α-Ta films on c-plane sapphire with enhanced microstructure Open Access

Epitaxial interfaces between metals and oxides play a key role in a wide range of applications in electronics and catalysis. However, complex interfacial reactions, wetting, adhesion, and band-bending between the metal film and the oxide substrate often lead to difficulties in achieving epitaxial growth with low defect densities in the film, in the substrate, and at their interface.^1,2

Driven by their potential use as superconducting layers in qubits, epitaxial films of tantalum are receiving increased attention. Tantalum crystallizes in two distinct crystallographic phases—a body-centered-cubic (bcc) α-phase and a tetragonal β-phase. The α-phase (lattice constant a = 3.3 Å) is a superconductor with a critical temperature (T_c) of 4.4 K, whereas the β-phase (lattice constants a = 5.3 Å and c = 9.9 Å) has a T_c of 0.5 K.³ The comparatively high T_c of the α-phase is of interest for its use in quantum electronics. Indeed, there has been a breakthrough in the fabrication of long-lifetime transmon qubits made of epitaxial α-Ta films.^4–6 In addition, α-Ta films offer a stable optical response to hydrogen over a large pressure range, making α-Ta a promising sensor material in the ever-growing hydrogen economy.^7–9 

However, the epitaxial growth of α-Ta films is demanding. The low vapor pressure of Ta implies that the deposition of Ta thin films is challenging for many film-growth techniques.¹⁰ Therefore, Ta films are mostly grown by sputter deposition.^{3–9,11–20} The resulting films, however, often suffer from the (partial) formation of the unwanted β-phase,^{3,6,11,12,14–18} from a mixed (110) (111) orientation,^3,6,15,17,19 from a high interface roughness,^4,5,15,16 and from a large number of domain or grain boundaries.^6,13,16,17 The latter has been shown to result in a decrease in the coherence length,¹³ which is potentially harmful to the performance of qubits. The low loss at high frequencies of sapphire suggests it as a substrate of choice for qubit applications.²¹ The growth of epitaxial Ta on a-plane sapphire instead of the commonly used c-plane sapphire has proven to be a successful strategy to achieve a pure (110) orientation. However, these films are still characterized by small domains and a correspondingly high density of domain boundaries.⁶ For growth on c-plane sapphire, the use of a few monolayers of Nb as a nucleation layer for the Ta film growth also yielded pure (110) Ta films and furthermore allowed a drastic reduction in the temperature necessary to suppress the formation of the β-phase. The Nb nucleation layer, however, negatively affected the crystal quality.¹⁷ 

In this work, we present a synthesis route providing on c-plane sapphire epitaxial α-Ta films, which are phase-pure, have a single out-of-plane orientation, a preferred in-plane orientation, a drastically increased grain size compared to conventional Ta films, smooth interfaces to the substrate, and smooth surfaces, and are of remarkable crystal perfection. The epitaxial Ta films are grown by thermal laser epitaxy (TLE).²² TLE offers two inherent advantages for the epitaxial growth of α-Ta: Ta is readily evaporated by the continuous wave (cw) source laser using a wavelength of ≈1 μm,^10,23 and high growth temperatures are readily achieved by laser heating the sapphire substrates using a wavelength of ≈10 μm.¹⁰ As we will show in this work, it is these high achievable temperatures that are key to epitaxially growing α-Ta films of high quality on c-plane sapphire.

The TLE used in this work is equipped with a TRUMPF fiber laser with λ = 1.07 μm and a spot size of 1 mm², operated at an output power P_L = 400 W to evaporate free-standing Ta cylinders. These cylinders have a purity of 3N8, a diameter of 3 mm, and a length of 6 cm. With a 60 mm distance between the hot Ta surface and the substrate, a growth rate of ≈0.8 Å/s is achieved under these conditions. A small melt pool at the source surface indicates a source temperature of ≥3000 °C. The chamber is equipped with a LN2 shroud such that even during deposition, the background pressure is <4 × 10⁻¹⁰ hPa. The Crystek 5 × 5 mm² c-plane sapphire substrates are heated by irradiating the backside with the λ ≈ 10 μm beam of a CO₂ laser; the substrate temperature T_sub is measured using a backside pyrometer. Sapphire substrates are prepared as described in Ref. 24. Films considered in this work have thicknesses between 50 and 120 nm as determined by x-ray reflectivity (XRR), Laue oscillations, and a Bruker DektakXT depth profilometer. Before and after growth, the samples are characterized by in situ reflection high-energy electron diffraction (RHEED). During growth, RHEED is turned off to avoid damage to the sample.^25–27 

After growth, the surface of the films is evaluated using an Asylum Cypher atomic force microscope (AFM) in tapping mode. θ–2θ scans, rocking curves, and XRR data were acquired by a PANalytical Empyrean diffractometer with monochromated Cu-K_α1 x rays. The hybrid monochromator is operated with a 1/16° divergence slit and 2 mm mask. A PIXcel 3D detector was operated with a 0D 0.5 mm receiving slit and a 1/4° anti-scatter slit. The time per step was 0.2 s, and the step size was 0.02°. XRR data were fitted with the X’Pert reflectivity software. The pole figures were acquired using a Bruker AXS D8 Discover diffractometer with an auto-absorber and a LYNXEYE detector. The step size was 0.02 in, and the time per step was 0.2 s. Electron-backscattering diffraction (EBSD) data were acquired using a Zeiss Merlin secondary electron microscope equipped with an Oxford symmetry detector driven by the Oxford AZtec software. Scanning transmission electron microscopy (STEM) specimens in the cross-sectional orientation were prepared by tripod polishing followed by argon-ion milling at liquid-N2 temperature.²⁸ STEM studies were conducted using a spherical aberration-corrected STEM (JEM-ARM200F, JEOL Co. Ltd.) equipped with a cold-field emission gun and a DCOR probe Cs-corrector (CEOS GmbH) operated at 200 kV. The STEM images were obtained using an ADF detector with a convergent semi-angle of 20.4 mrad and collection semi-angles of 70–300 mrad. In order to make precise measurements of lattice constants, ten serial frames were acquired with a short dwell time (2 μs/pixel), aligned, and added afterward to improve the signal-to-noise ratio and to minimize the image distortion of high-angle annular dark-field (HAADF) images. Electron energy loss spectroscopy (EELS) acquisition was performed using a Gatan GIF Quantum ERS imaging filter equipped with a Gatan K2 Summit camera with a convergent semi-angle of 20.4 mrad and a collection semi-angle of 111 mrad. Temperature dependent resistivity was measured in van-der-Pauw geometry utilizing a quantum design physical property measurement system. The data were taken with a cooldown rate of 4 K/min from 300 to 20 K and then a cooldown rate of 0.4 K/min to precisely record the phase transition.

It is well established that high substrate temperatures (T_sub) are beneficial for the deposition of phase pure α-Ta thin films.¹⁸  Figure 1 shows the surface quality of 65 nm-thick Ta films grown at increasing T_sub. Below 600 °C amorphous growth is observed. For 600 °C ≤ T_sub ≤ 900 °C, the RHEED pattern consists of transmission spots, indicating three-dimensional (3D) growth [Fig. 1(a)]. The resulting surface reveals small grains and a root mean square (RMS) value of = 2.95 nm [Fig. 1(c)]. At T_sub ≥ 1050 °C, the RHEED pattern drastically changes. For growth at such high temperatures, the RHEED pattern comprises sharp surface-diffraction streaks and spots, indicating a smooth, coherent two-dimensional surface structure. Moreover, these patterns reveal a reconstruction on the epitaxially grown α-Ta film. The spots span six times the fundamental spacing, with five regular reconstruction streaks in between each one, thereby revealing a sixfold superstructure.

The resulting surface topography of the film grown at 1050 °C is comparatively smooth with an RMS = 1.42 nm [Fig. 1(d)]. The grains are about ten times wider than the film thickness, and their lateral dimension is by far larger than that of state-of-the-art films.^4,17 The elongated grains are aligned along three in-plane orientations, which differ by 120° (compare Fig. 4).

Upon increasing T_sub further to 1150 °C, the grains grow even larger to attain micrometer sizes and interconnect [Fig. 1(e)], resulting in a reduced surface roughness with RMS = 0.52 nm.

For T_sub = 1250 °C, the grains are even larger, but they get separated by deep trenches. Interestingly, on top of these individual grains, steps of ≈2.3 Å are visible in the AFM micrographs [Fig. 1(f)]; this step height corresponds to the d₁₁₀ spacing of α-Ta. The surface morphologies of films grown at T_sub > 1250 °C become more irregular (Fig. S1 of the supplementary material).

Having found that T_sub = 1150 °C results in smooth and closed films of α-Ta, we now turn to the question whether these films also have favorable microstructural properties. The XRD θ–2θ scans taken on such films [Fig. 2(a)] reveal a (111) orientation for the film grown at T_sub = 900 °C. At T_sub = 1050 °C, additionally, the expected 110 and 220 peaks of epitaxial α-Ta are observed. For T_sub = 1150 and 1250 °C, only 110 and 220 peaks of epitaxial α-Ta are observed. These measurements thus uncover that for the growth conditions used, growth at T_sub > 1050 °C is required to grow α-Ta in single out-of-plane orientation on c-plane sapphire. Despite the mixed orientation of the films grown at T_sub = 1050 °C, all films were found to exhibit distinct Laue oscillations, which indicate a sharp substrate–film interface, as well as a good crystallinity of the film, and a homogeneous film thickness [Fig. 2(b)].²⁹  Figure 2(b) further reveals a slight shift of the 110 peak to lower angles with increasing T_sub. In Table I, the resulting lattice parameters are compared to the bulk value.³⁰ The slight increase in the lattice parameter is likely the result of the increased formation of Ta vacancies at increased T_sub, which are frozen by the rapid cooldown.

Figure 2(c) shows the rocking curves measured at the 110 peak for all films exhibiting this orientation. Astoundingly, the rocking curve peak is very sharp for all three T_sub. Indeed, its full-width half-maximum (FWHM) values are 0.038° ≤ Δω ≤ 0.041°. The rocking curve peaks are, thus, more than two times narrower than those of the best Ta-films published, which actually were obtained on a-plane sapphire.³¹ Interestingly, grain size and single orientation seem to not significantly impact the perfection of the crystal quantified by the FWHM.

As expected from the distinct Laue oscillations, XRR reveals extensive Kiessig fringes for the films grown at T_sub ≥ 1050 °C. These are a testament to the sharp substrate–film interface and the high quality of the film surface. Thus, it is no surprise that the fits of the XRR reveal a low Ta roughness (Table II). More interestingly, to fit the data, it is necessary to assume a thin oxide layer on top, which is expected to form as Ta gets exposed to air. Independent of the Ta T_sub, the oxide layer is ≈2.5 nm thick. For the fits, the formation of Ta₂O₅ was assumed. However, the density of the native oxide formed on the films does not correspond to that of Ta₂O₅ (8.2 g/cm³). As also the expected diffraction pattern of Ta₂O₅ is not observed, the thin native oxide layer is likely amorphous with a mixed oxidation state.

To gain insights into the epitaxial relationship between the film and the substrate and to decipher the in-plane orientation of the films, we present in Fig. 3 the XRD pole figures of a film grown at T_sub = 1150 °C and a substrate. The pole figures clearly reveal a Pitsch–Schrader epitaxial relation³² with very sharp reflections. Patterns reflecting other orientations¹⁹ are completely absent. For the Pitsch–Schrader orientation of (110) Ta on (0001) Al₂O₃, the in-plane direction [001] of the bcc Ta is parallel to the sapphire direction $11̄00$] or its equivalents $011̄0$] and $1̄010$], yielding three degenerate in-plane orientations. The corresponding arrangements of the atomic lattices are schematically depicted in Fig. 4. The Ta in-plane vectors $11̄l$] and $1̄1l$] (l = 1 or $1̄$⁠) are at an angle of 5.26° to the $1̄010$] direction or its equivalents of the hcp c-plane sapphire substrate. Interestingly, as shown in Fig. 3(a) for T_sub = 1150 °C, the diffraction spots at ϕ = 0° and ϕ = 180° are more pronounced than the others, disclosing that one orientation is preferred over the other two. The films grown at 1050 and 1250 °C have no strongly preferred orientation (Fig. S2 of the supplementary material). This can be quantified by the relative intensities of the underlying ϕ-scans found in Table III. The peaks are labeled as in Fig. S3 of the supplementary material. The degeneracy of the in-plane orientations is thus lifted in samples grown at T_sub = 1150 °C, which again shows that the temperature window for the ideal growth of epitaxial α-Ta is narrow. EBSD data (Fig. S4 of the supplementary material) show that the domains with preferred orientation merge to form very large single-crystalline domains. This may provide a starting point for realizing single-crystalline bcc films on hexagonal or triangular substrates.

Direct atomic-scale information on the crystal lattices and the substrate–film interfaces has been obtained from extensive STEM studies of the films. Figure 5(a) shows a wide-field HAADF image of a cross section through the film and substrate. The micrograph reveals a uniform thickness. The image also shows faint white lines, which likely indicate grain boundaries that meet at a 120° angle. The atomic resolution images shown in Figs. 5(c) and 5(d) directly confirm the high crystal quality on the nanoscale as well as the atomically sharp interface between the substrate and film. The rotation angle between the two images is ≈5.2°, providing direct real-space evidence for the film–substrate atomic arrangement depicted in Fig. 4. For clarity, the atomic arrangement of Fig. 5(c) is additionally depicted in Fig. 5(e). The atomic sharpness of the interface is also revealed by the EELS maps in Fig. 5(b), in which no intermixing between the film and substrate ions can be discerned, despite the high T_sub at which the films were grown.

Figure 6 shows the resistivity of the α-Ta films grown at different T_sub as a function of temperature. For comparison, this figure also shows the behavior of a bulk α-Ta sample as published in Ref. 33. Films grown at T_sub ≥ 1050 °C show T_c close to the bulk value. In fact, T_c of the samples grown at 1050 and 1150 °C is the same as the bulk value (≈4.4 K, see the inset of Fig. 6). It is to be noted that the sample grown at 1050 °C is by far the thickest sample (120 nm). Of all samples, it has the most noteworthy residual resistivity ratio (RRR) ρ_{300 K}/ρ_{5 K} of 33, which is much higher than any literature values for thin films. At the identified ideal growth temperature (1150 °C), the RRR ρ_{300 K}/ρ_{5 K} of the 80 nm thick film is 13. This value is comparable to literature values.⁶ In our case, however, this value is exhibited by films that are an order of magnitude thinner than the reference films, which greatly increases the relative contribution of the electron scattering at the film surface and at the film–substrate interface to the sample resistance at low temperatures. The RRR ρ_{300 K}/ρ_{5 K} of the film grown at 1250 °C is markedly worse (4.75), and its T_c is lowered. The observed grain separation at this temperature, thus, has a negative influence on the superconducting properties. The film grown at 900 °C, which is 111 oriented, has a much higher resistivity and a low RRR ρ_{300 K}/ρ_{5 K} (1.6), and its T_c is shifted to the limit of the measurement setup. This underlines the importance of the ability to grow 110 oriented α-Ta films at T_sub ≥ 1050 °C. The partial 111 orientation at T_sub = 1050 °C, however, has no significant negative effect on the superconducting properties. On the contrary, the RRR for this film is remarkably high.

This aim of this work is to find a new route to the epitaxial growth of α-Ta thin films, to avoid growth-related deficits. These defects include unwanted β-phase formation, a mixed (110) (111) orientation, a high interface roughness, and large numbers of grain boundaries. We have found that TLE offers the necessary parameter range for the growth of defect-poor α-Ta films, as Ta is readily evaporated by laser heating and high substrate temperatures are achievable by CO₂-laser heating while maintaining low background gas pressures.¹⁰ Our studies establish that substrate temperatures T_sub > 1050 °C are required to ensure a single (110) orientation of α-Ta films. The preferred substrate temperature for the growth of Ta films that have large domains, an atomically sharp interface with the underlying substrate, a singular out-of-plane and a preferred in-plane orientation is T_sub = 1150 °C. These films are closed, smooth, and free of β-Ta impurities.

These results have the potential to further drive the application of α-Ta in (quantum) electronics, where a high-quality film growth is generally desirable. Furthermore, high quality epitaxial α-Ta films may be a valuable starting point to unravel the impact of growth induced imperfections on optical hydrogen sensing capabilities.

See the supplementary material for the following: Fig. S1 shows the surface structure and RHEED of a Ta film grown at 1500 °C; Fig. S2 shows the pole figures of the Ta films grown at 1050 and 1150 °C; Fig. S3 shows Φ-scans extracted from Fig. 3; and Fig. S4 shows EBSD results for a Ta film grown at 1150 °C.

The authors thank Hans Boschker for scientific discussions and Gunther Richter and Gerd Maier for experimental support. The authors thank Ute Salzberger for assistance with TEM sample preparation.

The authors have no conflicts to disclose.

Lena N. Majer: Conceptualization (equal); Data curation (lead); Formal analysis (lead); Investigation (equal); Methodology (equal); Visualization (lead); Writing – review & editing (equal). Sander Smink: Conceptualization (equal); Data curation (equal); Formal analysis (equal); Investigation (equal); Supervision (equal); Visualization (equal); Writing – review & editing (equal). Wolfgang Braun: Conceptualization (equal); Data curation (equal); Formal analysis (equal); Investigation (equal); Methodology (equal); Supervision (equal); Visualization (equal); Writing – original draft (lead). Bernhard Fenk: Data curation (supporting); Methodology (supporting). Varun Harbola: Data curation (equal); Formal analysis (equal); Investigation (equal); Methodology (equal); Visualization (equal); Writing – review & editing (equal). Benjamin Stuhlhofer: Data curation (supporting); Methodology (supporting). Hongguang Wang: Data curation (supporting); Formal analysis (equal); Investigation (equal); Methodology (equal); Visualization (equal); Writing – review & editing (equal). Peter A. van Aken: Methodology (equal); Supervision (equal); Writing – review & editing (equal). Jochen Mannhart: Conceptualization (equal); Supervision (equal); Writing – review & editing (lead). Felix V. E. Hensling: Conceptualization (equal); Data curation (equal); Formal analysis (equal); Investigation (equal); Methodology (equal); Supervision (equal); Visualization (supporting); Writing – original draft (lead).

The data supporting the finding of this study are available within the paper, its supplementary material, and upon reasonable request.