Rocksalt-type heavy rare earth monoxides TbO, DyO, and ErO exhibiting metallic electronic states and ferromagnetism

Abstract

Solid-phase rare earth monoxides have been recently synthesized via thin film epitaxy. However, it has been difficult to synthesize heavy rare earth monoxides owing to their severe chemical instability. In this study, rocksalt-type heavy rare earth monoxides REOs (RE = Tb, Dy, Er) were synthesized for the first time, as single-phase epitaxial thin films. The REOs were characterized by hard X-ray photoemission spectroscopy, X-ray absorption spectroscopy, resonant photoemission spectroscopy, and magnetization measurements. These REOs showed metallic electronic states with the almost localized 4f states, indicating the 4fⁿ5d¹ electronic configurations of Tb, Dy, and Er ions. X-ray magnetic circular dichroism measurements of TbO evidenced the magnetic ordering of the 4f spins below the Curie temperature (T_C). The T_C was evaluated to be 233 K for TbO, 142 K for DyO, and 88 K for ErO, where the T_C decreased with the 4f electron number, approximately proportional to the de Gennes factor.

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Introduction

Chemically stable binary rare earth oxides are usually rare earth sesquioxides RE₂O₃ with RE³⁺ ions (RE = Sc, Y, lanthanoids), which are highly insulating and dielectric.¹ In contrast, rare earth monoxides REOs are metastable phases.² Among them, only EuO and YbO have been frequently studied because of the chemical stability of the Eu²⁺ and Yb²⁺ ions owing to the half or fully filled 4f orbitals.¹ Since EuO is an archetypal ferromagnetic semiconductor with the Curie temperature (T_C) of 69 K,³ other REOs with partially filled 4f orbitals are expected to exhibit intriguing electronic and magnetic properties. In the 1980s, light REOs (LaO, CeO, PrO, NdO, and SmO) were synthesized as metallic bulk polycrystals by high-pressure synthesis,² which were considered to be contaminated with rare earth metals.⁴ Recently, single-phase light REOs (LaO, CeO, PrO, NdO, and SmO) were successfully synthesized via thin film epitaxy.^5–9 LaO is a superconductor,⁵ CeO is a paramagnetic metal,⁶ PrO and NdO are low-T_C (28 and 19 K, respectively) ferromagnetic metals,^7,8 and SmO is a paramagnetic metal.⁹ Similarly, heavy REOs (GdO, TbO, HoO, and LuO) were synthesized as epitaxial thin films.^10–13 GdO, TbO, and HoO showed much higher T_C than light REOs (276, 231, and 131 K, respectively),^10–12 except for paramagnetic LuO with fully filled 4f orbitals of Lu²⁺ ions.¹³ These heavy RE ions with partially filled 4f orbitals possess large total angular momenta due to the parallel coupling of spin and unquenched orbital angular momenta as was seen in heavy RE nitrides RENs,¹⁴ in addition to the much higher T_C of heavy REOs than those of RENs. Such a high T_C, strong magnetization, and large spin–orbital interaction would be attractive for magnetics and spintronics. However, a sizable amount of the RE₂O₃ phase contained in most of the heavy REO thin films obscured their intrinsic physical properties, e.g. poor electrical conduction in GdO and TbO despite the presence of 5d electron carriers.^10,11 Also, DyO and ErO are expected to be high T_C ferromagnets, but have not been synthesized yet except for the minor byproducts in RE₂O₃ synthesized under high vacuum.^15,16 Thus, it has been strongly desired to unveil the intrinsic physical properties of single-phase heavy REOs and to elucidate the origin of their high T_C.

In the previous studies on light REO thin films, thin film epitaxy was effective to stabilize metastable REOs on suitable lattice-matched single crystal substrates. However, heavier REOs possess smaller lattice constants due to the lanthanoid contraction, which enhanced the lattice mismatch between REO thin films and commercial substrates. Previously, a rocksalt-type SrO buffer layer was used to obtain a highly crystalline EuO (lattice constant: 5.16 Å) epitaxial thin film,^17–19 although the SrO (5.14 Å) layer cannot be applied universally to other REOs whose lattice constants are broadly distributed: from LaO (5.295 Å)⁵ to YbO (4.87 Å).²⁰ And recently, a rocksalt-type CaO buffer layer was used to improve the crystallinity of the GdO thin film.²¹ In this study, we developed a rocksalt-type (Ca,Sr)O solid-solution buffer layer, whose lattice constant is tunable by the composition, being suitable for thin film growth of various heavy REOs (Fig. 1a), where the lattice mismatches between the thin film and the buffer layer for TbO and DyO except for ErO were significantly reduced in comparison with those between the thin film and YAlO₃ substrate (Table 1). As a result, single-phase TbO, DyO, and ErO were synthesized for the first time as epitaxial thin films. In contrast to the preliminary result of nonpure-phase TbO that was reported to be semiconducting,¹¹ these compounds possessed metallic electronic states, being consistent with the 4fⁿ5d¹ electronic configuration. In addition, DyO and ErO were discovered to be ferromagnetic. Their T_C was systematically changed with respect to the f-electron number: T_C = 233 K (TbO), 142 K (DyO), and 88 K (ErO).

Fig. 1 (a) Schematic crystal structure of TbO or DyO thin films on the (Ca,Sr)O buffer layer (left) and ErO thin film on the CaO buffer layer (right) on the YAlO₃ substrate. (b) XRD θ–2θ patterns for the TbO, DyO, and ErO thin films on the buffer layers. (c) Reciprocal space map around the TbO 224 diffraction and the YAlO₃ 334 diffraction peaks for the TbO thin film.

Table 1 Lattice constants of REO thin films, the YAlO₃ substrate, and buffer layers with their lattice mismatches

REO	Substrate/buffer layer	Mismatch [%]
TbO	YAlO3(×√2)	–4.24
c = 5.03 Å	5.26 Å
a = 4.90 Å	(Ca0.5Sr0.5)O	–0.20
V = 120.8 Å^3	5.04 Å
DyO	YAlO3(×√2)	–4.82
c = 5.00 Å	5.26 Å
a = 4.88 Å	(Ca0.5Sr0.5)O	–0.79
V = 119.8 Å^3	5.04 Å
ErO	YAlO3(×√2)	–5.39
c = 4.97 Å	5.26 Å
a = 4.84 Å	CaO	4.41
V = 116.4 Å^3	4.76 Å

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Experimental

Thin film growth

TbO, DyO, and ErO epitaxial thin films were grown by pulsed laser deposition with a KrF excimer laser (λ = 248 nm). A (Ca_0.5Sr_0.5)O buffer layer was used for the TbO and DyO thin films and a CaO buffer layer was used for the ErO thin film. Prior to deposition, the YAlO₃ (110) single crystal substrate was preannealed at 1200 °C for 4 hours to have a step-and-terrace surface.⁵ The buffer layers were grown on the substrates at 600–700 °C under 5.0 × 10⁻⁶ Torr of oxygen, where the (Ca_0.5Sr_0.5)CO₃ polycrystalline target was sintered with spark plasma sintering (Fig. S1†) and the CaO (99.9%) polycrystalline target was used for the (Ca_0.5Sr_0.5)O and CaO buffer layers, respectively. Subsequently, the REO thin films were grown on the buffer layers at 150 °C under 5 × 10⁻⁸ Torr of oxygen, where Tb, Dy, and Er metal (99.9%) targets were used for the corresponding REO thin films. The detailed growth conditions are summarized in Table 2. To prevent the film oxidation, 2–5 nm-thick amorphous AlO_x capping layers were in situ deposited on the thin films at room temperature. During deposition, film growth was monitored by in situ reflection high-energy electron diffraction (RHEED).

Table 2 Deposition parameters of buffer layers, REO thin films, and the capping layer

Layer	Target	T sub [°C]	P O2 [Torr]	E laser [J cm^−2]	Frequency [Hz]	Thickness [nm]
T sub: substrate temperature; PO2: oxygen pressure during growth; Elaser: energy density of the laser spot.
(Ca0.5Sr0.5)O	(Ca0.5Sr0.5)CO3	700	5.0 × 10^−6	0.50	10	5
CaO	CaO	600	5.0 × 10^−6	0.50	20	10
TbO	Tb	150	5.0 × 10^−8	0.70	10	10
DyO	Dy	150	5.0 × 10^−8	0.60	10	10
ErO	Er	150	5.0 × 10^−8	0.60	10	10
AlOx	Al2O3	R.T.	—	0.70	10	5

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Structure analysis

The crystal structure was evaluated by X-ray diffraction with Cu Kα₁ radiation (XRD, D8 Discover, Bruker AXS and SmartLab, Rigaku).

Synchrotron X-ray spectroscopy

For the electronic states, hard X-ray photoemission spectroscopy (HAXPES) was performed by using a Scienta R-4000 electron energy analyzer with a total energy resolution of 200 meV at a photon energy of 8 keV at BL09XU of SPring-8. Also, X-ray absorption spectroscopy (XAS) and RE 3d–4f resonant photoemission spectroscopy (RPES) were performed with total electron yield mode at BL-2A MUSASHI of Photon Factory, KEK. RPES spectra were recorded at various photon energies hν around the Tb, Dy, and Er 3d–4f thresholds for resonant photoemission measurement, where hνs of on- and off-resonances were determined by conducting XAS at the Tb, Dy, and Er M₅ absorption edges. The Fermi level (E_F) was calibrated by the measurement of a gold film that was electrically connected to the REO thin films. X-ray magnetic circular dichroism (XMCD) spectroscopy was performed with total electron yield mode to measure Tb M₄ and M₅ absorption edges at BL-16A of the Photon Factory, KEK.

Magnetization measurements

Magnetization was measured using a superconducting quantum interference device magnetometer (MPMS-3, Quantum Design). The thin films were mounted on straw holders. The diamagnetic signals from the holder and substrate were corrected by subtracting the magnetic-field-linear component of the magnetization fitted at a high magnetic field.

Results and discussion

Fig. 1b shows the XRD θ–2θ patterns of the TbO, DyO, and ErO thin films on YAlO₃ substrates with the buffer layers (Table 1). The REO 002 and 004 diffraction peaks overlapped with the (Ca,Sr)O 00l or CaO 00l diffraction peaks (Fig. S2†) were observed without any impurity phases. Out-of-plane lattice constants of the thin films were c = 5.03 Å (TbO), 5.00 Å (DyO), and 4.97 Å (ErO). Since the lattice mismatch of the ErO thin film was not significantly reduced by using the buffer layer (Table 1), not only the small lattice mismatch but also the rocksalt-type buffer layer was appropriate for the epitaxial stabilization of the REO thin films. The spot 224 peak of the TbO thin film in reciprocal space mapping (Fig. 1c) confirmed the epitaxial relationship of TbO [001]∥YAlO₃ [110] and TbO [110]∥YAlO₃ [001]. Epitaxial growth of the DyO (001) and ErO (001) thin films was also confirmed by reciprocal space mapping, in situ RHHED patterns, and φ scan measurements (Fig. S3–S5,† respectively) with the same epitaxial relationship. The lattice mismatch between the thin film and the buffer layer was larger in the order of TbO, DyO, and ErO thin films, resulting in less crystallinity, as seen in the broader XRD peak intensity, the broader RHEED streak patterns, and asymmetric peaks of the phi scan for DyO and ErO thin films (Fig. 1 and Fig. S3–S5†). The obtained lattice constants are summarized in Table 1. The ionic radii of the 6-coordinated RE²⁺ ions were 1.07 Å (Tb²⁺), 1.06 Å (Dy²⁺), and 1.04 Å (Er²⁺), evaluated from the cube root of the cell volumes. These values were approximately the same as ionic radii of the RE²⁺ ions calculated from Shannon's ionic radii of 6-coordinated RE²⁺ ions: 1.10 Å (Tb²⁺), 1.09 Å (Dy²⁺), and 1.06 Å (Er²⁺).²²

Fig. 2 shows the HAXPES spectra of the TbO, DyO, and ErO thin films at room temperature. The RE 3d core level peaks shifted monotonically with the atomic numbers (Fig. 2a). The Tb and Dy peaks were located between those of RE metal and RE₂O₃ (Table S1†),^23–25 indicating the ionic character of TbO and DyO (the formal ionic charge of Tb²⁺ and Dy²⁺ ions in TbO and DyO, respectively), while HAXPES 3d core level spectra of Er and Er₂O₃ have not been reported. The valence band spectra of REO thin films showed the RE 5d states distributed within 4 eV below E_F and the existence of Fermi edges (Fig. 2b), similar to CeO, PrO, and SmO epitaxial thin films,^6,7,27 while the O 2p state and RE 4f multiplet²⁸ of the REOs for 6–10 eV were probably overlapped with the O 2p states and Al 3s states due to the comparable photoionization cross-section of RE 4f and Al 3s.²⁹ This result indicates that the TbO, DyO, and ErO epitaxial thin films possessed metallic electronic states, where the 5d electrons formed the conduction band near E_F.

Fig. 2 HAXPES spectra of the TbO, DyO, and ErO thin films. (a) The RE 3d core level spectra and (b) the valence band spectra. The observed spectra between 6–10 eV are attributed to O 2p and RE 4f multiplet of the REOs, probably overlapped with that of the amorphous AlO_x capping layer, whose band gap was reported to be ∼3.64 eV.²⁶

Fig. 3a–c show XAS spectra for the TbO, DyO, and ErO thin films obtained at RE M₄ and M₅ edges together with the theoretical spectra calculated for the RE³⁺ states.³⁰ The good agreement of XAS spectra between the REO, the theoretical calculations, and RE metals³⁰ indicates that an outer electron configuration of RE ions was 4fⁿ5d¹6s² [the 4f electron number of RE ions in REO was 4fⁿ (n = 8, 9, and 11 for Tb²⁺, Dy²⁺, and Er²⁺, respectively)]. From the HAXPES results in Fig. 2, the presence of the 5d¹-electron-derived conduction band at E_F was confirmed. Accordingly, the electronic configuration of the RE ion side in REO is 4f⁸5d¹ (TbO), 4f⁹5d¹ (DyO), and 4f¹¹5d¹ (ErO), while the 6s² electrons are accommodated in the O 2p–RE 6s bonding states. Fig. 3d–i show the RE 3d–4f RPES spectra obtained at the corresponding photon energies for on- and off-resonances at the M₅ absorption edges. Owing to the giant resonance enhancement for the 4f states, the 4f derived states were extracted irrespective of the presence of the AlO_x capping layer. At E_F, the intensity for the on-resonance states was insignificantly enhanced in comparison with those of CeO and PrO,^6,7 representing a small 4f contribution to the states near E_F. Meanwhile, the 4f-derived multiplet structures at around 8 eV for TbO, at 8 eV for DyO, and at 12 eV for ErO were strongly enhanced for on-resonant spectra. These results indicate that the majority of 4f electrons are localized so that the hybridization between conduction and f electrons is weak in TbO, DyO, and ErO in contrast to CeO and PrO.^6,7

Fig. 3 XAS spectra for (a) TbO, (b) DyO, and (c) ErO thin films around the RE M₄ and M₅ edges (solid line) together with the calculated spectra (dashed line) for the RE³⁺ states. Valence band RPES spectra for (d and g) TbO, (e and h) DyO, and (f and i) ErO thin films with (g–i) the magnified data around E_F.

Fig. 4a–c show the temperature dependence of magnetization for the REO thin films under field-cooling (FC) and zero-field-cooling (ZFC). Since the onsets of magnetization were not clearly observed for DyO and ErO (Fig. 4b and c), the T_C values of DyO and ErO were evaluated from the onsets of dM/dT–T curves, whereas the T_C of TbO was evaluated from the negative peak of the dM/dT–T curve (Fig. S6†) like that of EuO,³¹ where the evaluated T_C was consistent with M⁻¹–T curves (Fig. S6†). The evaluated T_C was 233 K (TbO), 142 K (DyO), and 88 K (ErO). It is noted that the T_C was lower for the heavier REOs similar to rocksalt-type rare earth nitrides RENs,³² in spite of their significantly lower T_C: T_C = 42 K (TbN), 26 K (DyN), and 5 K (ErN). Fig. 4d–f show the magnetic field dependence of magnetization for the REO thin films at different temperatures. The magnetic hysteresis loops were observed up to 200 K, 100 K, and 80 K for TbO, DyO, and ErO, respectively, being consistent with the T_C described above and the temperature dependence of the coercive force (Fig. S6†). The saturation magnetization at 2 K and 7 T was 5.1μ_B/f.u. (TbO), 3.2μ_B/f.u. (DyO), and 5.6μ_B/f.u. (ErO). These values were smaller than the theoretical magnetic moments of the trivalent rare earth ions [9μ_B (Tb³⁺), 10μ_B (Dy³⁺), and 9μ_B (Er³⁺)]. The smaller magnetic moments have often been observed for rare earth nitrides,^33,34 possibly attributed to crystal field effects,^32,35 the noncollinear spin structure,³⁶ and/or antiferromagnetic superexchange coupling between RE–O–RE bonds. The smaller magnetization of DyO among them could be caused by the larger lattice mismatch (Table 1) as well as large anisotropy of the 4f orbital.³⁷ In Fig. 4d, wasp-waisted hysteresis was not observed in TbO in contrast to the previous study,¹¹ probably because the Tb₂O₃ impurity phase in the previous study served as a pinning center of the magnetic domain wall motion resulting in wasp-waisted hysteresis.³⁸

Fig. 4 (a–c) Temperature dependence of magnetization for (a) TbO, (b) DyO, and (c) ErO thin films under field-cooling (FC) and zero-field cooling (ZFC). A magnetic field of 0.1 T was applied along in-plane. (d–f) Magnetic field dependence of magnetization for (d) TbO, (e) DyO, and (f) ErO thin films at different temperatures. A magnetic field was applied along in-plane.

Fig. 5 shows XAS and XMCD spectra around Tb M₅ and M₄ absorption edges of the TbO thin film at 9 K and 1.25 T. The XAS spectrum of the TbO thin film was consistent with that in Fig. 3a. The TbO thin film clearly showed nonzero intensity in the XMCD spectrum, whose multiplet feature corresponded to that of the XAS spectrum. The multiplet feature, despite the different spectral shape, was similar to that of the Tb₃Fe₅O₁₂ garnet³⁹ rather than that of Tb based alloys.⁴⁰ These results reflect the ferromagnetic states of the RE³⁺ ions in REOs. The XMCD M₅ peak monotonically decreased with increasing temperature up to 100 K (inset of Fig. 5). Accordingly, the 4f and 5d global ferromagnetic order was formed below the T_C contrary to the previous study's suggestion that the 4f and 5d ferromagnetic sublattices form below 20 K and T_C, respectively.¹¹

Fig. 5 XAS and the corresponding XMCD spectra at 9 K of the TbO thin film. The inset shows the temperature dependence of the Tb M₅ peak in the XMCD spectrum. A magnetic field of 1.25 T was applied along out-of-plane.

The much higher T_C of TbO, DyO, and ErO than those of TbN (42 K), DyN (26 K), and ErN (5 K) would be attributed to the principal role of the 5d conducting carriers for the Ruderman–Kittel–Kasuya–Yosida (RKKY) interaction, taking into account the 4fⁿ5d¹ and 4fⁿ5d⁰ electronic configurations of the REOs and RENs, respectively. This is why the T_C of RENs is enhanced by increasing the carrier concentration through N deficiency.^41,42 The T_C of REOs decreased with increasing 4f electron number similar to heavy rare earth metals.⁴³ The T_C of heavy rare earth alloys is a monotonically increasing function of the de-Gennes factor ξ,⁴⁴ well scaled by ξ^2/3.⁴⁵ In contrast, the T_C of heavy REOs was rather proportional to ξ (Fig. 6), although this origin is unclear at present. It is also noted that the oxygen content in REOs could influence their electronic states and magnetism, since GdO was found to show significant dependence of electric and magnetic properties on the oxygen content.^10,21 Further investigation on this matter is necessary, although the precise control of the oxygen content in REOs is also challenging.

Fig. 6 Curie temperatures of heavy rare earth monoxides as a function of the de Gennes factor ξ = (g − 1)²J(J + 1), where g is Lande's g-factor and J is the total angular momentum quantum number. The T_C values of HoO and GdO are cited from previous studies.^10,12,21

Conclusion

Rocksalt-type heavy rare earth monoxides TbO, DyO, and ErO were synthesized as single phase epitaxial thin films by utilizing epitaxial force from the lattice-matching-tunable buffer layer. These rare earth ions possessed [Xe]4fⁿ5d¹ electronic configurations, and 4f and 5d electrons served as localized and conduction electrons, respectively, indicating the metallic electronic states even above the Curie temperature in contrast to the semiconducting EuO.³ Much higher Curie temperatures of TbO, DyO, and ErO than those of the corresponding heavy rare earth mononitrides suggest a significant influence of RKKY interactions owing to the 5d conduction electrons.

Author contributions

S. Sasaki: investigation, methodology, data curation, funding acquisition, writing – original draft, and writing – review & editing. D. Oka: investigation, data curation, funding acquisition, writing – original draft, and writing – review & editing. D. Shiga: investigation, data curation, funding acquisition, writing – original draft, and writing – review & editing. R. Takahashi: investigation, data curation, and writing – review & editing. S. Nakata: investigation, data curation, writing – original draft, and writing – review & editing. K. Harata: investigation, data curation, and writing – review & editing. Y. Yamasaki: resources, data curation, funding acquisition, and writing – review & editing. M. Kitamura: resources, data curation, and writing – review & editing. H. Nakao: resources, data curation, and writing – review & editing. H. Wadati: resources, funding acquisition, data curation, writing – original draft, and writing – review & editing. H. Kumigashira: resources, funding acquisition, data curation, writing – original draft, and writing – review & editing. T. Fukumura: resources, supervision, data curation, funding acquisition, writing – original draft, and writing – review & editing. All authors read and approved the final manuscript.

Data availability

Data for this article are included in the manuscript and the ESI† or are available upon request.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

The authors thank the Cooperative Research and Development Center for Advanced Materials, Tohoku University for technical assistance, and Dr Kenta Amemiya for technical support of XMCD measurements. The HAXPES experiment at SPring-8 was conducted with approval from the Japan Synchrotron Radiation Research Institute (Proposal No. 2022B1574). The work performed at KEK-PF was approved by the Program Advisory Committee (Proposals No. 2022G675, 2021S2-002, and 2021S2-004) at the Institute of Materials Structure Science, KEK. S. S. was supported by a Grant-in-Aid of Tohoku University, Division for Interdisciplinary Advanced Research and Education. This work was in part supported by JSPS-KAKENHI (Grant No. 20H02704, 21H05008) and JST SPRING (Grant No. JPMJSP2114).

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d4dt03214d

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