Growth of PdCoO₂ films with controlled termination by molecular-beam epitaxy and determination of their electronic structure by angle-resolved photoemission spectroscopy

Metallic oxides with the delafossite structure, shown in Fig. 1(a), have drawn significant attention in recent years due to their unique structural and electronic properties. Examples include PtCoO₂, which has the highest conductivity per carrier of all materials, and PdCoO₂, which has the longest mean free path (exceeding 20 μm at 4 K) among all known metals.^1–3 The in-plane electrical conductivity of PdCoO₂ at room temperature, which is about four times higher than that of palladium metal itself, has been argued to arise from electron–phonon scattering mainly occurring within a single, closed, highly dispersive band of primarily palladium character at the Fermi level (E_F).^1,4–8 The large spin-splitting of the surface state arising from the CoO₂ termination, in combination with the layered structure of PdCoO₂-based heterostructures makes this system ideal to study itinerant surface electrons driven by inversion-symmetry breaking.⁹ As for the magnetic properties of delafossites, PdCrO₂ is the only highly conducting delafossite material that orders magnetically; it orders antiferromagnetically (AFM) at around 37 K. Focusing on the electronic structure, the single band at the Fermi level with palladium character forms a reconstruction driven by the AFM order from the adjacent CrO₂ layer.^10–14 Comparing AFM PdCrO₂ with nonmagnetic PdCoO₂, the spins from Cr³⁺ interacting inside the CrO₂ layer with the palladium monolayers on either side of the CrO₂ layer play a critical role in the magnetic state of PdCrO₂.¹³ 

Angle-resolved photoemission spectroscopy (ARPES) is the premiere experimental tool to directly observe electronic structure in quantum materials. The combination of oxide molecular-beam epitaxy (MBE) with ARPES allows us to customize and study the electronic structure of correlated oxides. This setup has enabled understanding of how strain,^15,16 thickness,^17–19 hetertostructuring,^20,21 interfaces,^22,23 and terminations^24–29 can influence the electronic structure of thin films.

Due to the limited size of delafossite single crystals, the desire to explore the potential of metallic delafossites in electronic³⁰ and spintronic devices,³¹ together with the exciting opportunities that can be explored in delafossite heterostructures, metallic delafossites are being grown in thin-film form by reactive sputtering,³² pulsed-laser deposition (PLD),^30,33–36 and MBE.^37,38 Unfortunately, the transport properties of these films (so far) pale in comparison to the single crystals. Differences between the Fermi surface of the so-called Pd-terminated PdCoO₂ with CoO₂-terminated PdCoO₂ have recently been reported for epitaxial films of PdCoO₂ grown by PLD.³⁹ The claims of this prior study would be strengthened by improved evidence of surface termination control.

In this work, we describe an improved synthetic strategy for the growth of PdCoO₂ films with control of surface termination by MBE. Harnessing the ultra-high vacuum connection between our MBE and ARPES, we then study the electronic structure of Pd-terminated and CoO₂-terminated PdCoO₂. We find that our PdCoO₂ films exhibit similar bulk bands derived from palladium states but weak surface states compared to those in PdCoO₂ single crystals. Having succeeded in engineering the surface termination in PdCoO₂ over large areas, we then progress to investigate terminations of PdCoO₂ where we deliberately add unpaired electrons and study the resulting electronic structure by ARPES. Although we are able to epitaxially stabilize a variety of surface terminations involving iron substituting for cobalt in PdCoO₂, we do not see evidence of magnetic order.

Building upon our previous work,³⁸ thin films of PdCoO₂ were synthesized by MBE in a Veeco GEN10 MBE system on (001) sapphire substrates. Details of the film growth are provided in the supplementary material. Figure 1(a) shows the shutter timing diagram used to supply fluxes of the individual molecular beams to the substrate to form PdCoO₂. After growth, films were cooled down to 300 °C in the same ozone background pressure (around 5 × 10⁻⁶ Torr) in which they were grown and transferred under ultra-high vacuum conditions into an adjacent ARPES chamber. The reflection high-energy electron diffraction (RHEED) patterns acquired after deposition and the x-ray diffraction θ–2θ scans indicate the growth of single-phase PdCoO₂ films as shown in Figs. 1(b) and 1(c). The structure was characterized by a Panalytical Empyrean x-ray diffractometer utilizing Cu Kα₁ radiation. Electrical transport measured by a Quantum Design Physical Property Measurement System (PPMS) using a four-point van der Pauw geometry is shown in Fig. 1(d). The residual resistivity ratio (RRR = ρ_300K/ρ_4K) of this PdCoO₂ sample with a thickness of 52.1 nm is 3.3 in its as-grown state (i.e., no ex-situ postanneal). For comparison, the highest RRR previously achieved by MBE for a film in its as-grown state was 2.2 for a 10 nm thick PdCoO₂ film.³⁸ After an ex-situ anneal, films grown by MBE can reach a RRR of around 8 for a 50 nm thick PdCoO₂ film.³⁷ While these are the highest reported RRR values for films, PdCoO₂ single crystals can exhibit RRR as high as 400.¹ The resistivity of epitaxial PdCoO₂ films and single crystals are comparable at room temperature; the huge difference in resistivity emerges upon cooling. One likely defect responsible for this difference is the in-plane rotation twins present in all epitaxial delafossite films to date. The presence of 180° in-plane rotation twins in PdCoO₂ film grown on (001) Al₂O₃ substrates manifest in the x-ray ϕ scan and scanning transmission electron microscopy images shown in Ref. 38, as well as the atomic force microscopy images shown in the supplementary material (Fig. S3).

In-situ ARPES measurements are performed to study the effects of termination on the electronic structure of the PdCoO₂ films. Our lab-based ARPES system photoexcites electrons with a Scienta omicron VUV 5000 helium discharge lamp using He–I photons at 21.2 eV and He-II photons at 40.4 eV. The emitted electrons are detected with a VG Scienta R4000 electron analyzer. The ARPES is vacuum-connected to the MBE growth chamber via an ultra-high vacuum transfer chamber.

We first compare the Fermi surface of Pd-terminated and CoO₂-terminated PdCoO₂ films in Figs. 2(a) and 2(b), respectively. The sharp hexagonal pocket centered at Γ [illustrated in red in Figs. 2(a) and 2(b)] is observed in both Pd- and CoO₂-terminated PdCoO₂ films. Two smaller hexagonal pockets inside the bulk state pocket illustrated in green and blue are observed in the CoO₂-terminated PdCoO₂ film, in agreement with previous reports of splitting of the CoO₂ surface state driven by spin–orbit coupling.^40,41 For Pd-terminated PdCoO₂, we do not observe pronounced palladium surface states at E_F. Below E_F, however, there is some spectral weight possibly from the palladium surface state as described in Ref. 41. At He-II photon energy (40.4 eV), we observe stronger spectral weight below E_F, as illustrated in Fig. S5 of the supplementary material. This intensity below E_F might be related to the palladium surface state, but at this higher energy, still no palladium surface state appears at E_F.

In Figs. 3(a) and 3(c), we further compare the dispersion cut along the Γ-K direction of Pd- and CoO₂-terminated PdCoO₂, respectively. The fitted Fermi velocities (v_Fs) of the PdCoO₂ bulk state of Pd- and CoO₂-terminated films are all around 4.5 eV Å, as shown in Table I, in agreement with previous results measured on PdCoO₂ single crystals by ARPES.^9,13 Despite the invisible palladium surface state at E_F, two spin-split surface states from the CoO₂ termination show up at E_F. These are indicated by blue and green circles in Fig. 3(c). From the momentum dispersion curves (MDCs) comparison of Pd- and CoO₂-terminated samples in Fig. 3(e), it is easier to observe the CoO₂ surface states at E_F (indicated by the blue arrows) and very little of the palladium surface state is observed below E_F (indicated by the red arrow). The Fermi velocities of the CoO₂ surface states are 0.75 eV Å (blue) and 0.5 eV Å (green), respectively. This is roughly 10% of that of the bulk state, in agreement with the previous study of PdCoO₂ single crystals.^9,42 Dispersion along the K-M-K direction of Pd- and CoO₂-terminated PdCoO₂ films are shown in Figs. 3(f) and 3(h). Both terminations of the PdCoO₂ films show a split band at the M point 0.75 eV below E_F and 1.75 eV below E_F, as observed in PdCoO₂ single crystals.⁴¹ Interestingly, a hole band below E_F at the M point driven by the palladium surface state (as shown in Ref. 41) is not seen in our Pd-terminated films.

Comparing the electronic structure observed for our epitaxial PdCoO₂ thin films to that reported for PdCoO₂ single crystals with different terminations, our PdCoO₂ films have similar bulk state features to those of PdCoO₂ single crystals, but the surface states of our PdCoO₂ films are weaker or even disappear at E_F. Note that the alternating layers of Pd⁺ and $CoO2−$ along the c-axis of PdCoO₂ are not charge neutral. Doping of the surface by electrons arising from electronic reconstruction (i.e., no structural surface reconstruction) would generate the surface states observed on bulk single crystals. Ways in which the surfaces of our films differ from the single crystalsprovide routes to different or no surface states. For ARPES measurements of cleaved (001)-oriented PdCoO₂ single crystals, the polar surface exposed after cleaving may drive an electronic reconstruction of the surface or alternatively a mixture of termination regions, some of which are terminated by palladium and some of which are terminated by CoO₂, to alleviate the polar surface charge. To synthesize PdCoO₂ films, we use shuttered MBE growth to provide a full layer of palladium or CoO₂ as the terminating surface. This difference in the surface reconstruction structure of epitaxial PdCoO₂ films might result in a different electronic reconstruction from that of cleaved PdCoO₂ single crystals. We show the reconstruction of our films in low-energy electron diffraction (LEED) with different terminations in Fig. S4 of the supplementary material. Additional differences exist between our PdCoO₂ films and PdCoO₂ single crystals. For example, PdCoO₂ films grown on (001) Al₂O₃ substrates are known to have a high density of 180° in-plane rotation twins.³⁸ Furthermore, our films might contain oxygen vacancies to neutralize the surface polar effect. These may also play a role in the differences observed in the surface states between PdCoO₂ single crystals and our epitaxial films. In particular, the palladium surface state is very reactive; it can be essentially destroyed by temperature cycling of ARPES measurements.⁴² Prior ARPES results of a PdCoO₂ film grown by PLD,³⁹ where the palladium termination is confirmed by the absence of a CoO₂ surface state in the electronic structure, does not show the PdCoO₂ bulk state. In contrast, our Pd-terminated films show a strong PdCoO₂ bulk state without a CoO₂ surface state, but are missing the palladium surface state. One possible reason for this is the difference in sample quality, particularly of the sample surface and the ability of MBE to control the termination of the PdCoO₂ film. We note that our MBE and ARPES measurements are made immediately following film growth and that the ultra-high vacuum connection between our MBE and ARPES systems enables us to investigate the electronic structure of the pristine growth surface.

With the goal of introducing magnetic order into the surface of PdCoO₂, we attempt to replace the cobalt in the CoO₂ surface termination with a different transition metal. The low spin state (S = 0) of the d⁶ electron configuration of Co³⁺ in octahedral coordination underlies the lack of magnetic order in PdCoO₂. In contrast, PdCrO₂ is known to order antiferromagnetically at 37 K due to the unpaired spins (S = 3/2) arising from Cr³⁺ with its d³ electron configuration.^10,11 Considering what other transition metals are stable in the 3+ oxidation state under the highly oxidizing conditions needed to stabilize PdCoO₂ led us to attempt to substitute Fe³⁺ for Co³⁺. Other known iron-containing delafossites, such as the semiconductors AgFeO₂ and CuFeO₂ with bandgap larger than 1 eV, are known to exhibit magnetic phase transitions.^43–46 Combining magnetic FeO₂ with conducting palladium might create a new metallic delafossite with interesting magnetic properties. Note that PdFeO₂ is neither a known compound nor has it been suggested to form in the delafossite structure by prior crystal chemistry based suggestions of delafossites^47,48 nor first-principles suggestions for new delafossites.^49,50

Our attempts to terminate PdCoO₂ with a monolayer of FeO₂ were successful. To do so, we used epitaxial stabilization, a process in which lattice misfit strain energies and interfacial energies are exploited to favor a desired metastable phase over the equilibrium phase.^51–56 ARPES measurements reveal the FeO₂-terminated PdCoO₂ film to have a similar bulk band to PdCoO₂, but there is no extra surface state at the Fermi surface nor a reconstruction driven by AFM order like in PdCrO₂. In addition, we synthesized Pd(Co, Fe)O₂ films containing a solid solution of iron and cobalt in each CoO₂ layer of the Pd(Co,Fe)O₂ film. Electrical transport measurements on a series of 13 nm thick PdFe_xCo_1−xO₂ films with varying x (0 $<x≤$ 0.2) all show a drop in resistance at low temperature. Other than the dip shown in dρ/dT at low temperature, which is different from other known delafossites, no pronounced magnetic order is observed by magnetic susceptibility measurement down to 3 K. This behavior is in contrast to what is seen for PdCrO₂. Turning to the electronic structure of the PdFe_xCo_1−xO₂ films revealed by ARPES, a similar bulk band to PdCoO₂ is observed. No reconstruction appears at the Fermi surface nor is any temperature dependence of the electronic structure of PdFe_xCo_1−xO₂ seen.

We added a full monolayer of FeO₂ to a Pd-terminated PdCoO₂ film in an attempt to vary the termination of PdCoO₂ by introducing unpaired electrons (spin) from Fe³⁺. As ARPES is a surface sensitive measurement, if the unpaired electrons from Fe³⁺ with its d⁵ configuration in the surface FeO₂ layer interact within the FeO₂ layer and with the adjacent palladium layer like the CrO₂ layer does in the AFM metal PdCrO₂,¹³ we expect to see distinct features arise in the Fermi surface of the FeO₂-terminated PdCoO₂ film. The well crystallized FeO₂ termination determined by RHEED is shown in Fig. 5(a). Unfortunately, no difference is observed in the ARPES other than the similar PdCoO₂ bulk state appearing at the Fermi surface in the FeO₂-terminated PdCoO₂ film, as shown in Figs. 2(c) and 2(g). No reconstruction of the Fermi surface was seen at low temperature. The bulk band was also free of any temperature-dependent feature when we analyzed the MDCs in the Γ-K direction shown in Fig. 4(e). Thus, our epitaxial FeO₂ layer on the surface of PdCoO₂ film does not appear to create a spin interaction with the underlying palladium layer. On the other hand, PdFeO₂ is a metastable phase and we can only stabilize one formula unit of PdFeO₂. It is possible that the FeO₂ termination is insulating due to oxygen vacancies to neutralize its otherwise polar surface and preventing it from contributing to the electronic structure. The photoemission intensity data we collect include the film beneath the FeO₂ layer, which is PdCoO₂ itself.

In addition to replacing the entire CoO₂ monolayer with an FeO₂ monolayer, we also investigated the partial replacement of cobalt with iron hoping that the presence of iron in multiple layers of the Pd(Co,Fe)O₂ structure would enhance the chance of spin interaction between the Fe³⁺ and the adjacent palladium layer. In Figs. 2(d) and 2(h), one can observe that the band of the PdFe_0.17Co_0.83O₂ film still has similar features to the bulk state of PdCoO₂, but with significant noise. Meanwhile, similar to the FeO₂-terminated PdCoO₂ film, no temperature dependence of the bulk band of Pd(Co,Fe)O₂ is seen at E_F [see Figs. 4(c)–4(f)]. We compared the Fermi velocity (v_F) of the PdFe_0.17Co_0.83O₂ film with the other three different terminations of PdCoO₂ films in Table I, and within experimental error, they all have the same v_F value as that of PdCoO₂ single crystals.⁴¹ Thus, the partial replacement of cobalt by iron in each CoO₂ layer does not bring any ordered spin interaction between the (Co,Fe)O₂ layers nor do we observe any evidence of interaction with the electrons of Pd(Co,Fe)O₂, which contribute to the bulk band of PdCoO₂. The noise observed in the Fermi surface of the PdFe_0.17Co_0.83O₂ film might come from the increased disorder accompanying the replacement of cobalt by iron. With the d⁵ configuration of Fe³⁺ and the d⁶ configuration of Co³⁺, the Fe_0.17Co_0.83O₂ layer should, in principle, have a d^5.83 configuration resulting in unfilled d electron bands. Nonetheless, no new conducting band is seen at E_F as shown in Fig. 4. On the other hand, a hole doping scenario in which 0.17 electrons move from the palladium layer to the Fe_0.17Co_0.83O₂ layer would leave the Fe_0.17Co_0.83O₂ layer in an insulating state. Such a scenario would require hole doping of 0.17 holes on the palladium layer. A comparison of the momentum at E_F (k_F) shown in Fig. S7 of the supplementary material, reveals no pronounced difference between the k_Fs of a PdCoO₂ film and a Fe_0.17Co_0.83O₂ film.

Further characterization of the PdFe_xCo_1−xO₂ films is shown in Fig. 5. The maximum percentage of iron that we are able to incorporate into epitaxial PdFe_xCo_1−xO₂ films while retaining a single phase is x = 20%. RHEED of a single-phase, 13 nm thick PdFe_0.2Co_0.8O₂ film is shown in Fig. 5(a). The fringes in the x-ray diffraction θ–2θ scans of the PdFe_xCo_1−xO₂ films indicate the high structural quality of these films. Electrical transport measurement on the PdFe_xCo_1−xO₂ films is shown in Fig. 5(d). Note that the PdCoO₂ film shown in this comparison has a thickness of 13 nm, a quarter of the thickness of the PdCoO₂ film in Fig. 1(d). The upturn in resistivity of the pure PdCoO₂ film (0% Fe) seen in Fig. 5(d) below 20 K likely originates from localization in the ultrathin film. As more cobalt is replaced by iron, the absolute resistivity of the iron-doped PdCoO₂ film keeps increasing. Interestingly, instead of showing an upturn at low temperature like is seen in the pure PdCoO₂ film, the Fe-doped PdCoO₂ films show a drop at low temperature in electrical resistivity. Moreover, as the iron content (x) of the PdFe_xCo_1−xO₂ film is increased, a more pronounced drop in resistivity is seen. Derivatives of the resistivity as a function of temperature of these PdFe_xCo_1−xO₂ films are shown in Fig. 5(e). Strikingly, a dip at about 20 K is observed in the temperature derivatives of all iron-doped PdCoO₂ films, which is opposite to the dρ/dT in PdCrO₂ where a peak shows up at T_N driven by AFM order.¹⁰ The amplitude of the dip observed for PdFe_xCo_1−xO₂ increases with larger iron concentration.

A comparison of the temperature dependence of the Hall coefficient (R_H) between a PdCoO₂ film and a PdFe_0.17Co_0.83O₂ film is shown in Fig. 6(a). The R_H measurements are consistent with electrons acting as the carriers in both PdCoO₂ and PdFe_0.17Co_0.83O₂ films. The magnitude of R_H in the PdCoO₂ film is in agreement with prior reports from PdCoO₂ single crystals.⁵⁷ In contrast to Ref. 39, we do not observe an anomalous Hall effect in our PdCoO₂ films at low temperature. Hall resistivities of the PdCoO₂ film and the PdFe_0.17Co_0.83O₂ film are shown in Fig. S9 of the supplementary material. The PdFe_0.17Co_0.83O₂ film exhibits a larger R_H than does the pure PdCoO₂ film, which could be a result of carriers being trapped by iron-induced structural defects. For the PdCoO₂ film, the temperature dependence of R_H at low temperature becomes flat while for the iron-doped PdCoO₂ film the R_H starts increasing below 20 K, which is the same temperature at which the change in dρ/dT is observed in Fig. 5. One possibility for the observed resistivity anomaly at low temperature is that it is driven by iron disorder, since it is independent of the iron concentration. One scenario explaining why the R_H difference between of PdCoO₂ film and the PdFe_0.17Co_0.83O₂ film does not reflect on the band structure is that the electrons from iron do not interact with the electrons from palladium. Instead, iron clusters just trap the electrons from the PdCoO₂. Within this scenario, it is possible that iron disorder clusters in PdFe_0.17Co_0.83O₂ films are revealed by AFM in the supplementary material (Fig. S3). In Figs. 6(c) and 6(d), the magnetoresistance of the same PdCoO₂ film and PdFe_0.17Co_0.83O₂ film shows distinct magnetic dependences. The overall scale of magnetoresistance in the PdCoO₂ film is much smaller than that observed in PdCoO₂ single crystals.⁵⁸ The temperature dependence of the magnetoresistance of the PdFe_0.17Co_0.83O₂ film shows weak-localization-like behavior, which may arise from the magnetic disorder resulting from the addition of iron. The temperature dependence of the magnetic susceptibility of the PdFe_0.17Co_0.83O₂ thin film shows no transition or difference between the zero-field-cooled (ZFC) and field-cooled (FC) curves as shown in the supplementary material (Fig. S10). The observed behavior is in contrast to the splitting that is expected when AFM order is observed, such as in PdCrO₂.¹¹ Thus, the replacement of cobalt by iron does not appear to result in any spin order. Instead, only signs of magnetic disorder are seen.

In summary, we have synthesized high-quality PdCoO₂ films by MBE and harnessed the atomic layer control afforded by MBE to tune the termination of these films to study the resulting electronic structure by ARPES. On comparing the Pd-terminated and CoO₂-terminated PdCoO₂ films with those of PdCoO₂ single crystals, though the resistivity of our PdCoO₂ films are far higher than that of single crystals at low temperature, we find the PdCoO₂ bulk states in our films show features similar to those of PdCoO₂ single crystals, while the palladium surface state and CoO₂ surface state are not as strong as those of the PdCoO₂ single crystals. This difference might arise due to different electronic reconstructions. We also studied FeO₂-terminated PdCoO₂ films and find that the only remaining PdCoO₂ bulk state in the electronic structure is similar to that of PdCoO₂. In addition, we have successfully synthesized PdFe_xCo_1−xO₂ films. From the electric transport measurements, the addition of iron further increases the resistivity of PdCoO₂ films at room temperature. Meanwhile, we see different behavior at low temperature compared to pure PdCoO₂ films, but no magnetic ordering akin to what happens in PdCrO₂ is seen in our PdFe_xCo_1−xO₂ films. The electronic structure of a PdFe_0.17Co_0.83O₂ film shows a bulk state similar to that seen in pure PdCoO₂ films. Although we do not induce new spin order in delafossite films by replacing cobalt by iron in PdFe_xCo_1−xO₂ with x up to 0.2 or a pure FeO₂ terminating monolayer, our work invites further exploration of ways in which the electronic structure of delafossites can be perturbed by exploiting the ability of MBE to control atomic layering in combination with ARPES to measure its impact.

See the supplementary material for the growth methods of the PdFe_xCo_1−xO₂ (0 ≤ x ≤ 0.2) films and additional characterization by AFM, LEED, ARPES, and XPS, as well as the results of Hall effect and magnetization measurements.

The authors thank Alfred Zong, Juan Jiang, Yu He, and Liguo Ma for their insightful discussions. This paper was primarily supported by the U.S. Department of Energy, office of Basic Sciences, Division of Materials Science and Engineering under Award No. DE-SC0002334. This paper made use of the Cornell Center for Materials Research shared facilities, which are supported through the NSF Materials Research Science and Engineering Centers Program (Grant No. DMR-1719875). This paper also made use of the Cornell Energy Systems Institute Shared Facilities partly sponsored by the NSF (Grant No. MRI DMR-1338010) and the Kavli Institute at Cornell. Substrate preparation was performed, in part, at the Cornell NanoScale Facility, a member of the National Nanotechnology Coordinated Infrastructure, which is supported by the NSF (Grant No. NNCI-2025233). C.T.P. acknowledges support from Air Force Office of Scientific Research Grant No. FA9550-21-1-0168 and National Science Foundation Grant No. DMR-2104427. P.D.C.K. gratefully acknowledges support from the European Research Council (through the QUESTDO project, Grant No. 714193). Q.X. acknowledges support from the REU Site: Summer Research Program at PARADIM (Grant No. DMR-2150446). The authors thank Sean C. Palmer for his assistance with substrate preparation.