crease of K Fe ), it can be considered as being in a blocked state. eff Thus, below 16.1–18.7 K temperatures range the magnetization of one particle behaves according to a ‘‘superspin’’ configuration. In other words, the whole magnetization rotates at once within all coupled shells. As a consequence, even if HE increases with decreasing temperature (Fig. 13), HC field, which expresses the ‘‘opposition’’ to the magnetic field reversal, decreases due to the lack of intra-shell local fluctuations. This behavior is evidenced by the maximum seen, at low temperatures in Fig. 12 for FP-T sample. 4. Conclusions The reverse micelles method was used for preparation of Fe/Pt ‘‘core–shell’’ nanoparticles. Due to the specificity of method the Fe core is covered by a layer of amorphous iron oxide, so that the actual structure of nanoparticles with average size of about 3 nm is Fe@Fe-oxide@Pt. Starting from this type of particles by heat treatment in inert atmosphere (Ar) a complex system of ‘‘coremultishell’’ structure Fe@Fe3O4@Fe2O3@FePt@Pt was obtained. The magnetic properties of these two systems were analyzed and compared. The thermally treated nanoparticles show one order of magnitude higher coercivities than the initial Fe@Fe-oxide@Pt nanoparticles, preserving in the same time, at least at RT, an almost superparamagnetic behavior. At detailed analysis of hysteresis loops of annealed nanoparticles shows that HC has slightly different values for the magnetization curves as compared to the demagnetization ones thus proving the existence magnetically coupled multiple shells. The Verwey transition was evidenced at 113 K in case of Fe3O4 shell. An exchange bias seems to exists, bellow 260 K, at the inter-
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faces of Fe2O3 shell as well as between some of the surface magnetic moments of iron core and the Fe3+ moments from both A and B sublattices of the magnetite shell. Due to these couplings the Fe(0) cores of annealed nanoparticles have an increased effective anisotropy constant and became almost blocked at higher temperatures than in case of initial Fe@Fe-oxide@Pt nanoparticles (44 K instead of 17 K). Generally, when decreasing temperature, the evolution of the exchange field shows an increase, with two local maxima. One maximum is at the Verwey transition temperature and another one at a temperature corresponding to the blocking of the Fe cores. Thus the effective anisotropy constant of the iron core seems to be the sum of volume axial anisotropies and surface anisotropies originating from intrinsic and exchange contributions. Due to the complexity of the nanoparticles resulted from thermal treatment, at the present stage, a quantitative estimations for these contributions seems to be difficult to achieve. Acknowledgement This work was supported by the Romanian Ministry of Education and Research under the research Program POS-CCE-METAVASINT – 550/2010 and ID – 119/2011. Appendix A. Supplementary data Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j.jallcom.2013.05. 153. References [1] [2] [3] [4] [5] [6] [7]
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Contents lists available at SciVerse ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Synthesis and characterization of Fe–Pt based multishell magnetic nanoparticles O. Pana a, C. Leostean a,⇑, M.L. Soran a, M. Stefan a, S. Macavei a, S. Gutoiu a, V. Pop b, O. Chauvet c a
National Institute for R&D of Isotopic and Molecular Technology, 65–103 Donath St., 400295 Cluj-Napoca, Romania ‘‘Babesß-Bolyai’’ University, Faculty of Physics, 1 Mihail Kogalniceanu St., 400084 Cluj-Napoca, Romania c Institut des Matériaux Jean Rouxel (IMN)-UMR 6502, Université de Nantes, CNRS, 2 rue de la Houssière, BP 32229, 44322 Nantes Cedex 3, France b
a r t i c l e
i n f o
Article history: Received 7 March 2013 Received in revised form 20 May 2013 Accepted 22 May 2013 Available online 31 May 2013 Keywords: Core–shell Nanoparticles Inverse micelles Iron/platinum Magnetization
a b s t r a c t Iron/platinum based core–shell nanoparticles were obtained by the inverse micelles method in two stages. Due to its specificity, this method produces an intermediate amorphous iron oxide layer. In this way the nanoparticles architecture is Fe@Fe-oxide@Pt. By thermal treatment in inert atmosphere, the amorphous shell crystallize and an additional magnetically ordered FePt alloy shell is formed at the interface between the Fe oxide and the outer Pt shell (Fe@Fe3O4/Fe2O3@FePt@Pt). The properties of these composites nanoparticles are investigated by TEM, HRTEM, X-ray diffraction (XRD), X-ray Photoelectron spectroscopy (XPS) and superconducting quantum interference device (SQUID) magnetization measurements. The thermally treated nanoparticles show one order of magnitude higher coercivities than the initial Fe@Fe-oxide@Pt nanoparticles while preserving almost the same saturation magnetization. The complex magnetic behavior of these both multilayer coupled magnetic materials as a function of temperature and applied magnetic field is also discussed. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Increased interest in magnetic nanoparticles was observed in the last decade by virtue of their various potential applications in fields ranging from ultrahigh-density recording and catalytic chemistry to biology and medicine [1–4]. Concerning various applications in biological systems it worth to mention sensing devices [5] cancer treatment by hyperthermia and drug delivery [6–10], and imaging in vivo [11,12]. The use of magnetic nanoparticles in biology and medicine demands the development of multifunctional nanoparticles having magnetic, optical and stability properties, as well as good functionalization capabilities. A convenient structure for multifunctional nanoplatforms is a core/shell type configuration of nanoparticles. Usually, as a specific architecture, the magnetic core is coated by an outer shell made either from inorganic and/or organic materials. This type of particles has drawn the attention of physicist, chemists and material scientists since they allow the tailoring of the combined shell and core properties [13]. Recently, there has been an increased interest in studying platinum nanoparticles and platinum based magnetic nanocomposites because of their ability to combine the simplicity of bioconjugation
⇑ Corresponding author. Tel.: +40 264 584037; fax: +40 264 420042. E-mail address: [email protected] (C. Leostean). 0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.05.153
and catalytic properties with high magnetic moments making them suitable for various applications [14–17]. Iron/platinum based core–shell nanoparticles (Fe@Pt) offer an attractive potential for adjustment of various properties. For instance the outer Pt shell can be directly provided with various functions or combined with an additional controllable gold shell then used for NIR adjustable laser ablation of targeted cells. An additional silica shell will provide, among others, the possibility to attach florescent markers or to be used in drug delivery [8,18]. Regarding other medical applications, the iron cores, having high saturation magnetizations, low remanences, and a superparamagnetic behavior, can be directly involved in hyperthermia, for cell destruction, as well as in NMR imaging. In order to improve the efficiency of the processes, these biomedical applications also require, if possible, an adjustable increased coercivity. In case of superparamagnetic nanoparticles there is a small, but non zero, coercive field. It can be increased by realizing core–shell structures formed by exchange coupled hard/soft magnetic cores and/or shells. This could be achieved by adjusting the sizes of various coupled magnetic phases with small but controlled values. In the present work one intends that, starting from the core– shell Fe@Fe-oxide@Pt nanoparticles and by using a proper thermal treatment, an additional magnetically ordered FePt alloy shell to be realized at the interface between the Fe inner oxide and the outer Pt shell (Fe@Fe3O4/Fe2O3@FePt@Pt). This way one expects to produce a magnetically coupled core multi-shell system with both in-
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creased coercivity and high saturation magnetization having high thermal energy transfer capability. Fe-oxide notation will be used with respect to amorphous iron oxide phase existing inside the as prepared sample. For crystallized oxides we will use the nominal formula. The properties of these composites nanoparticles are investigated by TEM, HRTEM, X-ray diffraction (XRD), X-ray Photoelectron spectroscopy (XPS) and magnetization measurements. 2. Experimental Core–shell Fe@Fe-oxide@Pt (FP) nanoparticles were obtained by the inverse micelles method by using FeSO4 and K2PtCl6 as starting reagents. Two stages were used: (i) synthesis of Fe0 and (ii) the realization of Pt0 within the ‘‘core–shell’’ structure. In the first stage there were prepared the Fe0 nanoparticles by the reduction of Fe2+ in the presence of NaBH4. It was done by mixing and stirring the FeSO4, the celtyltrimethylammonium bromide (CTAB) and NaBH4, solved in butanol and octane mixture under Ar atmosphere, at room temperature for one hour. In the second stage, the solution of K2PtCl6 was added to the system together with a new quantity of CTAB and NaBH4 in butanol and octane mixture. The reaction was done by stirring under Ar atmosphere for about 5 h. Here, the NaBH4 was acting as the reducing agent and CTAB as surfactant. The following molar ratios were used: surfactant/K2PtCl6 – 12.5%, and FeSO4/K2PtCl6 – 1.8%. In order to verify if an intermediate magnetic FePt alloy layer can be formed between the Fe core and the Pt shell the FP nanoparticles were thermally annealed in Ar atmosphere at a temperature of 650 °C for half an hour. This way the final product should be a multiple core–shell composite of Fe@Fe3O4/Fe2O3@FePtA@Pt type (abbreviated as FP-T). Here FePtA designates an ordered FePt alloy phase of either Fe3P, FePt (L10) or FePt3 form and Fe3O4/Fe2O3 represents one and/or other of the crystalline phases of iron oxide. The structural characterization of samples was made by X-ray diffraction (XRD). The data were recorded at room temperature by using a Bruker D8 Advance X-ray diffractometer set-up, at 40 kV and 40 mA with vertical powder Bragg–Brentano geometry. Ni filtered Cu Ka radiation, k = 1.54178 Å was used for diffraction. The morphology of the composite nanoparticles was determined by transmission electron microscopy (TEM) and high resolution TEM (HRTEM). The TEM measurements were performed with JEM JEOL 1010 microscope, with an accelerating voltage of 20 kV. The powder was dispersed in ethanol by using an ultrasonic bath then deposited on carbon film copper grid of 400 meshes. The HRTEM images were collected with Hitachi H9000NAR transmission electron microscope. X-ray Photoelectron spectroscopy (XPS) associated with Ar ion etching was used for qualitative and quantitative compositional analysis of nanocomposites, using a SPECS custom built system. Al anode was used for the X-rays source excitation (ht = 1486.6 eV). The magnetic behavior was recorded on dried samples by using a superconducting interference device (SQUID) magnetometer system, MPMS XL7 from Quantum Design.
3. Results and discussion In Fig. 1a it is presented a TEM image of an ensemble of core– shell FP nanoparticles while Fig. 1b presents the HRTEM image of the same nanoparticles. One can see the tendency of nanoparticles to form agglomerations as a result of the used surfactant (CTAB). The size distribution of nanoparticles determined from TEM and HRTEM images is shown in Fig. 2. It was realized by considering approximately 450 nanoparticles. The large sized particles as seen in Fig. 1, were not counted here. Continuous line in this graph represents the best fit performed using the ‘‘lognormal’’ distribution function. The calculated fit parameters are D0 = 3.14 nm and the dispersion r = 0.22 respectively. The qualitative compositional analysis of Fe@Pt sample was performed by energy-dispersive Xray spectroscopy (EDX) associated to HRTEM equipment. The HRTEM image of FP-T sample is shown in Fig. 3a. In Fig. 3b is presented the EDX spectrum corresponding to the small ensemble of agglomerated nanoparticles, marked by a circle (Fig. 3a), where both Fe and Pt various emission lines can be observed. The XRD patterns of FP nanoparticles together with both Fe and Pt line indexes are shown in Fig. 4. Indexing was made considering the bcc and ccp cubic structures for iron and platinum respectively. In order to avoid fluorescence effects due to Fe component, the structural determinations by X-ray diffraction were performed by using a Co Ka1 anode (1.78892 Å) instead of a Cu Ka1 one. Pt diffraction peak (1 1 1) was used to calculate the Pt lattice constant and to evaluate the average sizes of nanocrystallites by applying
Fig. 1. TEM (a) and HRTEM (b) images of core–shell nanoparticle ensembles of FP sample.
Fig. 2. Size distribution of the FP core–shell nanoparticles. A mean diameter of 3.14 nm is obtained by the best fit using the lognormal distribution function.
Scherrer method. The following values were determined for lattice constant and mean crystallite size: aPt = 3.9 Å and DPt = 5.7 nm respectively. The peak centered at 2h = 54.4° contains partially
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(a) I (a.u.)
150
Pt M 1 Pt M 2 Pt M
Fe K 1
Fe K 1
100
Cu K 1
Cu K 1
(b)
Pt L 1 Pt L 2 Pt L 1 Pt L 2
50
0
10 nm
250
500
750
1000
1250
E (eV)
Fig. 3. HRTEM image of FP-T nanoparticles (a) and the EDX spectrum corresponding to a small ensemble of agglomerated nanoparticles (b).
Fig. 4. XRD diffraction pattern of the FP core–shell nanoparticles. The corresponding diffraction line indexes for pure iron and platinum are also indicated.
overlapping diffraction lines of Pt (2 0 0) and Fe (1 1 0) and was deconvoluted by using pseudo Voigt line profiles. In this case the calculated lattice constants for iron and platinum were aPt = 3.88 Å and aFe = 2.79 Å. The corresponding nanocrystallite sizes are DPt = 5.6 nm and DFe = 4.6 nm respectively. These values are a little higher than the mean size determined from HRTEM. The difference arises from large sized nanoparticles (D > 20 nm) as seen in Fig. 1a. The large nanoparticles, even few in number, give a significant contribution to the diffraction XRD line narrowing. For both diffraction peak analyzed, the lattice constants are slightly different from standard ones namely a = 3.92 Å for Pt and a = 2.86 Å for Fe. As observed in case of Fe@Fe-oxide@Au core shell nanoparticles, prepared also by the reverse micelles method [19] the intermediate shell of oxidized Fe should be mostly amorphous, thus is ‘‘unseen’’ by XRD. The XRD patterns of thermally treated FP-T sample are presented in Fig. 5. By annealing of the FP sample some of the iron atoms from the oxide shell will diffuse into the neighboring Pt shell with the formation, at the interface, of an ordered FePt alloy. In the same time, as a result of thermal treatment, the amorphous oxide shell will be crystallized into an ordered oxides shell. This can be seen in Fig. 5 where FePt alloy (L10 phase), Fe3O4, and Fe2O3 hematite diffraction patterns are present. In this case the iron core contribution is too small to be seen. Due to the thermal diffusion of iron into platinum shell, accompanied by the formation of the FePt alloy, the remaining oxide in the vicinity of Pt becomes depleted in Fe, thus hematite rather than Fe3O4 is most expected here. The lines are narrower than in case of FP sample indicating that, as a result of coalescence, some larger nanoparticles are produced in the system (20–30 nm). Due to the multiple peaks seen in Fig. 5 and to the noisy background an accurate determination by Scher-
Fig. 5. XRD diffraction pattern of the thermally treated FP-T core–multishell nanoparticles. The corresponding diffraction line indexes for FePt alloy L10, Fe3O4 and Fe2O3 hematite are also indicated.
rer equation is uncertain. The large nanoparticles, even if they are few in numbers, give an important contribution to the diffraction patterns since they contain an enough large fraction of the total number of atoms. For further compositional analysis of FP and FP-T nanocomposites XPS associated with Ar ions etching was used. The samples were subject of consecutive Ar ions etchings until by doing any additional sputtering the XPS spectra remained unchanged in shape and intensity. At this stage they reflect the real composition of the core–shell nanoparticles. Fe 2p, O 1s, C 1s, K 2p, Cl 4s as well as Pt 4d and Pt 4f core-levels lines were identified and can be seen in XPS survey spectrum of FP sample presented in Fig. 6. The existence of potassium and chlorine lines indicates that non-reacted
Fig. 6. XPS survey spectrum of FP sample.
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small quantities of K2PtCl6 are still there while C 1s core-level line is due to the remaining CTAB surfactant. For quantitative determinations the individual core-level lines were deconvoluted and fitted. As an example, the XPS spectra of Fe 2p as well as of Pt 4d core-level doublets from FP sample, together with the corresponding deconvolutions and fitted curves are shown in Fig. 7. As expected, besides the Fe(0) peaks two additional doublets corresponding to oxidized iron were seen at higher binding energies [20]. Depending on the synthesis conditions the oxide shell is more or less thick. In our case the best fit for the Fe 2p lines was obtained considering Fe2+ and Fe3+ oxidation states. The Fe 2p main lines are positioned at: 708.4 eV (2p3/2), 721.6 eV (2p1/2) for Fe(0) (denoted A), 709.5 eV (2p3/2), 722.9 eV (2p1/2) for Fe2+ (denoted B) and 711.9 eV (2p3/2), 725.3 eV (2p1/2) for Fe3+ (denoted C). Two pairs of iron shake-up satellite features at 714.9, 728.5 eV and 718.1, 732.2 eV are shown. The first one incorporates shake – up features associated to the Fe(0) and Fe2+ main lines and the second is associated to Fe3+. In this way we wanted to avoid placing too many lines in the spectrum. The Fe 2p doublet separation was constrained between 13.2–13.4 eV. Pt 4s core-level line which is superimposed to the Fe 2p spectrum is evidenced at 723.0 eV in Fig. 7a. In case of Pt 4d core-level spectrum, as shown in Fig. 7b, three doublet peaks can be seen: 312.9 eV (4d5/2), 330.0 eV (4d3/2) associated to the main platinum shell (label A), 309.1 eV (4d5/2), 326.5 eV (4d3/2) for platinum surface states (label B) and, at higher binding energies, 316.3 eV (4d5/2), 333.4 eV (4d3/2) are the more positive features belonging Pt ions surrounded by 6 negative chlorine from the K2PtCl6 residual precursor (label C). The quantity of platinum states associated to residual K2PtCl6 is in agreement with the potassium amount as determined from K 2p core-level doublet lines. The Pt 4d doublet separation was set between 17.0 and 17.1 eV.
(a)
(b)
For p, d and f core-levels lines the restrictions used for the fit core-level XPS spectra refer to the usual relation between areas of spin–orbit coupling doublet components. By using CASA software data base the integral intensities were normalized by dividing the raw XPS deconvoluted intensities to the corresponding real sensitivity, transmission and electronic mean free path factors. It was observed that, due to the core–shell structuring of nanoparticles, by increasing the etching time the XPS intensities of the Pt outer shell lines decrease while the line intensities of inner core Fe increase until the XPS signal, as coming from a very large number of sectioned nanoparticles, became stabilized [21]. As it was mentioned, XRD does not show any specific diffraction patterns of crystalline oxides thus the oxidized iron should to be mostly into an amorphous phase. It results that the real architecture of FP nanoparticles is, as expected, Fe@Fe-oxide@Pt. Molar and weight concentrations were evaluated dividing the integral intensities by the attenuation lengths (escape depths), kAL, specific for each component of the composite system [22]. They reflect the mean depth from which emerged photoelectrons are collected into the energy analyzer. The values of kAL for a given component depend on the material density, molar weight, and mean Z number of the compound as well as on the kinetic energy of the emerged photoelectrons [23]. The values of the attenuation lengths, calculated upon Ref. [23] are as following: 0.82 and 0.9 nm for Fe(0) and Fe3O4 respectively as one refers at Fe 2p core-level, 1.83 nm regarding K2PtCl6 at K 2p core-level line, 2.14 nm for CTAB at C 1s core-level line, and 0.85 nm for Pt 4d line. The weight content of different components from samples, xi, was evaluated by using the following relation:
li kIii Ii i li ki
Xi ¼ P
ð1Þ
Here Ii represents the integral intensity of the XPS core-level peak for a specific element of a molecule used for calculation and li the corresponding molar mass. The resulted molar and weight contents of Fe(0), Fe oxide, Pt and small quantities of as mentioned non-reacted components are presented in Table 1. The radial dimensions were calculated by considering a mean diameter of 3.14 nm as determined from size distribution of nanoparticles resulted from TEM. Taking into consideration both mean size (Fig. 2) and the composition of FP nanoparticles (Table 1) the mean volumes of the three components of the system were calculated as: 5.24, 7.87, and 3.1 nm3 for Fe core, amorphous iron oxide sell, and Pt shell respectively. The small value resulted for Pt outer shell thickness represents an average between small and large particles having different platinum coverage levels. This inhomogeneity in platinum coverage is a result of the reversed micelles method used for the preparation. One specificity of the method is that the larger the water micelle formed by surfactants greater the size of the resulted nanoparticles. Besides, the micellar surfactants have a decisive role in the platinum reduction. They act as growth-directing adsorbents on metal surfaces. As the size of the water droplet increases the micellar concentration decreases thus reducing the rate of metal reduction inside the confined hydrophilic volume. Due to higher concentration of hydrophilic terminations in small micelles, the small particles will be
Table 1 The calculated molar and weight content values of the FP sample as resulted from XPS.
Fig. 7. XPS spectra of Fe 2p (a) as well as of Pt 4d (b) core-level doublets from FP sample, together with the corresponding deconvolutions and fitted curves.
FP sample
Fe(0)
Feoxide
Pt
K2PtCl6
CTAB
yi (mol.%) xi (wt.%) Shell thickness (nm)
56.7 23.9 1.08 (radius)
13.5 23.7 0.38
28.0 38.6 0.11
3.3 12.3 –
0.5 1.5 –
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favored during the reduction process and much thick and complete outer Pt shell will result [24–26]. For larger nanoparticles a rather incomplete or cracked external shell is most probable. As it can be seen in the inset of Fig. 1b the interplanar distance evidenced by0 HRTEM in case of uniform small particles has the value of 2.2 Å A which corresponds to (1 1 1) crystalline plane family of Pt. Well defined Pt crystallites are also indicated by the XRD pattern from Fig. 4. The same XPS lines were also observed in case of FP-T sample. For instance, the corresponding Fe 2p and Pt 4d features of FP-T are presented in Fig. 8. The oxidized iron features are positioned at 709.2 eV (2p3/2) and 724.4 eV (2p1/2) for Fe2+ and at 711.1 eV (2p3/2) and 724.4 eV (2p1/2) for Fe3+ respectively. To avoid having too many peaks in the deconvolution, no separate peak doublets were set for each of the Fe3O4 and Fe2O3 components. The area of Fe3+ feature was left to evolve free with respect to the Fe2+ one in order to incorporate both Fe3O4 and Fe2O3 contributions. A more intense Fe3+ peak as compared to the Fe2+ one resulted accordingly. The Fe(0) from the inner core and Fe atoms incorporated into the alloy phase have binding energies of 708.4, 707.6 eV for 2p3/2 and 721.7, 720.8 eV for 2p1/2 respectively. Again, to avoid the use of too many components in the deconvolution the only one shake-up satellite peak was considered for every doublet component. Each satellite incorporates shake-up contributions from different principal lines. The Pt 4d core-level XPS spectrum of FP-T sample is presented in Fig. 8b. Here A and B doublet features denote the platinum atoms from FePt (L10) alloy and platinum outer shell, respectively. Additionally, an intense Ar 2s feature and a plasmonic peak are also observed. Due to the uncertainty regarding the type of potassium, chlorine and carbon end products resulted from the annealing of FP nanoparticles in the presence of highly catalytic nanostructured platinum as well as due to the presence of adsorbed/absorbed
(a)
(b)
Fig. 8. XPS spectra of Fe 2p (a) and Pt 4d (b) core-levels from FP-T sample. The corresponding deconvolutions and fitted curves are also presented.
residual Ar (from Ar atmosphere inside the kiln) a quantitative analysis, similar to that determined in case of FP sample, was no longer possible for FP-T nanoparticles. Nevertheless, considerations relative to the Fe(0), Fe3O4/Fe2O3, FePt (L10), and Pt from FP-T multi-shell sample can be made in correlation with the composition of FP nanoparticles. Thus, from the ratio of A and B peaks in Fig. 8b, it results that 33% from total Pt atoms were included into FePt alloy phase. Also it is reasonable to make the assertion that the Fe(0) core remains with the same volume as in the initial FP sample. An estimation regarding the relative content of the two iron oxides is rather difficult to be made. The magnetic properties of the FP nanoparticles were examined and compared with properties of the thermally treated FT-T nanoparticles. The magnetizations of FP compound as a function of the applied field H, at room temperature (RT) and at 4.2 K, are presented in Fig. 9a. The nanocomposite has with a quite small coercive field of 29 Oe at RT and, due to the blocking process, the coercive field increases to 282 Oe at 4.2 K. In case of nanoparticle ensembles, small coercive fields represents an indication that superparamagnetic behavior may be present in the system. The saturation magnetization of the iron core (Ms), normalized to the specific Fe weight content as resulted from XPS analysis (Table 1), is about 179 emu/g(Fe), which represents a typical value for iron nanoparticles [27]. The magnetic behavior changes after thermal annealing, as seen in Fig. 9b. When the temperature decreases from 298 K to 4.2 K, the coercive field increases and the saturation magnetization remain at almost the same level as for the FP sample. It is a prime indication that some of the magnetic components are magnetically coupled. The first derivative of demagnetization/magnetization curves of FP sample compared with the same first derivative FP-T sample at RT and 4.2 K are presented in Fig. 10. Since they are symmetric we
(a)
(b)
Fig. 9. Magnetization vs. applied magnetic field, taken at room temperature and 4.2 K, corresponding to: FP nanoparticles (a) and thermally treated FP-T nanoparticles (b). The coercive fields are shown in the insets.
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will refer only to the derivative of demagnetizing curves. The first derivative of demagnetizing curve shows the local slope of the M– H dependences and thus the number of peaks or bumps present in the dM/dH vs. H curve reveals the number of magnetic phases within the sample. The horizontal shift of these peaks from their nominal coercive field positions (for pure phases) is directly related to the strengths of the magnetic phases coupling. When hard and soft magnetic phased are exchange coupled the peaks or shoulders associated with hard magnetic phases are shifted from their nominal coercivities towards zero field while the peaks of the soft phase components are moved to higher fields with respect to their usual (uncoupled) coercive field positions. In case of fully coupled systems the two peaks can merge into a single broader one [28]. In case of FP-T sample the presence of the inflection points, at both RT and 4.2 K, it is an indication that several magnetically coupled phases are present into the system. The assignment of different contributions is based on the intrinsic coercivities expected from each component as (I) coupled ferromagnetic (FM) Fe(0), (II) coupled ferrimagnetic Fe3O4, (III) coupled FM FePt, and (IV) non-coupled FM FePt. They are indicated by arrows in Fig. 10. The last attributed position is based on the usual values of coercive field [29]. The so called DM plots were also calculated as in Ref. [30]. DM plots represent the difference between the ascending and descending parts of the hysteresis loop for B > 0. Like the first derivative of the demagnetization (magnetization) curves the derivative of these plots, dM/dB, reveal, with a more accuracy, mixed coercivity fractions. One maximum (or bump) means one coercivity fraction in the loop. The superimposed plots of FP and FP-T samples, obtained from RT and 4.2 K hysteresis loops, are presented In Fig. 11. One can see that the magnetic couplings appear both at RT and 4.2 K for the thermally treated sample. Fig. 11a shows an additional inflection point of the FP-T sample that is connected
(a)
(b)
Fig. 11. DM plot for the FP and FP-T nanoparticles obtained from room temperature (a) and 4.2 K (b) magnetic hysteresis measurements.
to the coupled week FM Fe2O3. Below 260 K hematite Fe2O3 it’s an antiferromagnet, hence this feature is not longer observed in Fig. 11b. In case of the FP sample in 4.2 K DM plot a weak coupling is observed at fields related to the coupled ferrimagnetic Fe3O4. This feature it is prove of the sensitivity of the DM plot method. One has to mention that XRD measurements didn’t revealed any crystallized form with regard to the Fe oxides inside the FP sample. The temperature dependences of the mean coercive field HC, defined as HC ¼ ðHleft Hright Þ=2, for both FP and FP-T samples are presented in Fig. 12. One can see that down to about 50 K the FP-T sample has much larger coercivities than initial FP sample. Bellow 50 K, due to both blocking process and energy density anisotropy increase, the coercive field of FP sample starts to rise. In case of FP-T sample the coercive field starts rising at about 110 K then, after reaching a maximum around 25 K, it begins to decrease. This maximum of the coercive field can be related to a max-
(b)
Fig. 10. Comparison of the magnetization derivative maxima corresponding to the FP-T nanoparticles at room temperature (a) and 4.2 K (b).
Fig. 12. Coercive field temperature dependence corresponding to the FP nanoparticles and to the thermally treated FP-T nanoparticles.
O. Pana et al. / Journal of Alloys and Compounds 574 (2013) 477–485
imum of the effective anisotropy density of the composite system. It can be also an indication that some of the magnetic components of FP-T sample are exchange coupled. At more detailed analysis one can see that HC has slightly different values in case of the magnetization curve as compared to the demagnetization one. In case of FP sample at 298 K HC is 24 Oe for the magnetization process and 34 Oe for the demagnetization curve. At 4.2 K the corresponding values are 272 and 293 Oe respectively. It is an indication that small quantities of crystallized ordered iron oxide are present even in the initial FP sample thus a small exchange is possible to occur here (as already evidenced by DM plot). In case of FP-T sample, as compared to the initial sample, a more intense exchange coupling appears here. At RT and 4.2 K the demagnetization–magnetization left–right coercive fields are 131–111 Oe and 534–491 Oe, respectively. The values of HE and HC are determined from the two fields (Hleft and Hright) of the hysteresis loop where the magnetization goes through zero. The following relation was used for the exchange field: HE ¼ ðHleft Hright Þ=2. Thus, a hysteresis loop which is shifted to the left of H = 0 is defined as having a positive exchange field [31]. The evolution of the exchange field, HE for the whole temperature range together with the temperature dependence of the field position of Fe(0) maximum (I) in the derivative of the demagnetization curve, as seen in Fig. 10, is presented in Fig. 13. Iron maximum position displacement to the left side with respect to H = 0 (derivative of the demagnetization curve) it is a measure of the coupling strength so that larger the displacement more intense the coupling with the neighbor magnetite shell. As one can see the Fe(0) maximum displacement follow the general tendency of the exchange field HE behavior as function of T. Both fields have a minimum at around 260 K which corresponds to the Morin transition of Fe2O3 component. Above this temperature the bulk hematite is a week ferromagnet (WF) while below it becomes an antiferromagnet (AF). For T < 260 K the hematite nanoparticles it may or may not exhibit a spin-flop Morin transition between WF and AF states, depending on its stoichiometry, crystallinity, surface condition, and microstructure. Generally particle size does not correlate directly with the presence or absence of a Morin transition but instead, that the Morin transition is controlled by the precise lattice parameters. The annealing of samples at temperatures above 550 °C, as it is our case, resulted in a reduction of defects and an increase of TM up to 260 K [32–34]. Once the Fe2O3 AFM phase is set, an exchange bias may be established with its ferromagnetic and ferrimagnetic neighbors, the outer FePt shell and the inner Fe3O4 one, respectively. As a consequence the overall exchange field starts to increase and, through the oxides layer (hematite and magnetite), a coupling, going from
483
the outer FePt shell down to the inner iron core, can be realized. In the vicinity of 120 K the Fe3O4 component undergoes the Verwey transition accompanied by a peak in the temperature dependence of the exchange field HE. At about 130 K the Verwey transition is preceded by the so-called magnetic isotropic point (Ti) [35]. One has to remember that the iron ions, in magnetite, are distributed between two sublattices. The sublattice A has Fe3+ occupying tetrahedral positions and sublattice B with equally distributed Fe2+ and Fe3+ in octahedral positions. The magnetic moments of Fe3+ ions from A and B sublattices are antiferromagnetically coupled. Within B sublattice Fe2+ and Fe3+ magnetic moments are ferromagnetically coupled [36]. It should be mentioned that below the Verwey transition temperature, the magnetite changes its crystal structure from inverse cubic spinel to monoclinic symmetry. The process is accompanied reorientation of the easy magnetization axis along c direction. The interplay between crystal structure, magnetism and conductivity associated to the Verwey transition are still under debate [37]. Since the dimensions of the FM components are at the nanometer scale they behave like monodomanins and are subject of thermally activated magnetization fluctuations. It appears that during the Verwey transition thermal fluctuations are reduced and the exchange field increases. At about 34.7 K both fields presented in Fig. 13 reach a local maximum. The meaning of these features can be understood in connection with the distribution of the energy barriers as determined from field cooled (FC) and zero field cooled (ZFC) temperature dependences of magnetizations. The magnetization versus temperature in ZFC–FC regimes were measured in 100 Oe magnetic field. In Fig. 14 there are represented the temperature dependence of the heights of the energy barriers for FP sample (a) and FP-T sample (b). They were deduced in accordance with Refs.
(a)
(b)
Fig. 13. The evolution of the exchange field, HE for the whole temperature range together with the temperature dependence of the field position of Fe(0) maximum (I) in the derivative plot of the demagnetization curve, as seen in Fig. 10.
Fig. 14. The energy barriers distributions that separates the two orientations of the nanoparticles magnetization corresponding to the FP (a) and thermally treated FP-T nanoparticles (b). They were calculated from FC and ZFC M = f(T) dependences shown in the insets.
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[19,38]. The FC and ZFC temperature dependences for the two samples are shown in the corresponding insets of Fig. 14a and b. One can see that a maximum appears in the ZFC curve from the inset of Fig. 14a at a temperature of about 110 K. It represents the mean blocking temperature TB of the magnetic nanoparticle system. Considering the small size of iron cores, the existence of a small coercive field, as indicated above in Fig. 9a, and the presence of a mean blocking temperature TB one can confirm that FP core–shell has a superparamagnetic character [19,21,22,38,39]. The temperature dependences energy barrier heights are connected to the Keff V product as:
K eff V ¼ kBT ln
s s0
ð2Þ
Here Keff represents the effective anisotropy energy density, V is the magnetically ordered volume of the nanoparticles and s and s0 are the observed and characteristic magnetization switching times respectively. The graph maximum from Fig. 14a, positioned at 17 K, corresponding to the energy barriers distribution of Fe(0) cores from FP sample [24] can be used to determine the effective anisotropy energy density constant K Fe eff [38]. Assuming an iron core volume of 5.24 nm3 as determined from XPS analysis a value of 6 3 K Fe eff 1:1 10 J=m results. This value is consistent with values determined for small sized Fe nanoparticles [24]. The energy barrier distribution is completely different in case of FP-T sample where two maxima are present into the graph. The intense maximum at 113 K is connected to Fe3O4 component and corresponds to the Verwey transition preceded at around 130 K by the inversion point of magnetite crystalline anisotropy [35]. The maximum at Tm = 44K corresponds to energy barriers distribution of iron cores, shifted, in FP-T sample, to higher T due to an increased K Fe eff value. The shift of the Fe(0) contribution can be attributed to the exchange existing between the iron cores and the newly formed magnetite intermediate shell. It represents the sum of several contributions: Fe K Fe eff ¼ K ð0Þ þ
6 ðK ðSÞ þ K ðSÞex Þ þ K ex : DFe
ð3Þ
Here K Fe ð0Þ is the axial intrinsic iron anisotropy energy density and Kex represents the exchange contribution. K(S) and K(S)ex are the intrinsic surface anisotropy and the exchange contribution to the surface anisotropy, respectively. Due to the roughness and inhomogeneities at the Fe(0)/Fe3O4 interface the exchange couplings are quite complex [37]. However, one can estimate that the coupling of surface Fe magnetic moments with Fe3+ occupying tetrahedral positions of the sublattice A or octahedral sites of sublattice B will have an exchange bias character thus contributing to K(S)ex. The coupling with Fe2+ states of B sublattice involves a direct exchange coupling and contributes to Kex. Also superexchange involving O 2p spin paired orbitals is possible and probably contributes to Kex. A quantitative estimation of various contributions seems to be difficult to be made at the present stage. At a detailed analysis, the maximum seen in Fig. 13 at T 0m ¼ 34:7 K (marked by an arrow) represents the same blocking process of the iron cores as observed at Tm = 44 K in energy barriers distribution from Fig. 14b. Due to the temperatures steps in recording different hysteresis loops its real position should be probably in-between the two consecutive points, namely 34:7 K T 0m 49:7 K. Further, considering the maximum associated with Fe(0) core from Fig. 14b, this temperature range can be reduced to 34:7 K T 0m T m . On the other hand, the graphs presented in Fig. 14 were deduced from thermo remnant magnetizations following [38,40], and as a consequence Eq. (2), which stands for H = 0, applies here. On the other hand the temperature dependence
of the exchange field, HE, was calculated from hysteresis loops taken at different temperatures, and thus Eq. (2) is no longer valid. Instead the following relation applies here [40]:
a H s K eff V 1 ¼ kB T ln HC s0
ð4Þ
Here HC is the field at which the given barrier vanishes. The exponent a takes values between 1.5 and22 [40]. For the a = 2 case and, if 0 < H/HC < 2, the factor 1 HHC act towards reducing the height of the energy barrier, thus decreasing the temperatures at which magnetization thermal fluctuations are blocked. At zero magnetization HE is the local field acting on Fe(0) cores while the field which cancels the magnetization is HC + HE, as recorded in the demagnetization curve. As a consequence the temperature associated to the maximum for which most of the iron cores are blocked is shifted to lower values as in Fig. 13. By considering Eqs. (2) and (3) the ratio between T 0m =T m became
2 T 0m HE ¼ 1 Tm HC þ HE
ð5Þ
Considering the corresponding values of HE and HC one obtains for T 0m ¼ 41:2 K which is a temperature within the expected range. The incompletely resolved maximum at low temperatures as seen in Fig. 14b can be attributed to the blocking of FePt thin shell. This assertion is sustained by calculating the blocking temperature of FePt alloy. As resulting from XPS analysis of spectrum shown in Fig. 8b it results that about one third of the initial platinum metal turns into FePt L10 alloy and thus its estimated mean volume is 6 3 about 1 nm3. Considering K FePt eff 6—7 10 J=m [29] one obtains FePt FePt a value of T B ¼ 16:1—18:7 K depending on T eff value and the value chosen for ln ss0 namely 25–27. At temperatures below these values the exchange field HE has a sharp increase due to the blocking of FePt alloy shell. Concerning the magnetite, below the Verwey transition temperature and, due to the neighbor couplings (high in3 O4 crease of K Fe ), it can be considered as being in a blocked state. eff Thus, below 16.1–18.7 K temperatures range the magnetization of one particle behaves according to a ‘‘superspin’’ configuration. In other words, the whole magnetization rotates at once within all coupled shells. As a consequence, even if HE increases with decreasing temperature (Fig. 13), HC field, which expresses the ‘‘opposition’’ to the magnetic field reversal, decreases due to the lack of intra-shell local fluctuations. This behavior is evidenced by the maximum seen, at low temperatures in Fig. 12 for FP-T sample. 4. Conclusions The reverse micelles method was used for preparation of Fe/Pt ‘‘core–shell’’ nanoparticles. Due to the specificity of method the Fe core is covered by a layer of amorphous iron oxide, so that the actual structure of nanoparticles with average size of about 3 nm is Fe@Fe-oxide@Pt. Starting from this type of particles by heat treatment in inert atmosphere (Ar) a complex system of ‘‘coremultishell’’ structure Fe@Fe3O4@Fe2O3@FePt@Pt was obtained. The magnetic properties of these two systems were analyzed and compared. The thermally treated nanoparticles show one order of magnitude higher coercivities than the initial Fe@Fe-oxide@Pt nanoparticles, preserving in the same time, at least at RT, an almost superparamagnetic behavior. At detailed analysis of hysteresis loops of annealed nanoparticles shows that HC has slightly different values for the magnetization curves as compared to the demagnetization ones thus proving the existence magnetically coupled multiple shells. The Verwey transition was evidenced at 113 K in case of Fe3O4 shell. An exchange bias seems to exists, bellow 260 K, at the inter-
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faces of Fe2O3 shell as well as between some of the surface magnetic moments of iron core and the Fe3+ moments from both A and B sublattices of the magnetite shell. Due to these couplings the Fe(0) cores of annealed nanoparticles have an increased effective anisotropy constant and became almost blocked at higher temperatures than in case of initial Fe@Fe-oxide@Pt nanoparticles (44 K instead of 17 K). Generally, when decreasing temperature, the evolution of the exchange field shows an increase, with two local maxima. One maximum is at the Verwey transition temperature and another one at a temperature corresponding to the blocking of the Fe cores. Thus the effective anisotropy constant of the iron core seems to be the sum of volume axial anisotropies and surface anisotropies originating from intrinsic and exchange contributions. Due to the complexity of the nanoparticles resulted from thermal treatment, at the present stage, a quantitative estimations for these contributions seems to be difficult to achieve. Acknowledgement This work was supported by the Romanian Ministry of Education and Research under the research Program POS-CCE-METAVASINT – 550/2010 and ID – 119/2011. Appendix A. Supplementary data Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j.jallcom.2013.05. 153. References [1] [2] [3] [4] [5] [6] [7]
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