Tuning the conductivity type in a room temperature magnetic oxide: Ni-doped Ga_0.6Fe_1.4O₃ thin films

Abstract

Ni-Doped thin films of the room temperature ferrimagnetic oxide Ga_0.6Fe_1.4O₃ were deposited by pulsed laser deposition and their electronic transport and structural and magnetic properties were studied. The actual insertion of the Ni cations within the Ga_0.6Fe_1.4O₃ structure has been checked by resonant X-ray scattering. A clear extremum is noticed for all properties for the 2% Ni doping: extrema in the crystallographic cell parameters of the films, maximum in the Curie temperature, and maximum in the electric resistivity. We also observed a change of conductivity type for this dopant concentration, from n-type for Ni contents below 2% to p-type for Ni contents above 2%. We explain this behavior by the existence of oxygen vacancies in the pulsed laser deposited Ga_0.6Fe_1.4O₃ thin films, which results in the reduction of some of the Fe³⁺ into Fe²⁺ cations, and n-type conduction via a hopping mechanism. The insertion of Ni²⁺ cations first deals with the presence of oxygen vacancies and reduces the number of n-type carriers in the films, in a compensation-like mechanism. When the number of introduced Ni²⁺ cations dominates the number of oxygen vacancies, conductivity becomes p-type and starts to increase again. We believe that the tunability of the conduction type and magnitude in thin films of a room temperature ferrimagnetic material paves the way towards new all oxide electronic devices.

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1. Introduction

The field of multiferroic and magnetoelectric materials is full of promise in terms of applications in new outstanding devices such as high density four state memories^1,2 or low power consuming magnetoelectric memories.^3,4 Such promises are however actually put off by the low number of room temperature multiferroic and magnetoelectric materials, on the one hand, and by the existence of important leakage currents in thin films of these materials, on the other hand. BiFeO₃ (BFO) is currently the most documented room temperature multiferroic material. Reduction of the leakage current density has been pointed out to be a major challenge that needs to be addressed before BFO-based films can be candidates for integrated microelectronic devices such as storage elements in non-volatile multiferroic memories.^5,6

Conductivity is usually attributed to the presence of oxygen vacancies in the thin films. The major options chosen to reduce the leakage currents in BFO have been annealing in oxygen⁷ or doping with various cations such as Ti,⁸ La⁹ or Cr.¹⁰ The number of oxygen vacancies can also be reduced by implantation¹¹ but this requires technological equipment which is not always available and might hinder the crystallographic quality of the films.

Here, we wish to present a room temperature multiferroic/magnetoelectric material which constitutes a possible alternative to BFO, the gallium ferrite Ga_2−xFe_xO₃ (0.8 ≤ x ≤ 1.4) (GFO), and to address, first, the issue of the reduction of the leakage currents in thin films of this material, and second, the possibility to tune its conductivity type at will, with a view to applications.

GFO was first described as a polar, magnetoelectric ferrimagnetic material in the 60s.^12–15 It has received a considerable renewal of interest the last 10 years after Arima et al. revisited its potential as a room temperature magnetoelectric compound for x > 1.3, with a Curie temperature of 370 K for x = 1.4.¹⁶ GFO does not have a perovskite derived structure as most of the promising multiferroic materials. It crystallizes in the orthorhombic space group Pc2₁n (33) for x between 0.8 and 1.4, the material adopting the β-Ga₂O₃ and α-Fe₂O₃ structures for x below 0.8 and above 1.4, respectively. The Pc2₁n structure of GFO is isostructural to ε-Fe₂O₃ and is the most interesting on the magnetic point of view. It is based on a double hexagonal compact ABAC oxygens stacking forming four types of cationic sites: a tetrahedral Ga1, and three octahedral Ga2, Fe1 and Fe2 sites (Fig. 1). The ferrimagnetic properties of GFO are due to the uncompensated antiferromagnetic ordering of cations in sites Ga1 and Fe1 with cations in sites Ga2 and Fe2, through superexchange interactions through the oxygens. Thanks to a cationic sites disorder, there is a resulting magnetization even for the x = 1.0 composition. However, both the magnetic ordering temperature and the resulting magnetization are maximum for the highest Fe content, i.e. for x = 1.4 (GFO1.4). We will therefore focus our study on GFO1.4 derived compounds, for their Curie temperature is well above room temperature. GFO1.4 lattice constants are a = 0.8765 ± 0.0002, b = 0.9422 ± 0.0002, and c = 0.5086 ± 0.0002 nm.¹⁷

Fig. 1 Projection of the crystal structure of Ga_2−xFe_xO₃ along the c axis (space group Pc2₁n). The image has been produced using the VESTA¹⁸ crystallographic software.

The polarization of GFO has been difficult to prove and quantify in thin films because of important leakage currents. We successfully decreased the conduction of GFO thin films through doping the films with bivalent cations such as Mg²⁺ or Co²⁺.^19,20 In those works, we focussed on the efficiency of the doping in reducing the leakage currents. We tentatively explained the phenomenon by hypothesizing that the conduction was due to a hopping mechanism between Fe²⁺ and Fe³⁺, the Fe²⁺ ions originating from an oxygen substoichiometry in the pulsed laser deposited thin films. The mechanism we suggested to explain this behaviour could be described using the Kröger–Vink notation. During the growth process, some oxygen vacancies are produced in Ga_0.6Fe_1.4O₃ according to

where is a vacancy in the oxygen site with a double positive charge. The consequence of the existence of such oxygen vacancies is the reduction of Fe³⁺ into Fe²⁺

2δe⁻ + 2δFe^×_Fe → 2δFe′_Fe. (2)

The simultaneous presence of Fe²⁺ and Fe³⁺ allows electrical conduction through a hopping process. We therefore suggested that doping with a bivalent cation would allow satisfying the charge unbalance due to the presence of oxygen vacancies without any reduction of Fe³⁺ into Fe²⁺. This would annihilate conduction through hopping and therefore strongly reduce the leakage currents. We have shown a proof of concept of this idea in a previous work in which we doped GFO with Mg²⁺.¹⁹ We indeed observed a reduction of the leakage currents with the Mg²⁺ content in the GFO cell up to 2%. Above this content, the leakage currents started increasing again. We attributed this increase to the introduction of an excess of charges related to an excess of bivalent cations with respect to the existing oxygen vacancies. Our model predicted an n- to p-type conduction change upon increasing the Mg²⁺ content. We however could not show a proof of this hypothesis. This is the subject of the present work. The introduction of Mg²⁺ into the GFO lattice led to a decrease of the ordering temperature and saturation magnetization. The introduction of magnetic Co²⁺ cations allowed the same efficiency in the decrease of the leakage currents, but without any decrease of the ordering temperature. This Co²⁺ doping still caused a decrease in the saturation magnetization though.²⁰ We therefore decided to study the effect of doping GFO thin films with other magnetic bivalent ions, the Ni²⁺ ions. Those cations are similar to the Fe³⁺ cations in size (Table 1), and possess a spin moment of 2 μ_B, while Fe²⁺ and Fe³⁺ have spin moments of 4 μ_B and 5 μ_B, respectively. Because of the ferrimagnetic nature of the compound, it is impossible to predict if the insertion of a cation bearing a lower moment will lead to a decrease or an increase of the total magnetization; it will depend upon the cationic distribution. Ni²⁺ ions are very unlikely to occupy the tetrahedral sites, for a d⁸ ion is stabilized in an octahedral site, but they may occupy any of the three available octahedral sites. We present here a study of the structural, magnetic, and electrical properties of the Ni-doped GFO1.4 pulsed laser deposited films.

Table 1 Cationic radii of the various ions that may be present in the GFO cell (pm)

	Coordination VI	Coordination IV
Ga^3+	62	47
Fe^3+	64 (high spin)	49
Fe^2+	78 (high spin)	63
Ni^2+	69	55

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2. Experimental section

Ni-Doped Ga_0.6Fe_1.4O₃ thin films were deposited on YSZ(100) substrates (Crystal Gmbh) by pulsed laser deposition (PLD) using a KrF excimer laser (λ = 248 nm) with a 10 Hz repetition rate, according to a procedure optimized in a previous work.²¹ The energy density of the laser on the target was tuned to 1 J cm⁻² and the distance between the target and the substrate was fixed at 4.5 cm. The depositions were performed at 800 °C under 1 mbar N₂/O₂ (80%/20%). After the deposition, the samples were cooled down to room temperature under the same gaseous atmosphere. Four sintered ceramics targets of compositions Ga_0.6Fe_1.4−mNi_mO₃ with m = 0, 0.01, 0.04, and 0.10 were used for the depositions, corresponding to global percentages of substitution of m/(0.6 + 1.4 − m + m) = 0, 0.5, 2.0, and 5.0, respectively. The targets were elaborated from the solid state reaction of Fe₂O₃, Ga₂O₃ and NiO powders according to a procedure whose optimization is described elsewhere.²² A stoichiometric mixture of the reacting powders was milled in an attritor in an ammoniac solution (pH = 9) for one hour. The solution was then placed in a drying oven until the full evaporation of the liquid part and the resulting powder was manually ground. An organic binder (polyvinylic alcohol) has been systematically added to improve the mechanical behaviour of the samples. The powders were finally uniaxially compacted into pellets and sintered at 1450 °C for 20 hours in air. The crystallographic structure of the deposited thin films was characterized by X-ray diffraction (XRD) using a Rigaku Smartlab diffractometer equipped with copper radiation (K_α = 0.154056 nm). The films thicknesses, determined from X-ray reflectivity, were of ca. 80 nm. Energy dispersive spectroscopy (EDX) coupled to a Scanning Electron Microscopy technique (JEOL 6700 F) was used to determine the composition of the samples and study the surface of the films. The analyses were performed at 5 keV, which ensures that the whole EDX signal can only originate from the GFO film and not from the substrate. SEM was also used to characterize the surface morphology. The composition of the films was checked by inductively coupled plasma atomic emission spectroscopy (ICP-AES, in the “Plateforme Analytique des Inorganiques”, Strasbourg) after microwave assisted dissolution of approximately 3 × 3 mm² samples in HNO₃ for trace analysis (TraceSELECT®, Fluka) (see ESI†). A deviation of less than 3% was observed between the expected composition and the measured one for all samples. The crystallographic structure of the films was observed using Transmission Electronic Microscopy (TEM) in cross views with a JEOL 2100 FCs corrected for spherical aberration at the condenser. Observations in the high angle annular dark field (HAADF) mode coupled to EDX allowed checking of the homogeneity of the Ni-doping throughout the films. Cross sections of the thin films were prepared by the focused ion beam (FIB) technique at the Service of Microscopy (D. Troadec) of the IEMN Institute in Lille, France. The valence state of Ni was checked using X-ray photoelectron spectroscopy (XPS). The measurements were carried out in an ultrahigh vacuum (UHV) setup equipped with a 150 nm VSW ClassWA hemispherical electron analyzer. A non-monochromatized Al Kα X-ray source (1486.6 eV; anode operating at 250 W) was used as the incident radiation. The binding energy scale was calibrated based on the adventitious carbon C 1s peak at 285 eV. Prior to individual elemental scans, a survey scan was taken to detect all the elements present (see ESI†). The actual insertion of Ni in the GFO cell was checked by resonant X-ray diffraction.²³ The intensity of three independent Bragg reflections of the GFO Pc2₁n structure were recorded as a function of the energy of the scattered photons crossing the atomic Ni L_2,3 edge (8.33 keV). The resonant X-ray diffraction experiments were performed on the CRG-BM02 beamline at the European Synchrotron Radiation Facilities (ESRF) in Grenoble using photon energies between 8.285 and 8.405 keV. The measured intensities were corrected for fluorescence. The magnetic properties of the films were determined using a superconducting quantum interference device magnetometer (SQUID MPMS XL, Quantum Design). Field-cooled (FC) and zero-field-cooled magnetizations were measured between 5 and 400 K in a field of 50 Oe applied parallel to the films. Hysteresis loops were measured at 300 and 5 K for magnetic fields comprised between plus and minus 7 T. Resistivity and Hall effect measurements were performed on the films by the van der Pauw method on a home built high temperature high impedance set-up, operating in the temperature range 80–880 K, using Ti/Au sputtered electrodes for ohmic contact. Seebeck effect measurements were carried out using home constructed equipment at 300 and 500 K, in order to confirm the nature of the carrier types.

3. Results and discussion

The films are well crystallized and oriented along the b direction, as demonstrated by the X-ray diffraction patterns of the films in θ–2θ mode (Fig. 2). Φ scans around reflections (057) and (240) of GFO and YSZ, respectively, show the existence of 6 in-plane variants, at 30° from each other. The orientations of the GFO variants are due to the different matching possibilities of GFO(0k0) onto YSZ(100): c_Ga0.6Fe1.4O3 = 0.5086 nm ≈ a_YSZ = 0.5139 nm (1% mismatch) on the one hand, and (1% mismatch), on the other hand.²⁴ The angle between c_Ga0.6Fe1.4O3 and the diagonal of the ac plane is , which explains the existence of six GFO variants located every 30°.²¹ The cell parameters were calculated from the θ–2θ measurement of the (040) and reciprocal space mappings of the (062) and (570) reflections. While the in-plane a and c parameters are lower than observed for the same Fe/Ga ratio in bulk,¹⁷ the out-of-plane b parameter is higher (Fig. 3). The mismatch between YSZ(100) and GFO(010) along the YSZ[010]//GFO[001] direction cannot account for the low c parameters observed. Indeed a_YSZ = 5.139 Å is higher than 5.086(2) Å, observed for c in bulk Ga_0.6Fe_1.4O₃ and this would therefore impose a tensile stress rather than a compressive one. The different thermal expansion coefficients between the substrate and the layer cannot account for the smaller in-plane parameters observed for GFO thin films either. The samples are deposited at 800 °C and then cooled down to room temperature. One could therefore expect a matching of c_GFO onto a_YSZ at 800 °C. The thermal expansion coefficient of YSZ being α_YSZ = 9 × 10⁻⁶ K⁻¹,²⁵ its cell parameter at 800 °C is 5.151 Å. Assuming a_GFO adopts this value when deposited, upon cooling, considering a thermal volume expansion coefficient of α_GFO = 2.6 × 10⁻⁵ K⁻¹,²⁶ it would become, if one considers that this thermal coefficient is isotropic, This value is higher than the experimentally observed one, even for undoped samples. The elongation of the GFO cell along its b axis is therefore more likely to be related to the existence of oxygen vacancies within the material. Some increases in the cell parameters have already been observed in materials like perovskites and explained by the oxygen vacancies resulting in the partial reduction of some of the cations, which having more electrons in their reduced form, have a bigger radius.^27–29 The cell parameters which are the closest to those observed for the bulk are obtained for the 2% Ni doping.

Fig. 2 (a) X-ray θ–2θ diffractograms of the Ni-doped and undoped films and (b) Φ scans around reflections (570) and (240) of GFO and YSZ, respectively, for the undoped GFO film, as representative of all films.

Fig. 3 Evolution of the (a) a, (b) b, and (c) c cell parameters with the Ni content in GFO.

The insertion of the Ni atoms within the crystallographic structure of GFO was checked by means of resonant diffraction at the Ni K edge. Three independent Bragg reflections of the GFO structure, (110), (211), and (510), were scanned. Anomalous diffraction could be observed in all cases (Fig. 4), after subtracting the fluorescence. It originates from a modification of the diffracted intensity due to resonant elastic X-ray scattering processes involving interactions between the X-ray beam and the Ni atoms in the sample and is a proof that Ni atoms contribute to the structure factor for this crystallographic reflection. This is therefore an unequivocal proof of the insertion of the Ni atoms within the GFO crystallographic structure.

Fig. 4 Intensity variations of the (a) (110), (b) (211), and (c) (510) reflections as a function of the photon energy at the Ni K edge. The different shapes of the variations are related to the phase of the Ni atoms complex structure factor, relative to other atoms in the crystal structure.

The homogeneity of the Ni-content over the entire thickness of the films was checked using EDX coupled with TEM. The thin films were observed in cross views in the STEM HAADF mode (Fig. 5). No accumulation of Ni at the surface or at the interface between the substrate and the films was observed. The value of the Ni content was confirmed for all compositions, within the error bar of EDX (3%).

Fig. 5 Cross view composition of the 2% Ni-doped GFO thin film determined from EDX analysis in STEM HAADF mode (C indicates a carbon layer deposited during the FIB preparation of the sample).

High resolution cross view TEM images show a columnar growth of the films with ca. 10–20 nm wide crystallites (Fig. 6). They confirm the Pc2₁n crystallographic structure of the samples and allowed observation of the contrast from the different variants.

Fig. 6 TEM cross view image of the 2% Ni-doped GFO thin film deposited on YSZ(001). GFO crystallites are visible in the (1) [100]_Pna21 and (2) [310]_Pna21 zone axes. The simulations are performed with the JEMS software package (P.A. Stadelmann) for a JEOL 2100 FCs corrected for spherical aberration at the condenser (defocus = 11 nm, thickness = 30 nm).

The Ni 2p_3/2 XPS peak (Fig. 7) shows a well-resolved two peak structure with a main peak at 855.2 eV and a satellite structure shifted 6.5 eV at the high binding energy side, which are characteristic of Ni in the +2 valence state.^30,31 In addition the Auger parameter (α) calculated using the peak maxima of the Ni 2p_3/2 photoelectron and L₃M₄₅M₄₅ Auger peaks is found to be 1698.0 ± 0.2 eV, within the range previously proposed for Ni²⁺.³² Overall, XPS results provide strong evidence that nickel cations are in the +2 valence state.

Fig. 7 XPS Ni 2p_3/2 spectrum for the 5% Ni-doped GFO thin film. The Auger parameter, (α), calculated by the Ni 2p_3/2 and Ni LMM peak maximum positions is included in the graph.

The Curie temperature of the GFO films was measured as the inflexion point of the field-cooled magnetization curve measured in a 50 Oe magnetic field (see ESI† for field-cooled magnetization curves of all the samples, measured in a 50 Oe magnetic field applied parallel to the films). It presents a maximum for the 2% Ni content (Fig. 8). The magnetic properties in GFO are governed by the superexchange interactions between neighbouring magnetic ions.³³ They will therefore be strongly affected by the M–O–M bond angles, where M is a magnetic cation, hence by the cationic ordering. In addition to the cationic site distribution, the distortion of the polyhedra also plays an important role. The 2% Ni-doped sample is the one for which the cell is the less distorted when compared to bulk (Fig. 3). The maximum in the Curie temperature is therefore probably due to a maximum in the orbitals overlapping. For lower doping, the cell is distorted by the presence of oxygen vacancies inducing the presence of Fe²⁺. One can assume that for the 2% Ni doping, all Fe²⁺ are replaced with Ni²⁺. This configuration is the closest to bulk GFO, which is free of oxygen vacancies. For higher doping, additional Ni²⁺ cations are introduced and replace Fe³⁺, which leads to another distortion of the cell.

Fig. 8 Evolution of the Curie temperature of the Ni-doped GFO films with the Ni content.

Hysteresis loops measured at 300 K (Fig. 9 and ESI†) indicate that the samples are more easily magnetized within the plane of the films (ac plane) than out of plane (b direction). This is consistent with the easy and hard magnetization axes observed in bulk: c is the easy axis while b is the hard axis.¹⁶ Because of the existence of 6 in-plane variants, the distinction between the different easy directions in plane, located every 30°, is not possible in the SQUID, in which such a precision in the positioning of the sample is not achievable. The room temperature saturation magnetization increases with the Ni content from 67 to 90 emu cm⁻³, for non-doped and 5% Ni-doped GFO, respectively (Fig. 10). Magnetization in GFO results from the uncompensated antiparallel ordering of two sublattices constituted by the Ga1 and Fe1 sites, on the one hand, and Ga2 and Fe2, on the other hand. It strongly depends upon the cation distribution on the four different sites.³⁴ The cations contributing to the samples magnetization are Fe³⁺ (5 μ_B), Fe²⁺ (4 μ_B), and Ni²⁺ (2 μ_B) – assuming a quenching of the orbital moment.

Fig. 9 Hysteresis loops of the undoped GFO thin film, as representative of all samples (see ESI† for a comprehensive presentation of the measurements), measured at 300 K in parallel (black filled squares) and perpendicular (red hollow squares) (enlargement in inset).

Fig. 10 Evolution of the room temperature saturation magnetization of the Ni-doped GFO films with the Ni content.

The total moment for pure GFO1.4 can be evaluated assuming a cationic distribution similar to the bulk one, without considering any oxygen vacancy. In that case, one expects 0.33, 0.79, 0.97, and 0.82 Fe cations in the Ga1, Ga2, Fe1, and Fe2 sites, respectively.³⁵ The sites Ga1 + Fe1 would therefore bear (0.33 + 0.97) × 5 = 6.5 μ_B, while the Ga2 + Fe2 sites represent (0.79 + 0.82) × 5 = 8.05 μ_B, yielding a total of 1.55 μ_B for 2 formula units. The Ni cations only bear 2 μ_B when compared to Fe cations which have 4 (Fe²⁺) or 5 μ_B (Fe³⁺). The increase in the saturation magnetization observed for increasing Ni content can therefore be explained by the positioning of these cations in the Fe1 or Ga1 sites. Ni²⁺, being a d⁸ cation, is strongly stabilized in octahedral sites. They will therefore most probably be located in the Fe1 sites.

Fig. 11 shows the temperature dependence of resistivity ρ(T) for Ni-doped GFO thin films in the 300–500 K temperature range. The measurement temperature was limited to 500 K, to avoid compositional transformation of the samples at higher temperatures. The resistivity of the films increases with decreasing temperature, indicating a semiconducting behaviour of the conductivity. Room temperature resistivity (ρ) values for different levels of Ni content are summarized in Table 2. Room temperature resistivity values vary between 5.5 × 10³ and 7.1 × 10⁴ Ω cm. The maximum value is obtained for the 2% Ni-doping.

Fig. 11 Temperature evolution of the resistivity of the films for various Ni contents.

Table 2 Room temperature resistivity values for different Ni contents in the Ni-doped GFO films

	Resistivity (ρ, Ω cm)
	300 K	500 K
Undoped GFO	5.0 × 10^4	2.8 × 10^3
Ni: 0.5%	3.2 × 10^4	4.8 × 10^3
Ni: 2%	7.1 × 10^4	1.3 × 10^4
Ni: 5%	5.5 × 10^3	1.4 × 10^2

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For the samples with 0–2% Ni the resistivity temperature dependence exhibits hopping conduction mechanism (resistivity does not changes significantly versus temperature) up to around 450 °C. Above 450 °C, the ρ(T) slope changes, which can be an indication of a band activation conduction mechanism. For the 5% Ni doped sample, resistivity is strongly temperature dependent, and an 0.24 eV activation energy of conductivity σ(T) can be estimated from the ln(1/σ) = ln(ρ) versus 1/T slope. Such behaviour for resistivity versus temperature can be explained by the substitution of a hopping conduction mechanism with a delocalized hole conduction.

Moreover, Seebeck and Hall effects measurements (Hall effect has been performed at 1.2 T when the magnetization of the samples is saturated) allowing unambiguous determination of the nature of the carriers, and of their concentration (Table 3). The nature of the carriers varies with the Ni content: for doping contents below 2%, the majority of carriers are electrons, while for those with Ni content higher than 2%, they are holes. The number of carriers is a minimum for samples with Ni content between 0.5 and 2% (Fig. 12). This confirms our assumption concerning the mechanism responsible for conduction in the films. The Hall mobilities are less than 1 cm V⁻¹ s⁻¹ for all samples, due to strong carrier localization. The highest resistivity observed for the 2% Ni doped sample, which has a higher carrier concentration then the 0.5% Ni doped sample, is explained by its lower mobility.

Table 3 Carriers types and concentrations at 300 and 500 K for different Ni contents in the GFO films, determined from Hall effect measurements in a 1.2 T magnetic field

Sample	Carriers concentration (cm^−3), (p): holes, (n): electrons
	300 K	500 K
Undoped GFO	N/A	(n) 2.80 × 10^16
Ni: 0.5%	N/A	(n) 1.20 × 10^15
Ni: 2%	N/A	(p) 3.60 × 10^15
Ni: 5%	(p) 1.2 × 10^15	(p) 1.04 × 10^17

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Fig. 12 Evolution of the carriers’ type and concentration with the Ni content in GFO thin films at 500 K.

At room temperature, the Hall effect is only valid for the 5% Ni doped sample, when hopping conduction gives way to band conduction. The substitution of Fe with Ni indeed leads to the creation of holes according to

where Ni′_Fe stands for Ni in the Fe site with an apparent negative charge and h˙ represents a hole. These holes can be recombined with electrons according to

e⁻ + h˙ → 0. (4)

The substitution of Fe with Ni therefore first leads to a decrease of the leakage currents thanks to the electrons–holes recombination mechanism. When the substituted Ni content m is too high compared to the number of oxygen vacancies δ, with Ni actively playing the role of acceptor, supernumerary holes are eventually created and the material becomes a p-type conductor. Since each oxygen vacancy generates 2 electrons (see eqn (1)), this is the case when m > 2δ.

4. Conclusions

We have observed the evolution of structural, electric and magnetic properties of thin films of a ferrimagnetic oxide, Ga_0.6Fe_1.4O₃, upon Ni²⁺ doping, for Ni content up to 5%. Resonant X-ray diffraction was used to confirm the actual insertion of the Ni cations within the GFO cell. The 2% Ni doping is clearly a particular value. The insertion of 2% Ni shifts the cell parameters of GFO thin films towards the values observed for the bulk, while they are different for Ni contents below and above this value. All the films are ferrimagnetic with a net room temperature saturation magnetization between 65 and 90 emu cm⁻³, for undoped and 5% Ni doped GFO samples, respectively. The Curie temperature reaches a maximum for the 2% Ni content, probably because its structure is the closest to the bulk one, allowing optimized orbital overlapping and superexchange interactions. A minimum conductivity was observed for the 2% Ni-doped sample, which shows hole conductivity. A conductivity type inversion, from n to p, is clearly evidenced when going from low to high Ni doping content. A conductivity type inversion in an another iron-based magnetic oxide, FeTiO₃, has already been reported by us.³⁶ The switching from p to n type necessitated a post annealing. In the present work the “bipolarity” is directly achieved by monoatomic doping, with no additional post-treatment. This can thus be considered as an improved method to tune conductivity. In summary, we believe that the semi-conducting behaviour of room temperature ferrimagnetic Ni-doped Ga_0.6Fe_1.4O₃ epitaxial thin films, with tunable type of conductivity, paves the way towards new all oxide electronic devices.

Acknowledgements

This work was done with the financial support from the CNRS, the Ministère de l’Enseignement Supérieur et de la Recherche (Paris), the international ANR DFG Chemistry project GALIMEO #2011-INTB-1006-01, and the laboratory of excellence Nanostructures in Interaction with their Environment (LabEx NIE 11-LABX-0058-NIE).

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Footnote

† Electronic supplementary information (ESI) available: Compositions of the Ni-doped films determined by ICP-AES, field-cooled magnetization curves and hysteresis curves of all films, XPS survey scan spectrum of the 5% Ni-doped GFO film together with the deduced stoichiometry. See DOI: 10.1039/c6ra01540a

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