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Electric and Magnetic Properties of ALD Grown BiFeO Films Benoît Marchand, Pasi Jalkanen, Vladimir Tuboltsev, Marko Vehkamäki, Manjunath Puttaswamy, Marianna L Kemell, Kenichiro Mizohata, Timo Hatanpää, Alexander Mikhailovitch Savin, Jyrki Räisänen, Mikko Ritala, and Markku Leskelä J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.5b11583 • Publication Date (Web): 17 Mar 2016 Downloaded from http://pubs.acs.org on March 21, 2016

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The Journal of Physical Chemistry

Electric and Magnetic Properties of ALD Grown BiFeO3 Films

Benoît Marchand*, Pasi Jalkanen*,Vladimir Tuboltsev, Marko Vehkamäki, Manjunath Puttaswamy, Marianna Kemell, Kenichiro Mizohata, Timo Hatanpää, Alexander Savin, Jyrki Räisänen, Mikko Ritala, and Markku Leskelä

Dr. B. Marchand, Dr. P Jalkanen, Dr. V. Tuboltsev, Dr. K Mizohata, Prof. J. Räisänen University of Helsinki, Department of Physics, Division of Materials Physics, P.O. Box 43, Helsinki FI-00014, Finland E-mail: [email protected], [email protected] Dr. M. Vehkamäki, Dr. M. Puttaswamy, Dr. M. Kemell, Dr. T. Hatanpää, Prof. M. Ritala, Prof. M. Leskelä University of Helsinki, Department of Chemistry, Laboratory of Inorganic Chemistry, P.O. Box 55, Helsinki FI-00014, Finland Dr. A. Savin Aalto University, School of Science, O.V. Lounasmaa Laboratory, Low Temperature Laboratory, P.O. Box 15100, Espoo FI-00076, Finland Keywords: multiferroics, domain wall pinning, polycrystalline, magnetic properties, strain

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Abstract The magnetization and electric polarization in thin bismuth ferrite films (BFO) films have been under extensive study for high technological potential of single phase multiferroic materials. Surpassing the antiferromagnetic nature and weak magneto-electric coupling of bulk BFO has required highly specialized substrates and epitaxial growth methods so far. Polycrystalline single phase multiferroic BFO (50-500 nm thick) films were grown by Atomic Layer Deposition (ALD) on technologically simple Pt/SiO2/Si substrates. The BFO films were found to exhibit strong saturating ferromagnetism and coercivity at temperatures ranging from cryogenic to room temperature even with 500 nm thick layers, a property which cannot be obtained with thick epitaxial films or bulk BFO. The magnetization mechanism was associated with magnetic domain wall dynamics and collapsing of the helimagnetic spin modulation. The electric properties were found to be strongly dependent on the film thickness. The film crystallization, composition and chemical state have been analyzed by various techniques. The magnetic and ferroelectric properties were determined by using a SQUID magnetometer and a ferroelectricity tester. The results of the work indicate clearly that ALD technique offers an efficient way for synthesis of polycrystalline BFO films and for tailoring their electro-magnetic properties.


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Introduction Bismuth ferrite BiFeO3 (BFO) has attracted a great deal of attention since it is the only known single phase (rhombohedral) multiferroic perovskite exhibiting coupled magnetic (G-type antiferromagnetic) and ferroelectric ordering with very high transition temperatures, viz. Curie and Néel temperatures being 1103 and 643 K, respectively.1-4 According to the conventionally accepted picture, in bulk BFO a small net magnetization due to the canting of adjacent Fe3+ spins is modulated by a long-range superstructure superimposed on the spin arrangement. The net magnetic moment is cancelled out due to the magnetization rotation following a cycloid with ∼62 nm periodicity propagating along the [110] direction (in the hexagonal representation).5,6 Recent studies, however, indicate that single crystals, mesoporous structures, various low-dimensional BFO systems, such as nanoparticles and thin films may exhibit unexpected structural and magnetic properties.7-11 The straightforward explanation of such properties is associated with the destruction of the modulation cycloid due to the spatial confinement in BFO nanomaterials. Also other reasons have been shown to affect multiferroicity in BFO such as stress induced phase transformation, domain boundaries, oxygen octahedron rotation, vacancy formation, Bi and Fe displacement.12-16 On the side of BFO’s multiferroicity, the photosensitive properties have gathered notable interest for photovoltaic, energy and catalytic applications. For instance, photoferroic devices, photocatalytic activity, and perovskite solar cells in general are part of the emerging technologies in which BFO, as a lead free material, is in central position. Aiming at possible BFO based applications, different approaches synthesis of high quality BFO films have been demonstrated, significant efforts being directed towards modifying and tuning of the materials photo sensitive magnetic and electric properties through crystallinity, strain engineering, porosity, ferroelectric 3 ACS Paragon Plus Environment


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and magnetic domain boundary control.8,14,17-25 Multiferroic BFO systems produced with different nano-motifs offer a diverse material platform enabling solutions and devices in micro- and nanoelectronics with multifunctionality embedded at the material level.14 This implies high demands in production methods and techniques compatible with the mass-production paradigm and enabling fabrication of BFO systems with precise control over the stoichiometry and critical dimensions. Various thin film deposition techniques have been studied during the last decade including pulsed laser deposition, radiofrequency magnetron sputtering, chemical solution deposition and chemical vapour deposition.17-21 Only recently first publications have appeared in the literature reporting on the attempts to employ Atomic Layer Deposition.22-24,26,27 ALD is well known to be an advantageous technique offering unprecedented degree of control over crystallinity, thickness and conformity in high aspect ratio structures.28 Developing new strategies, recipes and precursors for ALD of BFO thin films with tailored properties would certainly boost the application of this perovskite in microelectronics and photovoltaic applications.29-31 The weak magnetoelectric (ME) coupling between BFO’s magnetic and electric order has hindered efficient multiferroic applications. This might be changed by the recent implications that a strong magnetoelectric coupling can be established in ordered BFO nanocrystal configurations having some flexibility and strain to break rhombohedral symmetry. Assuming that mesoporosity is one factor determining efficient multiferrocity, ALD is well suited for covering ordered porous template/electrode structures. The ALD method provides thus an alternate method for highly controlled deposition of BFO over various geometries that can be applied for enhanced FME coupling. In our previous work, magnetic properties of 50 nm thick films were studied and BFO films with significant magnetization comparable to epitaxial films were demonstrated.8,31 The aim of the present work was to evaluate the electric and magnetic properties of BFO thin films of 50 – 4 ACS Paragon Plus Environment

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500 nm thickness grown by ALD on Pt electrodes. The multiferroic polycrystalline BFO/Pt/SiO2/Si stacks are technologically highly attractive being fully compatible with already available manufacturing processes and microfabrication techniques.

Experimental Section Sample Preparation All BiFeO3 films were grown by ALD on Pt/SiO2/Si substrates by using bismuth(III)2,3-dimethyl-2butoxide, iron(III)tert-butoxide and water precursors in an ASM Microchemistry F-120 flow-type reactor. Film deposition was done at 423 K (150 °C), known to be in true ALD regime for both FeOx and BiOx binary oxide processes. A decomposition limit of 170 °C has been observed for iron(III)tert-butoxide, and bismuth(III)2,3-dimethyl-2-butoxide is stable at least up to 200 °C.32 The process used herein is mainly limited by the Fe precursor, which is a solid and has limited thermal stability. The bismuth(III)2,3-dimethyl-2-butoxide source used here is conveniently a liquid up to the evaporation temperature of ca. 60-70 °C. The films were grown as binary oxide layers, with growth rates of 0.4 Å/cycle for BiOx layers and 0.2 Å/cycle for FeOx layers. Every FeOy layer consisted of 200 cycles of iron(III)tert-butoxide and water, and every BiOx layer of 180 cycles of bismuth(III)2,3-dimethyl-2-butoxide and water, resulting in 4 nm FeOx and 7.2 nm BiOx layers. This 4 nm + 7.2 nm sequence, empirically found to result in the correct stoichiometry, was repeated until desired film thicknesses were reached. Due to the low process temperature, the layers were completely amorphous in the as-deposited state. All films were processed so that a 200 cycles (4nm) of FeOx layer was deposited both at the start of the process and as the final top layer to limit Bi outdiffusion during high temperature annealing. The pulse durations were 0.4 s for the 5 ACS Paragon Plus Environment


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bismuth precursor, 0.5 s for the iron precursor and 0.5 s for water. The chamber was purged with nitrogen after each precursor cycle for 1.5 s after the metal precursor pulse and for 3 s after the water pulse. Samples with grown laminar layers were annealed at 773 K (500 °C) for one hour in nitrogen ambient at atmospheric pressure, resulting in layer interdiffusion and BFO phase formation. The annealing resulted in a volume shrinkage of ca. 17 %, for example a 60 nm asdeposited amorphous film formed a 50 nm polycrystalline BFO film. Sample Characterization Formation of BiFeO3 phase was assured by Grazing Incidence X-Ray Diffraction (GI-XRD) patterns recorded employing a PANanalitical X’pert Pro diffractometer (Cu Kα line λ = 1.5405 Å). To characterize the films microstructure, morphology and Bi/Fe ratio, Scanning Electron Microscopy and Energy Dispersive X-ray Spectrometry were performed using a Hitachi S-4800 SEM equipped with an Oxford INCA 350 spectrometer. A cross sectional sample was also prepared by a Quanta 3D 200i dual Focused Ion Beam (FIB) and analyzed by SEM. To ensure the stoichiometry by two independent techniques, Time-of-flight Elastic Recoil Detection Analysis (Tof - ERDA) has been performed using 50 MeV 127I 9+ ions. Finally, XPS was employed to evaluate the chemical state of Fe, Bi and O in the grown films. For this, the samples were irradiated by X-rays emitted from a Mg source (Kα line, λ = 9.890 Å) and photo electrons were analyzed by an Omicron Argus Spectrometer with a pass energy of 20 eV. Electron flooding was applied in order to avoid surface charging. Energy scale of the XPS spectra was calibrated using C 1s spectral component of adventitious carbon (284.8 eV).33 Electrical properties of the grown BFO films were evaluated using a probe station connected by coaxial cabling to a Radiant Multiferroics Premier II tester. BeCu contact needles ensured a firm contact with the samples under test.


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Ferroelectricity Measurements Ferroelectric properties were measured by using standard remanent polarization (Pr) and PUND (dPr) testing procedures. The test capacitors were fabricated by depositing 30 nm thick Pt electrodes on top of the grown films, surface area varying from 0.0007 to 0.02 cm2. The delay between the initial switching on/off pulse and the following measurement pulse was 1 Hz both for Pr and dPr testing modes. The optimal pulse frequency range resulting in strongest polarization was found to be between 10 –1000 Hz depending on the film leakage and polarization rates. The same experimental set-up was used for measurements of IV characteristics and capacitance.

Magnetic Measurements In order to characterize the magnetism in the grown films, magnetization was measured as a function of applied magnetic field and temperature. A Superconducting Quantum Interference Device (SQUID) magnetometer (Quantum Design MPMS-XL7) was employed enabling magnetization measurements in varying fields up to 70 kOe and at temperatures ranging from 1.8 K to 400 K. During the measurements, the magnetic field was applied perpendicularly to the films. Zero field-cooled magnetization was measured by, first, cooling down to 1.8 K in zero magnetic field and, then, warming up while measuring with the magnetic field switched on. In the fieldcooled measurements, samples were pre-cooled in the magnetic field. In order to assess dynamic magnetic properties of the films, the AC susceptibility was evaluated in a periodic field of 1 Oe centered at H = 0 Oe and oscillating at a frequency of 1-100 Hz.



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Results & Discussion Microstructure Figure 1 a) shows a characteristic x-ray diffraction pattern of the ALD grown 500 nm thick BiFeO3 film after annealing. Based on XRD, the near-surface part of the film is hexagonal BFO of R3c group with distorted rhombohedral perovskite structure without any major peak shifts compared to bulk [Further details of the crystal structure investigation(s) may be obtained from the Fachinformationszentrum Karlsruhe, 76344 Eggenstein-Leopoldshafen (Germany), on quoting the depository number CSD-99-000-0003]. There are, however, small peak shoulders at larger angles higher d-values. The depth range from the bottom Pt electrode up to ∼ 200 nm of the 500 nm thick BFO film was investigated with HR-TEM (see Figure 1 b)).


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Figure 1. A) GI-XRD (2θ) pattern of a 500 nm thick ALD grown BiFeO3 film. B) high resolution image of the selected grain showing fringing patterns.

The lattice period was found to differ from the hexagonal BFO bulk value of 3.96 Å. Close to inplane direction distances in the range 3.88 - 3.96 Å were observed, implying compressive strain, and elongated out-of-plane distances in the range 3.96- 4.03 Å also matching a compressive load. For comparison, Fu et al.34 observed an irreversible thickness dependent phase transition for epitaxial layers on STO/LSMO substrates under heavy compressive strain, in which the BFO nearest the substrate interface adopts the tetragonal-like phase (with corresponding in-plane distance lattice of 3.77 Å) and at film thicknesses above 220 nm fully relaxed rhombohedral BFO is 9 ACS Paragon Plus Environment


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formed. In our case, the strain is distributed more randomly in the polycrystalline matrix, but similar to the work of Fu et al. the near surface part of the film almost completely relaxed, as shown by XRD. No evidence of tetragonal-like BFO was found.25 Parasitic phases frequently reported in BFO produced by other deposition techniques are those of Bi2Fe4O9 (Fe rich) and Bi25FeO40 (Bi rich) as well as iron oxides α-Fe2O3, γ-Fe2O3 and Fe3O4.35-38 A small shoulder at ∼30° seen on the diffraction pattern can be attributed to strained BFO near the Pt electrode, and toa Bix-Pt1-x phase due to Bi diffusion into the Pt electrode under the BFO layer.39 Elastic recoil detection analysis (ERDA) confirmed the expected stoichiometry of the BFO film showing (20 ±2)% of Bi, (20±1)% of Fe and (60±1)% of O. Energy dispersive spectroscopy (EDS) measurements agreed well with the ERDA results yielding Bi/Fe ratio of 1.10±0.04 implying also enhanced Bi concentration presumably due to the formation of BixPt1-x. In order to verify this hypothesis, another sample without the platinum electrode was prepared and its composition was found to correspond to Bi/Fe ratio of 0.97±0.04. The film morphology is presented in Figure S1 of the supporting information (SI). From a detailed analysis of transmission electron microscopy (TEM) images it follows that the grown BFO films were nano-crystalline with grain sizes ranging from few tens of nm up to 60 nm with an average size of ∼40 nm. Overall the properties between 50 and 500 nm thick films turned out to be essentially similar in terms of x-ray diffraction (XRD) peak width and EDS results. Crystal sizes of the 50 nm films were somewhat smaller resulting in smaller average grain size and narrower size distribution compared to 500 nm thick films.31


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Chemical State The Figure 2 shows x-ray photoelectron spectroscopy (XPS) of the ALD grown films. Peak fittings were performed by using the well-known CASA XPS software.40

Figure 2. XPS analysis of a BiFeO3 film grown by ALD. The lines show fittings for the Fe 2p3/2 state and Shirley background A) and correspondingly, for the Bi 4f in B).

Two peaks for Bi 4f5/2 and 4f7/2 are centered at 158.9 and 164.2 eV in agreement with the values reported in the literature for Bi3+.41 The O 1s and C 1s states are split in to three spectral components (see Figure S3 in SI) due to lattice O2- and various carbohydrate related valence states.41-42 The positions of the Fe 2p3/2, satellite of the Fe 2p3/2 and Fe 2p1/2 peaks are at 710.4, 718.4 and 724.4 eV, respectively. The iron spectral component is more complicated to interpret as both Fe2+ and Fe3+ exhibit asymmetric peaks which cannot be fitted with a single component.33,43,44 A batch of peaks associated with Fe2+ and Fe3+ has been proposed in the literature.33 The recommended peak fitting procedure utilizes only the total area under all peaks and the FWHM of the main contribution as free parameters, constricting all other parameters (All other FWHM are defined as a function of the FWMH of the main contribution). The fitting procedure was applied to the Fe 2p3/2 peak after subtraction of a Shirley background and the result turned out to be in 11 ACS Paragon Plus Environment


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excellent agreement with the Fe3+ component without any need to involve Fe2+ (see Figure S2 in SI).33

Ferroelectric Properties The ferroelectric properties of the ALD grown BFO were analyzed by electric characterization of samples with different film thickness, viz., 50, 200 and 500 nm. Polarization, capacitance and leakage current were found to exhibit dependence on the film thickness of the polycrystalline films (see Figure 3). The measured capacitance values were 0.015 µF cm-2, 0.170 µF cm-2 and 0.370 µF cm-2 for the 50, 200, and 500 nm films, respectively. As it follows from Figure 3a), the BFO films turned out to be rather good dielectrics. Low leakage currents of the order of ∼10-8-10-6 Acm-2 were measured for 50 and 200 nm films. For 500 nm films the experimental value was ∼10-4 Acm-2 which is comparable with those reported for BFO films produced by different techniques.22 The films show a trend of decreasing leakage current with decreasing film thickness. The leakage current reduces by 2 orders of magnitude as the thickness decreases to 200 nm and by an additional 2 orders of magnitude for the 50 nm films. The overall strain in the thick films cause more prominent distortion over the film depth compared to the thin films. Strained films are reported to possess 1-2 orders of magnitude higher leakage current compared to their relaxed counterparts.45 Strain relieving boundaries, hillocks and delamination sites are present in the 500 nm films (see Figure S1 in SI) resulting in more conductive interfaces and paths through the films.46


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Figure 3. Electric properties of ALD grown BiFeO3 films: a) the log-log plot of leakage current with fitted slopes α (lines); b) remanent polarization; c) dPr values derived from PUND measurements.

The current dependence on the applied field for all BFO films is Ohmic (α ∼ 1) at low field strengths (< 200 kVcm-1) (Figure 3 a)). Additional features can be noted in the case of 50 nm films which turn from Ohmic to a space charge limited conductance at field strengths E > 1MVcm-1. The 13 ACS Paragon Plus Environment


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200 nm films show a transition to interface restricted Fowler-Nordheim tunneling type of conductance at E > 300 kVcm-1.47 The relatively high Ohmic conductance limits the applicable field range for the 500 nm films to less than 200 kVcm-1. Figure 3 b) shows remanent polarization measured at a frequency of 100 Hz. The values of 0.06 and 0.30 µCcm-2 were obtained for 50 nm and 200 nm films. The values were found to decrease by an order of magnitude as the pulse frequency increased to 1 kHz. A significantly higher remanent polarization value of 20 µCcm-2 was measured for 500 nm films at 1 kHz pulse frequency. Figure 3 c) presents the PUND results with dPr given as a function of field for the 50 nm films at 1 Hz and for the 200 and 500 nm films at 1 kHz pulse frequency. The 500 and 200 nm thick films have 2dPr values above 100 and 60 µCcm-2 against 0.12 µC cm-2 noted for 50 nm films just before the current breakthrough. The increase of polarization with increasing thickness was most likely due to relative reduction of the depolarization field strength and strained rhombohedral like structure.16,48-53 Also, crystal grains smaller than 30 nm have been shown to exhibit significant polarization suppression due to surface effects. Selbach et al. suggested that the critical crystal size for ferroelectric polarization is around 9 nm.48 For photovoltaic devices, such as solar cells, depolarization field enables efficient charge separation and photovoltaic output.30 Multiferroic devices require modification of the Pt-BFO interface properties for smaller screening effects.16,50-52 In polycrystalline ferroelectrics with grain size below ∼100 nm single ferroelectric domain states are expected. However, polydomain states were also reported for crystal sizes down to 50 nm.50 It is worth noting that depolarization field can induce a transition from a mono-to polydomain state as the film thickness is decreased.52 With polycrystalline films of few tens of nanometer sized crystals domain pinning at grain boundaries is likely to determine the domain state. Therefore, in our 50 nm BFO films only single domain states were expected. In thicker films (200 and 500 nm) 14 ACS Paragon Plus Environment

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including grains up to 60 nm in size, both single and polydomain states may exist resulting in a wider domain size distribution. The domain configuration favored by the nano-crystalline morphology is likely to affect also the film magnetic properties playing a role in the destruction of the antiferromagnetic cycloid and formation of the magnetic domain walls.

Magnetic Properties Figure 4 shows magnetization vs. magnetic field measured in the 500 nm BFO films at different temperatures. The magnetization is seen to exhibit a hysteric behavior with both coercivity and saturation depending on the temperature. The magnetization behavior is essentially similar to a saturated ferromagnetic (FM) hysteresis measured in thinner BFO films (50 nm) grown by ALD and presented in detail elsewhere.31 The magnetic characteristics derived from the loops, viz., saturation magnetization (Ms), remanent magnetization (Mr), coercive field (Hc) and exchange bias (He) are summarized in Table 1. The hysteresis loops clearly indicate the existence of a ferromagnetic state in the films. The main contributions to the magnetization are expected to arise from the FM spin canting and ferromagnetic domain walls.4,13



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Figure 4. Magnetization as a function of applied field measured in 500 nm BFO films at different temperatures. T

Ms

Mr

Hc

He

[K]

[emucm-3]

[emucm-3]

[kOe]

[kOe]

(500/50nm) (500/50nm) (500/50nm) (500/50nm) 1.8

29.1/36.2

16.6/21.7

0.66/0.45

0.28/0.04

100

25.5/31.3

9.1/20.4

0.27/0.29

0.03/0.01

300

22.2/24.4

2.8/0.9

0.06/0.01

0.01/0.01

Table 1. Parameters characterizing the magnetic properties of ALD grown BFO nano-crystalline 50 and 500 nm thick films.31


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We assume that the cycloid-like magnetic modulation was broken or significantly distorted in the nano-crystalline grains of size less or comparable with the wavelength of the helimagnetic spin order (∼62 nm). The magnetic properties of the films dramatically differing from the bulk BFO antiferromagnetism are presumably associated with the nano-structured morphology of our films. There is a growing body of evidence for size effects leading to the enhancement of the net magnetic moment and ferromagnetism in various nano-structured BFO systems produced by different techniques.8-11,20 Essential contribution to magnetization can arise from the domain walls even in those multiferroics which are not intrinsically ferromagnetic.13,14,31 The nano-crystalline morphology is likely to modulate ferroelectric domain sizes and the corresponding domain boundary network leading to high density of ferroelectric 71°, 109° and 180° domains. The domain boundaries can be ferromagnetic, in particular, the 180° domain wall, even if the domain itself is antiferromagnetic.13 This implies a strong correlation between film morphology, ferroelectricity and magnetization which is discussed further later in terms of the magnetic domain wall pinning. Our results are consistent with data reported by other authors indicating a prominent relevance of the BFO film nano-crystalline morphology to their magnetic properties. The exchange bias detected in the films implies existence of either a ferromagnetic / antiferromagnetic interface, spin glass or ferromagnetic domain interactions.54

Magnetic Impurities Since magnetic properties of BFO-based materials are of great interest, we have examined carefully the origin of the high magnetization in our case to rule out possible magnetic impurities



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(especially, transition elements) and non-stoichiometric phases. The expected stoichiometry of BiFeO3 through the film thickness has been ensured by ERDA implying a single crystal phase to be present in the films. The XRD analysis and XPS suggest that within the measurement sensitivity no other crystal phases or impurity related Fe oxidation state is present. Besides, zero field cooled / field cooled (ZFC/FC) magnetizations (see Figure 5) also confirm the absence of parasitic ferromagnetic phases. Therefore, it is unlikely that the exchange bias detected in the films is due to interaction with ferromagnetic impurity phases.16,55 As mentioned above, two bismuth iron oxide parasitic phases could be formed during the BFO film synthesis, viz., Bi2Fe4O9 and Bi25FeO40. Han et al.55 have studied magnetism of these phases and have shown that the Néel temperature of the Fe rich phase is 263 K, whereas the Bi rich phase is paramagnetic above 5 K.57 Therefore, possible existence of the parasitic phases in our films cannot explain the peculiarities of the magnetization temperature dependence. Iron oxide impurity is another possible source of magnetism. Most of the Fe2O3 polytypes are antiferromagnetic and, therefore, cannot contribute to the magnetization. γ-Fe2O3 is known to be weakly ferromagnetic but nanoparticles show a blocking temperature of ∼70 K.56,58 Nano-inclusions of magnetite phase (Fe3O4) could, in principle, give rise to rather strong magnetism exhibiting also temperature and particle size dependence.57,59 However, our XPS measurements ruled out presence of significant amount of Fe2+ in our films (see Figure S2 in SI). Besides, it has been shown that Fe3O4 nanoparticles as small as 18 nm have a blocking temperature above the room temperature.59 The saturation magnetization of ∼ 410 emucm-3 exhibited by Fe3O4 is extremely high.59 Therefore, if magnetization of the present films is due to embedded magnetite nanoclusters then approximately 5 % of the total volume should be occupied by Fe3O4 and, hence, should have been


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detected by XRD and HR-TEM imaging. The same was found for the 50 nm BFO films for which XRD and HR-TEM has revealed no parasitic phases.

Uncompensated Spins Zero field cooled (ZFC) and field cooled (FC) magnetization of the ALD grown BFO films is presented in Figure 5 as a function of temperature. Magnetization is seen to increase with the applied magnetic field. The data for 20 kOe is plotted against the right hand scale for the sake of clarity.

Figure 5. Temperature dependence of magnetization (500 nm film) measured in different applied magnetic fields. The lower and upper branches of each paired curve show ZFC and FC 19 ACS Paragon Plus Environment


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magnetizations, respectively. The lower graph is a magnification of the ZFC curve obtained below 40 K at 0.05 and 0.1 kOe showing cusps. Vertical arrows in the upper graph show the irreversibility temperature Tf. In Figure 5, the ZFC magnetization curves measured in the fields of 50 and 100 Oe show obvious cusps at ∼15 K. We suggest that the cusps are due to the uncompensated spins which are assumed to be frozen at 1.8 K and which start to align with the field when the temperature is increased. Above ∼15 K the magnetization is seen to decrease which is most likely because of the thermal excitations gradually degrading the alignment of the weakly bound uncompensated spins. This is consistent with our AC susceptibility measurements shown in Figure 6.

Figure 6. Real part of the AC susceptibility measured in ALD grown BFO films. The measurements were performed in an external magnetic field of 1 Oe and frequencies of 1, 10 and 100 Hz. The insert shows the AC susceptibility for 10 Hz over the extended temperature range.


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Magnetic Viscosity A temporal aspect in magnetization measurements may shed light on the processes governing the magnetic behavior of the BFO films. The time dependence exhibits relatively slow magnetization dynamics in a constant magnetic field that is consistent with the existence of a frustrated state (see Figure S4 in SI). The magnetic viscosity can be explained by either a spin-glass behavior or relaxation of domain walls pinned to defects.59,60,61 In both cases, magnetization is expected to change with time following Equation 1:62 M (t ) = M 0 + S ln(t ),

(1)

where M0 is the initial magnetization before the magnetic field increase, S is the magnetic viscosity and t is time. In fitting a theoretical curve M vs. t calculated using Equation 1, the best fit parameters are M0 = 16.281 ±0.001 emucm-3 and S = (2.03 ±0.02)x10-2 emucm-3. The excellent agreement between the calculated M(t) and experiment, R2 = 99.6%, supports the existence of the frustrated magnetic state in our BFO films.

Magnetic Domain Wall Pinning As it follows from Figure 5, ZFC and FC curves split at an irreversibility temperature Tf that is a signature of a frustrated state existing in the films. Figure 7 a) shows how Tf was derived from the (MFC - MZFC) vs. T plots for 50 and 100 Oe fields. The estimated values for Tf were obtained by extrapolating the linear part of the curves.



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Figure 7. Difference between FC & ZFC magnetization for an ALD grown BFO film (500 nm) measured in external magnetic fields of 50 and 100 Oe is presented in a). The irreversibility temperature Tf vs. magnetic field is plotted in b) for 50 and 500 nm thick films.

The frustrated state, associated with the observed ZFC and FC magnetization curve splittings, in BFO films has been commonly attributed to either superparamagnetism,63 spin glass or domain wall depinning.64,65 In order to discriminate between the mechanisms, Figure 7 b) shows that Tf vs. magnetic field can be fitted by the following Equation 2:66

[

]

H = H 0 exp − bT f (H ) →

dT f dH

−1

= (bH )

(2)

where b is a fitting parameter and H0 is the strength of the applied field at zero temperature required for depinning of the magnetic domains from ZFC state to FC.66 The exponential dependence with the same value for b equal to 0.0108 K-1 is found for both film thicknesses which is consistent with the film magnetization controlled mainly by the evolution of domain walls pinned at defects,65 thus ruling out the significant contribution of other mechanisms. The differing pre-exponential factors H0, viz., 3.8 kOe and 10 kOe for 50 and 500 nm films implies a lower pinning energy for the thinner film, presumably, due to the lower surface anisotropy (energy) of smaller randomly oriented crystals and, hence, domains in the 50 nm film.67 22 ACS Paragon Plus Environment

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As it follows from Table 1, magnetization is, on the average, stronger in 50 nm films, but coercivity is lower if compared to the 500 nm films. We assume that this is due to the smaller average crystal grain size in thinner films and, hence, a stronger contribution of the surface/interface spins and suppressed spin cycloids. The exponential dependence H(Tf ) in Figure 7 clearly indicates that pinning of the magnetic domain walls played a central role in magnetization of our films at temperatures below ∼100 K.16,68 The magnetization by magnetic domain walls can be estimated by Equation 3:13,14

M DW =

M 0d

ω

.

(3)

Where MDW is the magnetization of the 1800 domain walls as described in reference (13), M0 = 130 emucm-3, the domain wall width d = 12 nm and the domain size ω. The polycrystalline 50 and 500 nm thick BFO films consist grains with average size of 30 and 40 nm.31 The crystal grains are expected to be in (FE) mono domain state in grains smaller than 50 nm. Therefore, grain boundaries constrain the smallest domain size to the size of the grain that in turn sets the upper limit for grain boundary density.69 Substitution of grain size by domain size in Equation 3 yields domain wall magnetizations of 44 and 33 emucm-3 for the 50 and 500 nm BFO films. The values are two times higher than the value at 2 K where the remanent magnetization is the highest. In the calculation, it was assumed that all domain walls are magnetic corresponding to the boundary between the 1800 -oriented FE domains. The randomly oriented polycrystalline BFO structure consists equally 710, 1090 and 1800 domain walls, in which the 710 walls are not considered to be magnetic,13 and magnetization of 1090 walls is between those of 710 and 1800. Assuming 1/3 density for the 1800 domain walls yields 14 and 11 emucm-3 for the 50 and 500 nm films. The measured magnetization is acquired assuming that the 1090 domain walls can have half of the



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magnetization strength of the 1800 domain walls yielding an additional magnetization of 7 - 5 emucm-3 for the 50 and 500 nm films. Alternatively, the magnetization by canted spins is reported to be around 6 - 8 emu cm-3.13 The spin canting is affected by the breakage of the spin cycloids due to grain size leaving only a weak remanent magnetization in spite of the increase in saturation magnetization.9,68 Therefore, we conclude that the remaining magnetization is mainly due to weakly ferromagnetic 1090 domain walls. In 500 nm films, Mr decreases as the temperature rises from 2 to 100 K because of depinning of the magnetic domain walls between the grains. In comparison, the narrow size distribution of crystal grains in the 50 nm film restricts the number of possible domain configurations resulting in steeper reduction in Mr as the threshold temperature is reached. At room temperature the domain pinning and crystallinity have only a weak contribution to the magnetic domain size.

Magnon Propagation The shape of the FC magnetization curves presented in Figure 5 provides some insight into the magnetic behavior of the ALD grown BFO films. For example, the FC curve measured at 10 kOe can be subdivided into two temperature regions. A high temperature region from ∼100 K to 400 K exhibits a slow decrease in the magnetization as the temperature increases. The temperature dependence within this range can be explained by the existence of spin waves (magnons) inside the material.12 This part of the magnetization curve can be fitted by the Bloch’s law ( Fitting parameters in Table S1 in SI):70

(

)

M Bloch (T ) = M 0 B 1 − (T TBloch ) . α

(4)


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Where M0B is the saturation magnetization as the temperature approaches zero, TBloch is the Bloch temperature at which the magnetization vanishes completely and α is the Bloch exponent. In the low temperature region (1.8-70 K), the magnetization is seen to decrease exponentially. Two reasons are conventionally associated with such a temperature dependence. First, this can be explained by low density of magnon states at low temperatures in low-dimensional nanostructured magnetic systems such as nano-sized grains, as it was suggested, for example, in case of BiFeO3-CuO nanocomposites.71 An another reason can be associated with magnetic behavior of the uncompensated surface spins in magnetic nanoparticles.60 Thermal excitation upon heating was suggested to break the weak ferromagnetic coupling of the uncompensated spins, thereby, leading to fast decrease of magnetization. Both mechanisms predict an exponential behavior of magnetization, so that it is difficult to discriminate between them. Nevertheless, the latter is consistent with the observed (ZFC) magnetization cusp in low field strengths (Figure 5). Therefore, the FC magnetization temperature dependence can be fitted (see Table S1 in SI) by using a modified version of the Bloch’s law including an additional exponential term as in Equation 5:72 T M (T ) = M Bloch (T ) + A exp( − T freezing ),

(5)

where A is a magnetization induced by uncompensated spins and Tfreezing a characteristic temperature. The agreement between the modified Bloch’s law and the experimental data is rather good, R2 = 99% (see Figure S5 in SI). For high magnetic fields (10 and 20 kOe), the α value ( 1.7 - 1.8 ) is a bit higher than the theoretically predicted value of 3/2 for ferromagnetic materials. This deviation has been already reported for some magnetic materials and attributed to reduced dimensions of nano-sized grains as in our nano-crystalline BFO films.73



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ALD technique was demonstrated to be capable of delivering a material with non-trivial and versatile electromagnetic properties that might be useful for microelectronic and photovoltaic applications. The polycrystalline BFO is considered to be viable material for solar cell mass production and ALD is a flexible method for production of different complex electrode structures that may be needed for efficient energy harvesting.30 Based on the active nature of ferromagnetoelectric domain walls of ALD grown films we believe that films can be modified for photoferroelecric –and magnetic applications as well.74 The ALD of BFO/Pt films can provide a highly tunable method for mesoporous BFO films that in turn are shown to exhibit the much desired strong ferroelectric-magnetic coupling.8 Formation of intrinsic strain in polycrystalline BFO gives an interesting possibility to modify and enhance the band gap properties for photovoltaic applications.29,30 The high degree of control over crystallization process, grain size and strain provided by an ALD process can be used for tailoring multiferroic properties. Being fully compatible with the mass-production paradigm, ALD offers unprecedented uniformity and conformality in device design of 3D and high aspect ratio structures adding extra functionality to the current planar technology in microfabrication.


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Conclusions Multiferroic BiFeO3 films with thickness from 50 to 500 nm were synthesized by ALD using successive deposition of alternating BiOx and FeOx layers. The polycrystalline films were found to exhibit non-trivial thickness and morphology dependent ferromagnetic and ferroelectric properties. The observed saturated ferromagnetic state exhibited coercivity even at high temperatures. Enhanced magnetization of the ALD grown films as compared to the bulk counterpart was shown to result from the formation of magnetic domain walls and breaking of the helimagnetic AFM order due to the phase confinement within the nanocrystalline morphology. ZFC-FC magnetization measurements revealed the presence of a frustrated state associated with complex dynamics of pinned magnetic domain walls. At low temperatures, uncompensated spins associated with nano-grain surfaces/interfaces were found to contribute to the magnetization and to the frustration of the magnetic system. Ferroelectric polarization found in thick films turned out to be strongly suppressed in thin films corresponding to an increase in the depolarization field and surface effects. It is found that strain is developed within crystal grains close to the bottom Pt electrode which can be highly useful for practical applications.

Supporting Information In the supporting information SEM imaging results and details of the XPS fittings for Fe 2p3/2, O 1s and C 1s valence states are presented. The Bloch fitting parameters are tabulated, figures presenting the Bloch fit and magnetic viscosity are included. This material is available free of charge via the internet at http://pubs.acs.org/.



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Acknowledgements Funding from Academy of Finland within the Finnish Centres of Excellence in Atomic Layer Deposition and in Low Temperature Quantum Phenomena and Devices is greatly acknowledged. The authors are grateful to the Electron Microscopy Unit of the Institute of Biotechnology, University of Helsinki for access to the HRTEM high resolution transmission electron microscopy.

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TOC figure

TOC Figure. The ALD synthesis of BFO films via alternating layer method.



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Figure 1. A) GI-XRD (2θ) pattern of a 500 nm thick ALD grown BiFeO3 film. B) high resolution image of the selected grain showing fringing patterns. 82x110mm (300 x 300 DPI)

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Figure 2. XPS analysis of a BiFeO3 film grown by ALD. The lines show fittings for the Fe 2p3/2 state and Shirley background A) and correspondingly, for the Bi 4f in B). 82x46mm (300 x 300 DPI)



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Figure 3. Electric properties of ALD grown BiFeO3 films: a) the log-log plot of leakage current with fitted slopes α (lines); b) remanent polarization; c) dPr values derived from PUND measurements. 80x164mm (300 x 300 DPI)


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Figure 4. Magnetization as a function of applied field measured in 500 nm BFO films at different temperatures. 82x104mm (300 x 300 DPI)



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Figure 5. Temperature dependence of magnetization (500 nm film) measured in different applied magnetic fields. The lower and upper branches of each paired curve show ZFC and FC magnetizations, respectively. The lower graph is a magnification of the ZFC curve obtained below 40 K at 0.05 and 0.1 kOe showing cusps. Vertical arrows in the upper graph show the irreversibility temperature Tf. 82x120mm (300 x 300 DPI)


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Figure 6. Real part of the AC susceptibility measured in ALD grown BFO films. The measurements were performed in an external magnetic field of 1 Oe and frequencies of 1, 10 and 100 Hz. The insert shows the AC susceptibility for 10 Hz over the extended temperature range. 82x62mm (300 x 300 DPI)



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Figure 7. Difference between FC & ZFC magnetization for an ALD grown BFO film (500 nm) measured in external magnetic fields of 50 and 100 Oe is presented in a). The irreversibility temperature Tf vs. magnetic field is plotted in b) for 50 and 500 nm thick films. 82x40mm (220 x 220 DPI)


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TOC Figure. The ALD synthesis of BFO films via alternating layer method. 84x27mm (300 x 300 DPI)


Electric and Magnetic Properties of ALD-Grown BiFeO3 Films - The

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