Intermetallic Nb3Sn alloys have long been believed to form through Sn diffusion into Nb. However, our observations of significant oxygen content in Nb3Sn prompted an investigation of alternative formation mechanisms. Through experiments involving different oxide interfaces (clean HF-treated, native oxidized, and anodized), we demonstrate a thermodynamic route that fundamentally challenges the conventional Sn diffusion mechanism for Nb3Sn nucleation. Our results highlight the critical involvement of a SnOx intermediate phase. This new nucleation mechanism identifies the principles for growth optimization and new synthesis of high-quality Nb3Sn superconductors.

Nb3Sn1–3 is a crucial and promising superconductor that finds widespread applications in various modern technologies, including superconducting radio-frequency (SRF) resonators,4–6 high-field magnets,7 and emerging electronics, such as quantum devices and high-capability detectors.8 Of particular interest are SRF resonant cavities that accelerate charged particle beams by coupling RF superconductors with RF waves. These cavities are critical components for modern particle accelerators, with applications spanning from photon science,9–11 high-energy and nuclear physics,12,13 advanced materials discovery,9,10 isotope production,14 to quantum computing.15 

For SRF applications, Nb3Sn possesses advantages over the conventional material niobium (Nb). Nb3Sn exhibits a high predicted superheating field (400 mT), twice that of Nb (200 mT), allowing for large accelerating gradients up to 100 MV/m.16 Nb3Sn has a higher critical temperature (Tc) of 18 K compared to Nb’s 9.2 K, resulting in lower surface resistance and higher quality factors. The higher Tc of Nb3Sn enables the replacement of costly helium cryogenics (operated at 2 K) with commercial cryocoolers (operated at 4.2 K), thereby facilitating the development of turn-key, compact accelerators.17 With excellent capabilities, Nb3Sn-based cavities are poised to advance toward practical accelerator applications.18 

To date, efficient SRF cavities that utilize Nb3Sn films are dominantly produced through tin (Sn) vapor diffusion on Nb surfaces.18–21 Recently, a new process involving Sn electroplating on Nb followed by thermal annealing has been demonstrated,22 with other methods, e.g., sputtering, being explored.

The vapor diffusion process for synthesizing Nb3Sn involves heating Nb in the presence of vaporized Sn at ∼500 °C for ∼4 h for nucleation, followed by heating at ∼1100 °C for 2–3 h for grain growth. To optimize the process, various improvements have been implemented, including switching from native Nb-oxidized surfaces to anodized surfaces and from the pure Sn source to an additional high-vapor-pressure SnCl2 source. Regardless of these changes, the observation of thick oxides and high subsurface oxygen concentrations in vapor-diffused Nb3Sn22,23 has raised questions about the source of oxygen during nucleation.

Since the 1980s, the proposed mechanism for intermetallic formation of Nb3Sn has been based on the diffusion of Sn through the grain boundary of Nb or an initially nucleated thin interface of NbxSny. This leads to further growth of Nb3Sn or NbxSny and slow bulk diffusion to produce Nb3Sn.24,25 Recent atomic simulations have confirmed that grain boundary diffusion in Nb3Sn is significantly faster than bulk diffusion, e.g., 105 times faster at 1100 °C.26,27 This diffusion-based mechanism has been widely used to analyze and optimize Nb3Sn growth.20,28–31

The early stage of Nb3Sn nucleation has been investigated using the Sn vapor process.29–31 This study stands apart from previous research in several ways, challenging existing assumptions.

First, previous investigations followed the assumption that Sn is the primary component involved in the nucleation process and identified a pure kinetic problem to address non-uniform nucleation issues. However, this work reveals that the thermodynamic route actually involves an intermediate phase of SnOx in the nucleation process.

Second, vapor-based investigations observed the formation of droplet regions, likely representing the initial nuclei.29–31 However, accurate identification of these droplets remains unverified due to several challenges in characterization. Electron dispersive spectroscopy analysis of these droplets is plagued by unreliable measurements caused by interference from oxygen signals. Cross-sectional transmission electron microscopy imaging poses challenges: sample preparation involving focused ion beam polishing under 5 × 10−7–10−6 Torr introduces errors related to the presence of residual oxygen in the chamber; the presence of non-crystalline oxides further complicates the identification of these nuclei.

Here, we design the sample architectures to ensure a good capping of the interface between the precursor SnCl2 and the substrate, preventing exposure to air. Probing the interface using X-ray photoelectron spectroscopy (XPS) depth profiling under 10−10–10−9 Torr provides semi-in situ identification of the nuclei with more reliable results to date.

Three different layouts of SnCl2/(NbxOy)/Nb were fabricated: (i) a clean Nb surface treated with hydrofluoric acid (HF) followed by SnCl2 deposition (SnCl2/Nb); (ii) a native oxidized surface comprising a ∼7 nm oxide stack23 followed by SnCl2 deposition (SnCl2/thin, native oxide/Nb); and (iii) an anodized surface with a 60 nm-thick, porous oxide layer followed by SnCl2 deposition (SnCl2/thick, porous oxide/Nb).

High-purity Nb substrates were used with a large grain size of > 50 mm and a residual-resistivity ratio of >300. The 1 × 1 cm2 substrates were electropolished using a mixture of 1:9 HF (48%) and sulfuric (98%) acids, resulting in an average surface roughness of 40 nm after a 100 μm polishing. The anodization process utilized the electropolished Nb as the anode and Pt as the cathode in a two-electrode control. The solution used was concentrated sodium hydroxide, and the voltage applied was up to 30 V, producing porous oxides as imaged by scanning electron microscope (SEM) in Fig. 1(a).

FIG. 1.

(a) SEM surface image and (b) XRD pattern of the anodized Nb.

FIG. 1.

(a) SEM surface image and (b) XRD pattern of the anodized Nb.

Close modal

Afterward, the substrates were transferred to a nitrogen-flown glovebox with O2 and H2O levels below 0.5 ppm, and all subsequent operations were performed in the glovebox. One type of substrate was soaked in 2% HF for 30 min and dried. SnCl2 dihydrate microparticles were dissolved in methanol at a concentration of 1 mg/ml. The drop casting technique was employed to deposit 2 ml of SnCl2 solution onto the substrate surface, followed by solvent evaporation. Uniform coverage of white films was observed on the anodized sample, while it took several attempts to obtain comparable coverage for the other two conditions, which showed poor adhesion.

Following baseline characterizations, the samples underwent annealing in a Lindberg tube furnace under a vacuum of 5 × 10−7 Torr at 300 °C for 4 h. While this temperature is lower than the typical nucleation temperature for vapor diffusion (500 °C for 4 h under ∼10−6 Torr), it was chosen to capture the early stages of transformation.

Chemical states, elemental compositions, and phase structures were determined using X-ray photoelectron spectroscopy (XPS) with depth profiling and high-resolution X-ray diffraction (XRD). To accommodate thick films reaching up to 30 μm, XPS survey scans were performed with sputtering intervals ranging from 1 to 327 nm using a Surface Science Instruments SSX-100 ESCA spectrometer. Monochromatic Al k-alpha X-ray (1486.6 eV) photoelectrons were collected from an 800 μm analysis spot at a 55° emission angle under a vacuum of 10−9 Torr. The scan parameters were set to 150 eV pass energy, 1 eV step size, and 100 s/step. For depth profiles, a 4 kV Ar+ beam with a spot size of ∼5 mm was rastered over a 2 × 4 mm2 area.

Phase structures were further analyzed using a Rigaku SmartLab XRD instrument. The X-ray generated by a Cu target was converted into a parallel, monochromatic beam and directed onto the sample surface after passing through a 5 mm divergence slit. The Kα signal with a wavelength of 0.154 nm was collected by filtering the diffracted beam. The parallel slit analyzer and 5° Soller slits were used to resolve high-resolution signals. The 2θ scan was performed with a step size of 0.0032°.

To gain insight into the structural changes and reactions between multiple elements, XPS spectra of Sn 3d, Nb 3p, O 1s, and Cl 2p (e.g., Fig. 2) were collected as a function of depth for the as-deposited and 300 °C annealed samples. The binding energies are determined with reference to Table I. To accurately interpret the results, we combine the information from concentration profiles (Fig. 3), FWHM (full-width-at-half-maximum) changes (Fig. 4), and binding energy shifts (Fig. 5), while accounting for the sputtering effects and resolution challenges.

FIG. 2.

XPS spectra for the anodized sample after 300 °C annealing, taken at a depth of 200 nm. The intensity units are arbitrary.

FIG. 2.

XPS spectra for the anodized sample after 300 °C annealing, taken at a depth of 200 nm. The intensity units are arbitrary.

Close modal
TABLE I.

Binding energies (eV) of relevant motifs in the literature.23,32–35

Sn 3d5/2484–485 (metallic)486–487 (SnO/SnO2)488 (SnCl2)491.2 (charging)
Nb 3p3/2 360.5–361 (metallic) 363–365 (Nb oxides) 365 (NbCl5 
O 1s 529.5–530.5 (oxides) 532–534 (hydroxides) > 537 (charging)  
Cl 2p3/2 198.5–199.5 (chlorides)    
Sn 3d5/2484–485 (metallic)486–487 (SnO/SnO2)488 (SnCl2)491.2 (charging)
Nb 3p3/2 360.5–361 (metallic) 363–365 (Nb oxides) 365 (NbCl5 
O 1s 529.5–530.5 (oxides) 532–534 (hydroxides) > 537 (charging)  
Cl 2p3/2 198.5–199.5 (chlorides)    
FIG. 3.

Atomic concentrations of Sn, Nb, O, and Cl as a function of depth for the as-deposited [(a)–(c)] and annealed [(d)–(f)] samples. [(a) and (d)] Clean Nb, [(b) and (e)] native oxidized, and [(c) and (f)] anodized. The concentration uncertainty is <5%. The depth resolution is affected by the sputtering effect.

FIG. 3.

Atomic concentrations of Sn, Nb, O, and Cl as a function of depth for the as-deposited [(a)–(c)] and annealed [(d)–(f)] samples. [(a) and (d)] Clean Nb, [(b) and (e)] native oxidized, and [(c) and (f)] anodized. The concentration uncertainty is <5%. The depth resolution is affected by the sputtering effect.

Close modal
FIG. 4.

XPS peak FWHM of (a) Sn 3d5/2 and (b) Nb 3p3/2 as a function of depth, with an intrinsic width limit of ∼2.3 eV.

FIG. 4.

XPS peak FWHM of (a) Sn 3d5/2 and (b) Nb 3p3/2 as a function of depth, with an intrinsic width limit of ∼2.3 eV.

Close modal
FIG. 5.

XPS binding energies for (a) Sn 3d5/2, (b) Nb 3p3/2, (c) O 1s, and (d) Cl 2p3/2, plotted as a function of depth. The error bars represent a conservative uncertainty estimate of 1 eV.

FIG. 5.

XPS binding energies for (a) Sn 3d5/2, (b) Nb 3p3/2, (c) O 1s, and (d) Cl 2p3/2, plotted as a function of depth. The error bars represent a conservative uncertainty estimate of 1 eV.

Close modal

The oxygen-free interface on clean Nb [Fig. 3(a)] shows the appearance of NbClx and NbxSny following SnCl2 deposition. Incorporation of these two new Nb and Sn structures results in FWHM increases up to 6.5 and 4.5 eV, respectively (Fig. 4). The 485 eV Sn binding energy strongly indicates the generation of NbxSny alloying upon deposition [Fig. 5(a)]. The 365 eV Nb binding energy is attributed to NbClx [Fig. 5(b)], as no oxygen signals were detected, ruling out the possibility of Nb oxides that have characteristic peaks in the same positions.

In contrast, the presence of oxygen in the native oxidized and anodized interfaces causes a clear disappearance of chlorine [Figs. 3(b) and 3(c)]. The formation of NbClx and NbxSny is still observed, as evidenced by the similar FWHM spikes (Fig. 4) and characteristic binding energies [Figs. 5(a) and 5(b)].

Figures 3(d)3(f) demonstrate the absence of chlorine and the conversion into Sn oxides. To support the phase identification, high-resolution XRD patterns are presented in Fig. 6 and referenced in Table II.

FIG. 6.

High-resolution XRD patterns for the 300 °C annealed samples: (a) clean Nb, (b) native oxidized, and (c) anodized. The data were processed using the GENPLOT smoothing algorithm.

FIG. 6.

High-resolution XRD patterns for the 300 °C annealed samples: (a) clean Nb, (b) native oxidized, and (c) anodized. The data were processed using the GENPLOT smoothing algorithm.

Close modal
TABLE II.

XRD diffraction angles for Nb (Im3̄m), SnO2 (P42/mnm), SnO(P4/nmm), and Nb3Sn (Pm3̄n) in the database.36 

Nb38.4° [110]55.4° [200]69.4° [211]82.2° [220]
SnO2 26.4° [110] 33.7° [101] 38.7° [111] 51.5° [211] 54.5° [220] 57.5° [002] 61.5° [310] 
 64.3° [112] 71.4° [320] 78.2° [321]     
SnO 29.7° [101] 33.2° [110] 36.8° [002] 47.7° [200] 50.4° [112] 57.2° [211] 61.9° [103] 
Nb3Sn 33.7° [200] 37.8° [210] 41.5° [211] 60.2° [320] 62.9° [320] 70.78° [400]  
Nb38.4° [110]55.4° [200]69.4° [211]82.2° [220]
SnO2 26.4° [110] 33.7° [101] 38.7° [111] 51.5° [211] 54.5° [220] 57.5° [002] 61.5° [310] 
 64.3° [112] 71.4° [320] 78.2° [321]     
SnO 29.7° [101] 33.2° [110] 36.8° [002] 47.7° [200] 50.4° [112] 57.2° [211] 61.9° [103] 
Nb3Sn 33.7° [200] 37.8° [210] 41.5° [211] 60.2° [320] 62.9° [320] 70.78° [400]  

The clean Nb sample, with a Sn concentration of ∼10 at. % and a thinness of ∼10 nm, exhibits the lowest amount of SnOx [Fig. 5(a)]. On the surface, higher-order Nb oxides of ∼3 nm are present. Underneath this layer, metallic Nb [Fig. 5(b)] is mixed with SnOx, resulting in increased Sn FWHM and decreased Nb FWHM (Fig. 4). A low amount of NbxSny may be present, as suggested in the XRD pattern [Fig. 6(a)], but the characteristic peak of metallic Sn is not prominent. The rapid decrease in oxygen concentration within ∼5 nm for this condition suggests the source of oxygen should be the residual gas in the furnace chamber with a vacuum of 5 × 10−7 Torr, which differs from the two oxidized samples where oxygen is evenly distributed at 50 at. % throughout the samples.

The native and anodized samples with initial oxides exhibit similarity in the concentration profile [Figs. 3(e) and 3(f)]. Both conditions yield significant amounts of SnOx in addition to NbxSny nuclei, as evidenced by the persistence of Sn binding energies at 487 eV associated with SnOx [Fig. 5(a)] and the observed diffraction patterns of SnO2 and SnO [Figs. 6(b) and 6(c)]. NbxSny starts to emerge at a depth of ∼200 nm, as indicated by the considerably broadened width of both Sn and Nb XPS peaks in Fig. 4. These results provide compelling evidence that SnOx is a necessary intermediate phase for alloying.

Nevertheless, the anodized sample differs from the native oxidized sample due to the presence of residual SnCl2 on the surface [Figs. 3(f) and 5(d)], resulting in a higher Sn binding energy [Fig. 5(a)] and a width spike in the Sn peak [Fig. 4(a)] at the < 200 nm surface. Within the same depth range, the anodized oxides persist and convert to a larger quantity of higher-order oxides, as evidenced by the 365 eV binding energies [Fig. 5(b)] and the largest Nb peak width with all charge states [Fig. 4(b)]. The adoption of anodization for vapor diffusion produces a smoother Nb3Sn surface. Our data suggest that the underlying mechanisms involve two aspects: (i) a larger and more uniform adsorption of SnCl2 due to the porous oxide structures, as shown in Fig. 3(c), and (ii) a prolonged reaction time with additional SnCl2 source at a later stage and additional Nb oxides for producing the intermediate SnOx phases necessary for generating the intermetallic alloy.

The early stage of NbxSny nucleation (e.g., at 300 °C) crucially determines the quality of the final Nb3Sn product, especially in applications involving vapor diffusion in SRF cavities. Achieving a smooth surface in Nb3Sn relies on the uniformity of the initial nuclei, which serve as a seed layer for promoting uniform grain growth. Moreover, Nb3−nSn has stoichiometric boundaries of 18 and 25 at. % Sn, with the desired stoichiometry of 25 at. % Sn corresponding to the critical temperature limit of 18 K. Correcting the nucleation from an initial 18 at. % stoichiometry to the desired 25 at. % is challenging. Previous nucleation studies assumed Sn to be the main component involved, considering it to be a purely kinetic problem dependent on the availability of sufficient Sn supply.29–31 However, our observations challenge the application of this assumption to the nucleation process, while it remains unclear to what extent pure Sn diffusion, without the assistance of oxygen, contributes to subsequent grain growth.

Our findings reveal the critical and dominant involvement of SnOx as a significant intermediate step during the early nucleation stage of NbxSny. Our experimental design has addressed several confusions related to nucleation during vapor diffusion while minimizing uncertainties associated with previous studies. The results directly clarify previous observations of nuclei-like droplet features, providing clear evidence that these features are not Sn droplets, but rather SnOx entities that contain converted NbxSny. This rectifies several proposed reactions that aimed to balance the formulas with Sn as the end product, e.g., in Ref. 31. In future work, thermodynamic considerations should be prioritized, and models for SnOx/Nb and SnOx/NbOx should be constructed in experiments and simulations.

These findings also explain the high oxygen content that remains in the final Nb3Sn product. Moreover, while investigating grain growth in Nb3Sn is beyond the scope of this study, the connection between the high oxygen content in Nb3Sn and the involvement of SnOx implies that oxygen is highly likely involved in the grain growth process. This point should be further verified.

Furthermore, in addition to providing insights into vapor diffusion, this work has the potential to open up new avenues for generating Nb3Sn, moving beyond vapor diffusion.

To summarize, this work demonstrates the thermodynamic route for Nb3Sn nucleation. Our results reveal the crucial role of a SnOx intermediate phase, leading to alloy nucleation. This mechanism is particularly pronounced at the native oxidized and anodized interfaces, as well as clean Nb interfaces due to residual oxygen in the annealing vacuum furnace (5 × 10−7 Torr). Our results also suggest the anodized surfaces that produce smoother Nb3Sn rely on increasing the quantity and even distribution of adsorbed SnCl2, and enhancing the reaction kinetics through excessive SnCl2 and Nb oxide sources. These insights into the nucleation mechanisms of intermetallic Nb3Sn alloys will guide the development of optimized growth processes and new synthesis recipes for producing high-quality Nb3Sn superconductors.

This work was supported by the U.S. National Science Foundation under Award No. PHY-1549132, the Center for Bright Beams. This work made use of the Cornell Center for Materials Research Shared Facilities which are supported through the NSF MRSEC program (Grant No. DMR-1719875).

The authors have no conflicts to disclose.

Zeming Sun: Conceptualization (lead); Data curation (lead); Formal analysis (lead); Investigation (lead); Methodology (lead); Validation (lead); Writing – original draft (lead); Writing – review & editing (lead). Darrah K. Dare: Investigation (equal); Methodology (supporting); Writing – review & editing (equal). Zhaslan Baraissov: Investigation (supporting); Methodology (supporting). David A. Muller: Funding acquisition (supporting); Project administration (supporting); Supervision (supporting). Michael O. Thompson: Formal analysis (supporting); Investigation (supporting); Methodology (supporting); Writing – review & editing (supporting). Matthias U. Liepe: Investigation (equal); Methodology (supporting); Project administration (lead); Supervision (lead); Writing – review & editing (supporting).

The data that support the findings of this study are available from the corresponding author upon reasonable request.

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