Enhancing the dual magnetic and optical properties of co-doped cerium oxide nanostructures

Abstract

In spite of the potential biomedical application, ceria nanoparticles suffer from poor optical and magnetic properties. Here we report the structural, optical and magnetic properties of iron (8 at%) and europium (1 at%) co-doped ceria nanoparticles with respect to annealing temperature. A cubic fluorite structure of ceria was observed on co-doping but an additional α-Fe₂O₃ phase emerged on annealing at 700 °C and above. An increase in the mean crystallite size from 6.5 to 37 nm and a corresponding reduction in strain was observed with annealing temperature. Surface area and scanning electron microscopic studies indicated a porous structure which became dense upon annealing. Raman spectroscopic studies reveal the presence of oxygen vacancy defects arising from the combination of intrinsic and extrinsic defects due to the presence of Ce³⁺ and dopant, respectively. X-ray photoelectron spectroscopic results confirm the oxidation of Ce³⁺ to Ce⁴⁺ state while the trivalent state of iron was retained on annealing at higher temperatures. On annealing above 500 °C, reduction in defect concentration improved the emission intensity primarily through magnetic dipole transition. A typical ferromagnetic behaviour was noticed in all the samples and shift a from soft to hard magnetic behaviour upon annealing.

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1. Introduction

Optical, chemical and electronic properties of lanthanides have attracted much attention due to the presence of their 4f sub-shell electrons leading to various applications.^1–4 In particular cerium oxide (ceria, CeO₂) nanoparticles have shown interesting biomedical properties such as antioxidant, antibacterial activity and being considered as a therapeutic agent in the treatment of reactive oxygen and nitrogen radical mediated diseases.^5–7 The unique properties of CeO₂ arise from the existence of cerium ion in dual oxidation states of +3 and +4 which is reversible. As a result cerium ion acts as a oxygen storage buffer and changes the reactivity with respect to environment. In the nano regime, formation of Ce³⁺ is favoured and to compensate the charges, oxygen vacancies are generated as shown by the following Kröger–Vink notation

The biocatalytic properties of CeO₂ can be engineered through the formulation of oxygen vacancies.⁸ Since CeO₂ exhibit weak emission in the visible region using these nanoparticles for theranostics becomes difficult. The emission properties of CeO₂ can be tuned by either down or up conversion processes.^9–11 In both the processes lanthanide based dopants are generally introduced into the ceria matrix to provide the visible emission. Europium (Eu) is one such dopant reported for improving the optical properties and it also provides information about the local symmetry environment in a crystal.⁹

Magnetic nanocrystals have been investigated for its biomedical application especially in the area of targeted drug delivery. Tumour targeting using magnetic fields to direct the movement and localisation of drugs to solid tumours have generated much interest.¹² Sundaresan et al., observed that the room-temperature ferromagnetism (RTFM) of many metal oxide nanoparticles have been induced by doping with transition metals which has partially filled d-shells due to exchange between electron spin moments caused by oxygen vacancies (V_O). In the nano regime, by engineering the defects it is possible to introduce defect induced magnetism.¹³ Cerium oxide nanoparticles are known to impart room temperature ferromagnetism and the correlation between defects and ferromagnetism by doping with nonmagnetic cation (Cr³⁺) doping have been reported for spintronic applications.^14,15 Wang et al., investigated the effect of iron concentration on electronic and structural properties of CeO₂ nanoparticles using X-ray absorption and X-ray emission spectroscopy.¹⁶ Doping low concentration of iron resulted in the ferromagnetism while at higher iron concentration lead to an antiferromagnetic behavior due to the reduced distance between iron atoms.

Dual imaging through optical and magnetic means offer more information than single probe imaging. Though many reports are available on tuning either optical or magnetic properties in CeO₂, no reports are available on combining both the properties. The present work focuses on the synthesis of iron and europium co-doped cerium oxide nanostructures and correlate the size, structural, surface, optical and magnetic properties with respect to annealing temperature.

2. Experimental

2.1. Preparation of co-doped CeO₂ nanostructures

Iron (8 at%) and europium (1 at%) co-doped ceria was prepared by low temperature combustion route using cerium(III) nitrate hexahydrate (Ce(NO₃)₃·6H₂O), ferric(III) nitrate hexahydrate (Fe(NO₃)₃·9H₂O), and europium(III) nitrate pentahydrate (Eu(NO₃)₃·5H₂O) as metal precursors (oxidizer) and glycine as fuel.¹⁷ The calculated amount of nitrate and glycine was dissolved in distilled water and stirred for 30 minutes. The solution was heated over a hot plate at 80 °C which resulted in a highly viscous liquid. The viscous liquid was heated in a furnace at 150 °C and auto ignition of the mixture lead to the evolution of large volume of gases to produce a voluminous powder. The oxidizer to fuel ratio (O/F) was kept constant at 1 : 1.2 for the combustion synthesis.¹⁷ The overall reaction can be represented by the following equation

The resultant powder was washed with water to remove traces of any unreacted nitrate/glycine and then dried at 100 °C for 2 hours in air. To understand the influence of temperature on properties, the as prepared powder (EFCR) was annealed at 300, 500, 700 and 900 °C (sample coded as EFC3, EFC5, EFC7 and EFC9, respectively). For comparison pure cerium oxide (CeO₂) and iron oxide (Fe₂O₃) nanoparticles were also prepared by the similar procedure.

2.2. Characterization and measurements

X-ray diffraction (XRD) was recorded using a Rigaku Ultima IV X-ray diffractometer with monochromatic CuKα radiation in the range of 20–80° with a scan rate of 0.02° per second for the phase identification. Philips CM200 ATEM operated at 200 keV was used for transmission electron microscopic (TEM) studies. Energy dispersive X-ray spectroscopic (EDAX) studies were carried out to quantify the elements using Oxford X-Max 80 SDD detector attached with TEM. The surface morphology of the powder was studied by scanning electron microscope (SEM; Hitachi S-3400N). The specific surface area of the as-synthesized powder was measured by the standard Brunauer–Emmett–Teller (BET) technique with N₂ adsorption using Gemini VII 2390 and BJH method was used for the determination of pore volume. Optical absorption spectra were recorded by Perkin Elmer Lambda 650-S. Horiba Scientific Fluoromax-4 spectrofluorometer was used for studying photoluminescence (PL) properties under the excitation wavelengths of 330 and 465 nm. Magnetic measurements were carried out using Lake Shore 7404 Vibrating Sample Magnetometer (VSM) at room temperature. Renishaw Laser confocal Raman Microscope RM 2000 was used for recording Raman spectra with 514 nm Ar⁺ laser with a power of 0.5% focused on the sample with beam diameter of 1 μm, exposed for 30 s and data was collected at the intervals of 1 cm⁻¹. X-ray photoelectron spectroscopic (XPS) studies were carried out using Al K_α source in Specs (Germany).

3. Results and discussion

3.1. Structural and surface characterization

Fig. 1(a) shows the powder XRD pattern recorded for the samples with respect to annealing temperature. Bulk ceria has a cubic fluorite structure where each Ce⁴⁺ ions are coordinated to 8 oxygen atoms in a cubic octahedral symmetry. All the samples showed peaks corresponding to (111), (200), (220), (311) and (222) reflection planes of ceria (ICDD card no: 00-034-0394).¹⁸ XRD spectra of the as prepared iron oxide under similar processing conditions resulted in the formation of mixed maghemite γ-Fe₂O₃ (cubic) and haematite α-Fe₂O₃ (rhombohedral) phases matching with the ICDD card no 01-089-5892 and 01-079-007, respectively. Though α-Fe₂O₃ forms at higher temperature, the liberation of larger amount of heat during the combustion process lead to the formation of α-Fe₂O₃ along with γ-Fe₂O₃.¹⁸ The slow scan XRD results did not exhibit any peak corresponding to iron oxide formation for the as prepared and the samples annealed upto 500 °C which indicates the presence of iron as dopant in the CeO₂ lattice. However a weak peak emerged upon annealing at 700 and 900 °C due to the formation of α-Fe₂O₃ (Fig. 1b).

Fig. 1 (a) X-ray diffraction pattern for as prepared and annealed samples along with undoped oxides (^#indicates α-Fe₂O₃ and the rest of the peaks in Fe₂O₃ corresponds to γ-phase) and (b) slow scan XRD analysis in the range of 34–38° showing the presence of α-Fe₂O₃ (110).

In comparison to CeO₂, other samples (EFCR to EFC9) exhibited a shift in peak position towards lower angles. The ionic radii for Ce³⁺, Ce⁴⁺, Fe³⁺and Eu³⁺ ions are 0.1034, 0.092, 0.064 and 0.121 nm, respectively.²⁰ In contrary, it was expected that the replacement of higher ionic radii (Ce⁴⁺) with smaller ionic radii (Fe³⁺) ion should lead to a shift in peak position towards higher angle. The observed changes in XRD can be attributed to the size and strain effect arising upon doping. To understand the effect of doping on strain Williamson–Hall plot was used

where β is the full width half maxima, λ is the wavelength of X-ray, D is the crystallite size and ε is the residual strain. The calculated size and strain values are shown in Fig. 2. As prepared CeO₂ had a larger mean crystallite size (13.6 nm) than that of the as prepared co-doped sample EFCR (6.5 nm). The smaller crystallite size emphasizes the role of dopant present between the surfaces of the CeO₂ nuclei thereby reducing the growth process.¹⁸ Formation of smaller crystallite favours the reduction of Ce⁴⁺ to Ce³⁺ ion and facilitates the oxygen vacancy generation as shown by eqn (1). Deshpande et al., reported the correlation between crystallite size, lattice parameter and Ce³⁺ concentration in ceria nanoparticles.²¹ It was observed that the smaller crystallite size lead to peak shift towards lower angles than bulk ceria due to the formation of higher concentration of Ce³⁺ which has larger ionic radii. Though many reports are available showing the higher angle shift of ceria peaks on iron doping but that is true for similar or comparable crystallite sizes.²² In the present work, presence of Fe³⁺ as dopant shift the peak to higher angles while the prevalence of Ce³⁺ shift towards lower two theta position. The observed lower angle peak position than the expected higher angle shift for EFCR in comparison to CeO₂ can be attributed to the influence of smaller crystallite size in EFCR which results in higher Ce³⁺ concentration. The calculated strain shows a higher value for EFCR due to the presence of co-dopants as well as the larger ionic radii of Ce³⁺ than that of CeO₂. A marginal shift in strain was observed for EFC3 and EFC5 compared to EFCR (Fig. 2).

Fig. 2 Size and strain for samples calculated from XRD data.

Analysis of XRD data in the range of 34–38° showed a weak peak for EFC7 and EFC9 at 35.6° (Fig. 1b) corresponding to (110) plane of α-Fe₂O₃.¹⁹ The formation of α-Fe₂O₃ in EFC7 and EFC9 was further confirmed by our Raman spectroscopy results discussed later. It has been reported that annealing above 500 °C initiates the phase change from γ-Fe₂O₃ to α-Fe₂O₃ and complete transformation to α-Fe₂O₃ at 700 °C.²³ Due to doping in the as prepared and samples annealed upto 500 °C no traces of any of the phase of iron oxide was observed in XRD. But annealing above 500 °C resulted in the segregation and outward diffusion of iron ion and subsequent oxidation to form α-Fe₂O₃. A further shift in peak position towards lower angles with a reduction in peak width was observed for EFC7 and EFC9. Also a drastic reduction in strain and an increase in crystallite size (27–37 nm) was noticed in EFC7 and EFC9 due to the phase segregation of iron oxide and thermally assisted particle growth (Fig. 1b). Annealing at 700 °C results in the oxidation of Ce³⁺ to Ce⁴⁺ which shifts the peak position to lower angles. Further increase in crystallite size of EFC9 can be attributed to the formation of liquid phase of iron oxide at 900 °C which favours the growth and coalescence of the particles.²⁴

Surface area calculated from BET method was found to be 32.2, 26.1 and 9 m² g⁻¹ for EFCR, EFC5 and EFC9, respectively. Also, increase in annealing temperature lead to the reduction in pore volume from ∼0.14 cm³ g⁻¹ (EFCR and EFC5) to 0.004 cm³ g⁻¹ (EFC9). The ratio (φ) between the size calculated from BET surface area (D_BET) and XRD (D_XRD) showed a value of about 2.21. The deviation of φ from the ideal value of 1 indicates the presence of porous agglomerates supporting our SEM results (discussed below).

TEM analysis was carried out to understand the size and phase changes associated with the samples. Fig. 3 shows the TEM images of the as prepared co-doped sample (EFCR) and after annealing the same at 900 °C (EFC9). TEM analysis of EFCR showed a crystallite size of 7–8 nm matching with that of XRD. But EFC9 exhibited a strongly fused particles due to iron oxide mediated liquid phase growth.²⁴ Further the colour of the sample changed from yellow (EFCR) to red colour (EFC9) due to the segregation of iron oxide as supported by our XRD results. Selected area electron diffraction (SAED) of EFCR showed polycrystalline nature and the calculated d-spacing from the ring patterns correspond to (111), (200), (220) and (311) planes of ceria while EFC9 showed only one visible weak ring pattern from (111) plane.

Fig. 3 Transmission electron microscopic images of EFCR (left) and EFC9 (right) and the inset shows the corresponding SAED pattern. On annealing colour of the sample changed from yellow (EFCR) to red (EFC9).

SEM studies on EFCR, EFC5 and EFC9 were carried out to understand the surface morphology, elemental chemistry and distribution (Fig. 4 and 5). A porous spongy like structure was observed for EFCR that can be attributed to the gases evolved during the combustion process matching with XRD and surface area results (Fig. 4a). The presence of Ce, Eu and Fe in co-doped sample EFCR was confirmed by EDAX analysis (Fig. 4b). The elemental composition was found to be approximately 91.5, 1 and 7.5 at% for Ce, Eu and Fe, respectively, matching with the theoretical stoichiometry. Though EDAX analysis showed the presence of Ce, Eu and Fe to determine the distribution elemental mapping was carried out. Fig. 5 shows SEM micrographs for EFCR, EFC5 and EFC9 along with the corresponding elemental distribution for Ce, Eu, Fe and O. Dopant elements had a uniform distribution without any segregation. Due to annealing, EFC9 exhibited a dense network which is in agreement with TEM observation. In EFC9 an increase in oxygen count was observed that can be attributed to the increase in oxygen concentration from the oxidation of Ce³⁺ to Ce⁴⁺ and iron ion to the corresponding oxide.

Fig. 4 SEM showing (a) surface morphology and (b) EDS spectra for EFCR.

Fig. 5 SEM micrographs and the corresponding elemental mapping for EFCR, EFC5 and EFC9 (scale bar – 10 μm).

Fig. 6 depicts the Raman spectra of the samples recorded at room temperature using a laser source with an excitation wavelength of 514 nm. Bulk CeO₂ shows a single Raman active peak at 465 cm⁻¹ from symmetric breathing mode of O atoms around each Ce⁴⁺.²⁵ The Raman active F_2g mode is primarily observed for metal oxides having fluorite structure. To identify the peak position a Lorentzian peak fitting profile was used. In the present work a red shifted Raman peak at 458 cm⁻¹ was observed for CeO₂ which can be attributed to the effect of lattice strain and phonon confinement due to nanoparticle size. Upon co-doping EFCR exhibited an asymmetric, broad peak which shifted to lower energy (455.8 cm⁻¹) than CeO₂. The shift can be induced by the presence of strain and smaller crystallite size in EFCR than CeO₂ supporting our XRD results (Table 1). EFC7 and EFC9 showed larger shift towards higher energy than any of the samples which can be attributed to the increase in particle size and reduction in strain.

Fig. 6 Raman spectra for the as prepared and annealed samples obtained using an excitation wavelength of 514 nm.

Table 1 Calculated bang gap energy, defect concentration from F_2g mode of ceria and observed magnetization values for samples

Sample	Optical	Raman		Magnetic
	Direct band gap (eV)	Peak position (cm^−1)	Defect concentrations (×10^21 cm^−3)	Magnetization (×10^−4 emu g^−1)	Coercivity (G)
CeO2	3.08	458.1	1.28	6.56	97.9
EFCR	2.60	455.8	2.90	35.7	108.6
EFC3	2.56	457.5	2.35	207	90.2
EFC5	2.49	457.9	1.86	267	113.7
EFC7	2.84	458.3	0.79	199	444.1
EFC9	2.96	458.4	0.74	53.7	516.8

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The concentration of oxygen vacancy was estimated using the spatial correlation model. Raman line narrowing can be described by the dependence of its half width at half-maximum (HWHM) on the grain size. Grain size (d_g) and HWHM (Γ) are inter-related by the correlation proposed by Weber et al.,²⁶

Γ (cm⁻¹) = 5 + 51.8/d_g (4)

The defect concentration is related to the correlation length and grain size through the following equation.

where L is correlation length (average distance between two lattice defects), α is radius of CeO₂ units (0.34 nm). The defect concentration (N in cm⁻³) can be calculated from the correlation length L (nm) as a function of grain size

N = 3/4πL³ (6)

The calculated defect vacancy concentration is given in Table 1. EFCR showed higher defect concentration than that of CeO₂. The defect concentration decreased with annealing temperature due to the migration of oxygen vacancies towards the surface and subsequent annihilation. EFC7 and EFC9 showed defect concentrations lower than that of CeO₂ due to the increase in particle size and conversion of Ce³⁺ to Ce⁴⁺.

In addition to F_2g mode, two peaks are observed in Raman spectra at 550 and 600 cm⁻¹ which correspond to intrinsic (due to Ce³⁺) and extrinsic (due to doping) oxygen vacancies, respectively.^25,26 A perfect stoichiometric CeO₂ possess negligible defect concentration due to the presence of cerium completely as Ce⁴⁺ ion. To correlate the role of Ce³⁺ and dopant towards the defect concentration, the two peaks around 550 and 600 cm⁻¹ were deconvoluted to calculate the area A_Ce³⁺ and A_dopant, respectively. The peak areas of A_Ce³⁺ and A_dopant plotted against area of the F_2g mode (A_host) is shown in Fig. 7 and the inset represents the deconvoluted plot for EFCR. The contribution towards oxygen vacancy formation from Ce³⁺ in CeO₂ and EFCR are found to be similar while the introduction of dopant ions (from Fe³⁺ and Eu³⁺) into the lattice resulted in the generation of additional oxygen vacancies in EFCR. This observation implies that two Ce⁴⁺ ions have been replaced by two Fe³⁺ or Eu³⁺ ions to created one oxygen vacancy (Table 1). EFC5 showed a decline in the ratio of the area peaks around 550 and 600 cm⁻¹ in comparison to F_2g due to the partial segregation of iron and oxidation of Ce³⁺ to Ce⁴⁺. It can be seen that 600 cm⁻¹ peak appeared only for EFCR, EFC3 and EFC5 but disappeared in EFC7 and EFC9. EFC7 and EFC9 did not exhibit 600 cm⁻¹ peak indicating the absence of iron ion in the ceria lattice as substitutional dopant due to the outward segregation to form iron oxide which supports our XRD observation. The variation in defect concentration has been found to be marginal in EFC7 and EFC9 due to the presence of cerium primarily in the form of Ce⁴⁺.

Fig. 7 Ratio of defect concentration for peaks derived from Raman spectra and inset shows the representative deconvoluted peaks for EFCR (R² = 0.9973).

The Raman spectrum of Fe₂O₃ (Fig. 6) exhibited several peaks positioned between 217 to 1320 cm⁻¹.^27,28 For α-Fe₂O₃ seven Raman active phonon modes are expected from two A_1g modes (225 and 498 cm⁻¹) and five E_g modes (247, 293, 299, 412 and 613 cm⁻¹) while a three broad structures around 350, 500, 700 cm⁻¹ are reported for γ-Fe₂O₃. Due to two magnon scattering an additional peak results at 1320 cm⁻¹ for α-Fe₂O₃.²⁸ In the present work Raman peaks of pure Fe₂O₃ indicate the presence of both α- and γ-Fe₂O₃. On the other hand Raman spectra of EFCR, EFC3 and EFC5 did not show any peaks due to the absence of Fe₂O₃ phase. But EFC7 and EFC9 exhibited intense peaks at 217, 289, 395 and 1320 cm⁻¹ which can be attributed to the presence of α-Fe₂O₃. The results from XRD and Raman spectroscopy clearly demonstrates the presence of α-Fe₂O₃ in EFC7 and EFC9, thereby indicating the phase separation from the doped structure. Further XPS studies were carried out to quantify the changes in oxidation states (ESI, Fig. S1–S3†). A significant reduction in Ce³⁺ concentration was observed on annealing (EFCR: 29% and EFC9: 9%) while the iron remained in the +3 oxidation state for all the samples. The XPS and Raman spectroscopic results complement the oxidation state and defect concentration changes.

3.2. Optical properties

To evaluate the dopant effect on optical properties optical absorption and photoluminescence studies were carried out. Absorption characteristic of ceria depends on size and oxidation state of cerium (+3 or +4). The optical absorbance spectra of the pure and doped samples are shown in Fig. 8. CeO₂ showed strong absorption in the UV region with two overlapping broad bands around 250–280 and 340–360 nm. UV absorption of CeO₂ do not occur by inner band transition between Ce 4f and Ce 3d levels but by a charge transfer from O²⁻ (2p) to Ce⁴⁺ (4f) bands.²⁹ Unlike CeO₂, the co-doped samples exhibited a broader but stronger absorption across the visible region and the absorption increased with annealing temperature. The bands featured in CeO₂ red shifted significantly in EFCR, EFC3 and EFC5 which can be attributed to the higher concentration of Ce³⁺. Pure Fe₂O₃ exhibited several absorptions in the UV region (<300 nm) due to O → Fe³⁺ charge-transfer bands of isolated Fe ions in tetrahedral and octahedral coordination.³⁰ This transfer occurs from the highest lying O 2p orbital to the half-filled Fe 3d orbital (t_2g → e_g in O_h symmetry and e → t₂ in T_d symmetry). The main absorption band around 530 nm, commonly associated to d–d (⁶A_1g → ⁴T_1g and ⁶A_1g → ⁴T_1g) transitions in α-Fe₂O₃.²⁸ EFC7 and EFC9 showed similar absorption profile as that of Fe₂O₃ which is a characteristic of d–d transition in α-Fe₂O₃ confirming the segregation of iron oxide.

Fig. 8 Room temperature optical absorbance spectra for samples.

Using absorption data, the direct band gap (E_g) of particles of pure and doped CeO₂ was calculated by Tauc plot using the following equation

where α is the absorption coefficient, hν is the photon energy, E_g is the band gap energy for direct transitions (transitions between the conduction band and the valence band with photon emissions) and C is a constant. CeO₂ showed a band gap of 3.08 eV (Table 1) which is lower than that of bulk ceria (3.15 eV).³¹ Upon co-doping the band gap was reduced to 2.6 eV for EFCR. Annealing upto 500 °C (EFC5) resulted in the red shift in band gap energy than that of CeO₂ but the band gap shifted to higher energy above 500 °C (for EFC7 and EFC9). The observed shift in band gap energy can be explained in terms of size and oxidation state of cerium. Smaller the size and higher concentration of Ce³⁺ in CeO₂ enhances the absorption in visible region. Co-doping (EFCR) reduced the size and lead to a red shift in the visible region compared to that of CeO₂. Due to the increase in size and oxidation of Ce³⁺ to Ce⁴⁺ a blue shift was observed that can be attributed to the presence of Ce⁴⁺ state.

PL emission spectra recorded at the excitation wavelengths of 330 and 465 nm are shown in Fig. 9(a) and (b), respectively. The wavelength of 330 and 465 nm corresponds to charge transfer absorption band of ceria arising from Ce⁴⁺–O²⁻ transition and ⁷F₀ – ⁵D₂ absorption from Eu³⁺, respectively. CeO₂ did not exhibit strong emission due to the presence of empty 4f shell in Ce⁴⁺.³¹ A typical Eu³⁺ based emission was observed due to various transitions of ⁵D₀ – ⁷F_J (J = 0, 1, 2, etc.) and the intensity of the peaks increased with an increase in annealing temperature. The observed emission peaks at 545, 572 and 591 nm can be assigned to transitions ⁵D₁ – ⁷F_0, ⁵D₀ – ⁷F₀ and ⁵D₀ – ⁷F_1, respectively. The peaks observed at 611 and 632 nm arise from ⁵D₀ – ⁷F₂ transition. The splitting of structure sensitive ⁵D₀ – ⁷F₂ transition indicates the presence of Eu³⁺ ion in at least two structural arrangement differing in symmetry. According to Judd–Ofelt (J–O) theory for rare earth ion present in a host matrix, the emission peak appear from the cumulative effect of the magnetic dipole (MD) and electric dipole (ED) transitions. As per J–O theory ED transition positioned at 611 and 632 nm (⁵D₀ – ⁷F₂) is only allowed in the absence of inversion symmetry and is sensitive to the local electric field. The MD transition at 591 nm (⁵D₀ – ⁷F₁) is a magnetic dipole allowed which is insensitive to the crystal environment.^32,33

Fig. 9 Photoluminescence spectra of samples excited at (a) 330 nm and (b) 465 nm.

In the present work PL excitation at 330 nm for EFCR, EFC3 and EFC5 showed weak peaks only at 545 and 591 nm.^19,28 The absence of ⁵D₀ – ⁷F₂ transition within 4f subshell which is parity forbidden (Laporte forbidden) can be attributed to the presence of defects or Eu³⁺ ions occupying the high inversion symmetry sites in the ceria lattice. The presence of higher oxygen vacancy concentration due to co-doping (Table 1) in EFCR, EFC3 and EFC5 reduces the intensity of emission by trapping the radiation thereby dissipating the energy and reducing the efficiency of overall transition. EFC7 and EFC9 exhibited a strong ED transition due to the reduction in defects, increased crystallization and structural rearrangement disrupting the inversion symmetry around Eu³⁺ ions. A weak MD transition at 591 nm was observed in EFCR, EFC3 and EFC5 while an intense peak was observed for EFC7 and EFC9 due to the presence of local magnetic field introduced from the formation of iron oxide as well as reduction in defect concentration.³⁴

When excited at 465 nm peaks were observed at 591, 611 and 632 nm only for samples annealed at 700 °C and above (EFC7 and EFC9) due to the reduction in oxygen vacancy concentration. Compared to 330 nm excitation at 465 nm resulted in broader peaks which can be assigned to the low transition probabilities upon exciting the dopant directly. The intensity ratio between I(⁵D₀ – ⁷F₂)/I(⁵D₀ – ⁷F₁) known as asymmetry ratio (R) is sensitive to the local symmetry environment around Eu³⁺ – oxygen coordination and have low value for high local symmetry around Eu³⁺ ion.^35,36 The calculated R value was 0.9 and 0.76 for EFC9 on excitation at 330 and 466 nm, respectively. Relatively low value of R than that reported in the literature indicated the presence of Eu³⁺ in high symmetry with inversion centre.^31,33,34 It has been reported that at lower concentrations, Eu is mainly located in sites with inversion symmetry leading to higher probability for MD transition.³³ The presence of europium at low concentration (1 at%) lead to the MD transition in the present work. However no significant emission above 640 nm was observed for the samples independent of excitation wavelengths.

3.3. Magnetic properties

Fig. 10 shows the room temperature magnetization of the samples measured as a function of field and the calculated saturation magnetization (M_s) and coercivity values are given in Table 1. All the samples showed a symmetric hysteresis loop of a typical of ferromagnetic material. Ceria and co-doped ceria exhibited ferromagnetic behaviour in the as-prepared and annealed conditions. The saturation magnetization increased gradually upon co-doping in the ceria lattice on annealing upto 500 °C and subsequently decreased at higher annealing temperatures.

Fig. 10 Magnetization vs. magnetic field plot for various samples at room temperature.

Though stoichiometric ceria is diamagnetic, a weak RTFM has been observed in the present work. Various researchers have reported RTFM in ceria nanoparticles.³⁷ The origin of RTFM is attributed to the oxygen vacancy present in CeO₂ nanoparticles. At the nano regime Ce⁴⁺ gets partially reduced to Ce³⁺ leading to the existence of cerium cation in mixed valency which deviates from perfect stoichiometry. To maintain the electrical defects are introduced in the lattice of ceria as represented by eqn (1). Our Raman results support the presence of oxygen vacancy in ceria (Table 1). This defect can induce magnetic moment in the nearby Ce³⁺ atoms. The Ce³⁺ have an electronic configuration belonging to partially filled 4f orbital with one unpaired electron contributing towards RTFM while Ce⁴⁺ with empty 4f orbital which do not contribute to RTFM. The interactions between the spin moments of Ce³⁺ and oxygen vacancy arise in accordance with F-centered exchange leading to the ferromagnetic character.^38,39

Saturation magnetization for EFCR (5.37 × 10⁻³) was found to be higher than that of pure CeO₂ (6.56 × 10⁻⁴). An increase in defect concentration has been observed on co-doping iron and europium as supported by our Raman spectra. The trivalent iron and europium doping in the ceria induces additional oxygen vacancies along with the transition metal iron ions contributing towards enhancing the magnetization in the material. The FM behaviour can be attributed to the presence of magnetic ions mediated by oxygen vacancies as reported in ceria doped with transition metal (TM) such as cobalt and nickel.^40,41 In both the studies the behaviour was associated with the FM exchange coupling of F-centres, in which the oxygen vacancies and TM dopants were involved. On doping Fe³⁺, every iron ion has one unpaired electron in low spin state (3d⁵) that can influence in ferromagnetic ordering. So the observed magnetism can be due to presence of F-centers and defects present in (Ce³⁺–V_O–Ce³⁺) complexes where V_O is the oxygen vacancy.⁴²

EFC3 and EFC5 showed higher magnetization among the samples but a linear paramagnetic component was also observed in addition to the ferromagnetic behaviour. The increase in magnetization can be attributed to the higher amount of oxygen vacancies present in these samples as shown from Raman spectroscopy. This higher amount of defects in the lattice enables efficient coupling of cerium and dopants through F-centers.³⁸ The paramagnetic contribution present in the M–H curve for EFC3 and EFC5 was found to be absent in EFC9. Similar paramagnetic influence in the M–H curve has been reported for Fe³⁺ doped CeO₂. The observed paramagnetic effect can be due to the presence of a fraction of interstitial positioned not magnetically ordered iron. The absence of paramagnetic signal in EFC9 may be due to the structural rearrangement of interstitial iron atoms on annealing at 900 °C. A lower magnetization value for EFC7 and EFC9 was noticed than that of EFC5. This can be attributed to the reduction in oxygen vacancies that are responsible for the ferromagnetism. Also these samples showed a reduction in magnetization along with the suppression of paramagnetic signal. This reduction in magnetization can be due to the absence of oxygen vacancy to form F-center through which two nearby Fe atoms can interact via super exchange mechanism.^38,43 From XRD and Raman analysis it can be seen that iron segregates as α-Fe₂O₃. Though segregation lowers the defects but could not destroy the ferromagnetic ordering completely and hence it retains the ferromagnetic character and the magnetization was found to be higher than that of pure CeO₂ and EFCR.

The calculated coercivity value for the samples are given in Table 1. A significant difference was observed in the coercivity value between Fe₂O₃ and EFCR which rules out the presence of any Fe₂O₃ phase in the as prepared condition. Coercivity value was found to be increasing with annealing temperature from 108.56 G (for EFCR) to 516.84 G (EFC9). A transition from soft magnetic to hard magnetic character was observed for high temperature annealed samples (EFC7 and EFC9). Though many reports claim the influence of oxygen valency and transition metal ions such as Co, Mn or Fe is responsible for the observed room temperature ferromagnetism, in our case both oxygen vacancies as well as transition metal dopant together contributed towards the observed room temperature ferromagnetism in the ceria lattice.¹⁶

4. Conclusion

In summary, iron and europium co-doped ceria nanoparticles with porous morphology have been synthesized by combustion method. The XRD analysis indicated that the dopants present in the ceria lattice segregated into α-Fe₂O₃ upon heating above 500 °C. The defect concentration decreased with annealing temperature due to the oxidation of Ce³⁺ to Ce⁴⁺ and conversion of Fe³⁺ present as a dopant to α-Fe₂O₃. Presence of iron ion enhanced the optical absorption in the visible region. Due to the presence in lower concentration, Eu³⁺ is present in high symmetric environment though the higher oxygen vacancy concentration quenches the emission. Ferromagnetic and a transition from soft to hard magnetic character has been observed due to the defect mediated F-centre exchange.

Acknowledgements

Authors are grateful for the financial support provided through Start-up grant (PU/PC/Start-up Grant/2011-12/305) of Pondicherry University. Authors thank the services rendered by Central Instrumentation Facility (CIF), Pondicherry University, Puducherry for the characterization of the samples and Mr P. K. Parida, Microscopy and Thermophysical Property Division, Physical Metallurgy Group, IGCAR for TEM sample preparation and characterization. One of the author Avinash Kumar wish to ackonwldge Mr P. Arunkumar and Mr G. Vinothkumar for the discussions.

References

1. S. Majeed and S. A. Shivashankar, J. Mater. Chem. B, 2014, 2, 5585–5593.
2. Z. X. Yang, B. He, Z. Lu and K. Hermansson, J. Phys. Chem. C, 2010, 114, 4486–4494.
3. N. Mahato, A. Banerjee, A. Gupta, S. Omar and K. Balani, Prog. Mater. Sci., 2015, 72, 141–337.
4. F. Wang and X. Liu, Acc. Chem. Res., 2014, 47, 1378–1385.
5. K. S. Babu, M. Anandkumar, T. Y. Tsai, T. H. Kao, B. Stephen Inbaraj and B. H. Chen, Int. J. Nanomed., 2014, 9, 5515–5531.
6. M. Anandkumar, C. H. Ramamurthy, C. Thirunavukkarasu and K. S. Babu, J. Mater. Sci., 2015, 50, 2522–2531.
7. J. M. Dowding, T. Dosani, A. Kumar, S. Seal and W. T. Self, Chem. Commun., 2012, 48, 4896–4898.
8. D. R. Mullins, Surf. Sci. Rep., 2015, 70, 42–85.
9. A. Kumar, K. S. Babu, A. S. Karakoti, A. Schulte and S. Seal, Langmuir, 2009, 25, 10998–11007.
10. K. S. Babu, J.-H. Cho, J. M. Dowding, E. Heckert, C. Komanski, S. Das, J. Colon, C. H. Baker, M. Bass, W. T. Self and S. Seal, Chem. Commun., 2010, 46, 6915–6917.
11. V. Turkowski, S. Babu, D. Le, A. Kumar, M. K. Haldar, A. V. Wagh, Z. Hu, A. S. Karakoti, A. J. Gesquiere, B. Law, S. Mallik, T. S. Rahman, M. N. Leuenberger and S. Seal, ACS Nano, 2012, 6, 4854–4863.
12. A. Sundaresan and C. N. R. Rao, Nano Today, 2009, 4, 96–106.
13. A. G. El Hachimi, H. Zaari, M. Boujnah, A. Benyoussef, M. Yadari and A. Kenz, Comput. Mater. Sci., 2014, 85, 134–137.
14. S. Y. Chen, C. H. Tsai, M. Z. Huang, D. C. Yan, T. W. Huang, A. Gloter, C. L. Chen, H. J. Lin, C. T. Chen and C. L. Dong, J. Phys. Chem. C, 2012, 116, 8707–8713.
15. S. Y. Chen, K. W. Fong, T. T. Peng, C. L. Dong, A. Gloter, D. C. Yan, C. L. Chen, H. J. Lin and C. T. Chen, J. Phys. Chem. C, 2012, 116, 26570–26576.
16. W. C. Wang, S. Y. Chen, P. A. Glans, J. Guo, R. J. Chen, K. W. Fong, C. L. Chen, A. Gloter, C. L. Chang, T. S. Chan, J. M. Chen, J. F. Leed and C. L. Dong, Phys. Chem. Chem. Phys., 2013, 15, 14701–14707.
17. P. Dutta, S. Pal, M. S. Seehra, Y. Shi, E. M. Eyring and R. D. Ernst, Chem. Mater., 2006, 18, 5144–5146.
18. K. S. Babu, R. Kumar, T. Inerbaev, A. Masunov, A. Schulte and S. Seal, Nanotechnology, 2009, 20, 085713.
19. K. Deshpande, A. Mukasyan and A. Varma, Chem. Mater., 2004, 16, 4896–4904.
20. R. D. Shannon and C. T. Prewitt, Acta Crystallogr., Sect. B: Struct. Crystallogr. Cryst. Chem., 1970, 26, 1046–1048.
21. S. Deshpande, S. Patil, S. Kuchibhatla and S. Seal, Appl. Phys. Lett., 2005, 87, 133113.
22. L. Li, H. K. Yang, B. K. Moon, Z. Fu, C. Guo, J. H. Jeong, S. S. Yi, K. Jang and H. S. Lee, J. Phys. Chem. C, 2009, 113, 610–617.
23. A. Sutka, S. Lagzdina, T. Kaambre, R. Parna, V. Kisand, J. Kleperis, M. Maiorov, A. Kikas, I. Kuusik and D. Jakovlevs, Mater. Chem. Phys., 2015, 149–150, 473–479.
24. P. Arunkumar, S. Preethi and K. S. Babu, RSC Adv., 2014, 4, 44367–44376.
25. M. Guo, J. Lu, Y. Wu, Y. Wang and M. Luo, Langmuir, 2011, 27, 3872–3877.
26. W. H. Weber, K. C. Hass and J. R. McBride, Phys. Rev. B: Condens. Matter Mater. Phys., 1993, 48, 178–185.
27. Z. Dohcevic-Mitrovic, N. Paunovic, M. Radovic, Z. V. Popovic, B. Matovic, B. Cekic and V. Ivanovski, Appl. Phys. Lett., 2010, 96, 203104.
28. D. L. A. de Faria, S. V. Silva and M. T. de Oliveira, J. Raman Spectrosc., 1997, 28, 873–878.
29. Z. Wang, Z. Quan and J. Lin, Inorg. Chem., 2007, 46, 5237–5242.
30. A. S. Reddy, C. Chen, C. Chen, S. Chien, C. Lin, K. Lin, C. Chen and S. Chang, J. Mol. Catal. A: Chem., 2010, 318, 60–67.
31. M. Balestrieri, S. Colis, M. Gallart, G. Schmerber, M. Ziegler, P. Gilliot and A. Dinia, J. Mater. Chem. C, 2015, 3, 7014–7021.
32. D. Avram, C. Rotaru, B. Cojocaru, M. Sanchez-Dominiguez, M. Florea and C. Tiseanu, J. Mater. Sci., 2014, 49, 2117–2126.
33. S. Shi, M. Hossu, R. Hall and W. Chen, J. Mater. Chem., 2012, 22, 23461–23467.
34. X. Ren, L. Tong, X. Chen, H. Ding, X. Yang and H. Yang, Phys. Chem. Chem. Phys., 2014, 16, 10539–10547.
35. A. V. Thorat, T. Ghoshal, P. Carolan, J. D. Holmes and M. A. Morris, J. Phys. Chem. C, 2014, 118, 10700–10710.
36. P. Primus, T. Ritschel, P. Y. Siguenza, M. A. Cauqui, J. C. Hernandez-Garrido and M. U. Kumke, J. Phys. Chem. C, 2014, 118, 23349–23360.
37. S. Phokha, S. Pinitsoontorn and S. Maensiri, Nano-Micro Lett., 2013, 5, 223–233.
38. J. M. D. Coey, A. P. Douvalis, C. B. Fitzgerald and M. Venkatesan, Appl. Phys. Lett., 2004, 84, 1332–1334.
39. M. Radovic, Z. Dohcevic-Mitrovic, N. Paunovic, M. Scepanovic, B. Matovic and Z. V. Popovic, Acta Phys. Pol., 2009, 116, 84–87.
40. Q. Y. Wen, H. W. Zhang, Y. Q. Song, Q. H. Yang, H. Zhu and J. Q. Xiao, J. Phys.: Condens. Matter, 2007, 19, 246205.
41. A. Thurber, K. M. Reddy, V. Shutthanandan, M. H. Engelhard, C. Wang, J. Hays and A. Punnoose, Phys. Rev. B: Condens. Matter Mater. Phys., 2007, 76, 165206.
42. P. C. A. Brito, D. A. A. Santos, J. G. S. Duque and M. A. Marcedo, Phys. B, 2010, 405, 1821–1825.
43. Z. V. Popovic, N. Paunovic and Z. Dohcevic-Mitrovic, J. Phys.: Condens. Matter, 2012, 24, 456001.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c5ra15336k

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