Screening of 3d metals as A-elements in MAX phase Nb₂SnC and their effects on the magnetic properties of the solid solutions of Nb₂(Sn_1−xA_x)C

Abstract

The chemical versatility of MAX phases has increased almost exponentially over the last two decades, especially because of the synthesis/discovery of new solid solution phases. Many elements challenge the traditional ternary MAX phase compositions because they can be incorporated into the structure by alloying with another element, despite not forming MAX phases by themselves. Examples are (mid-to-late) transition metals that can adopt the A-site in MAX phases, some even as the sole A-element (Fe through Zn). For solid solutions, Sn has proven to be a willing A-site partner for many transition metals (Mn, Fe, Co, and Ni), even for more than one at a time. Where is the limit of transition metals that form A-site solid solutions? In this work, we demonstrate a screening of all 3d metals, except for Sc, partially substituting Sn in MAX phase Nb₂SnC. We show the successful incorporation of up to 40% of V to Zn 3d metals (except for Ti), where V and Cr occupy the A-site of the MAX phase, which is demonstrated for the first time. The full series of solid solution Nb₂(Sn_1−xA_x)C with A = V, …, Zn is studied in detail by diffraction, microscopy, and spectroscopy techniques, as well as DFT calculations, and the influence of the A-elements on the magnetic properties is discussed.

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Introduction

MAX phases are a large and growing family of nanolaminate ternary carbides, nitrides, and carbonitrides with metallic and ceramic properties. Following the synthesis of around 60 members of what Jeitschko, Nowotny, and coworkers initially called H-phases in the 1960s,^1,2 extensive research on MAX phases began in 1996 when Barsoum et al. investigated the distinctive properties of phase-pure Ti₃SiC₂.³ They coined the name MAX phases and recognized that these materials exhibit a range of characteristics such as high-temperature stability, excellent thermal shock resistance, good mechanical properties, easy machinability, electrical/thermal conductivity, and even superconductivity^4–6 stemming from their inherent crystal structure. Represented by the general formula M_n+1AX_n, where M is an early-to-mid transition metal, A is a main group element, typically of groups 13–15, or a late transition metal, X is mainly C and/or N (B in rare cases), and n ranges from 1 to 4.^7,8 They possess a layered structure with edge-sharing M₆X octahedra interleaved by a single sheet of A atoms, resulting in the hexagonal space group P6₃/mmc.⁹ Nowadays, over 120 ternary MAX parent phases exist, as summarized in the latest review article.⁷ It is important to note that among the 3d transition metals, Mn, Fe, Co, Ni, Cu, and Zn have been observed to occupy the A-site of ternary MAX phases.¹⁰ To date, no earlier 3d transition metals have been reported on the A-site in any type of MAX phase.

The synthesis of MAX phase solid solutions by mixing elements on M, A, and/or X sites extends beyond the traditional ternary compounds, leading to the discovery of over 150 additional MAX phases synthesized to date.⁷ Of these, more than 80 compounds are M-site solid solutions, the most common form of alloying. A-site solid solutions are less common, with approximately 46 known examples. The rest are X-site and a few multi-site solid solutions. The A-site elements, including Al, Si, Ga, and Sn, are commonly used in solid solutions, particularly in the MAX phase, where the M-site is occupied by Ti, Zr, Hf, V, Nb, or Ta.⁷ In 2013, Cabioch et al.¹¹ demonstrated the potential of using an A-site solid solution to tune the properties of MAX phases. They synthesized Cr₂(Al_1−xGe_x)C, exhibiting composition-dependent thermal expansion coefficients. This discovery demonstrated how alloying or replacing specific elements at the A-site can influence key materials properties of MAX phases. Li et al.¹⁰ further demonstrated that A-site alloying can influence the magnetic properties by synthesizing V₂(A_xSn_1−x)C (A = Fe, Co, Ni, and/or Mn). By combining Sn with up to four different A elements, they synthesized V₂(Sn_0.67Mn_xFe_yCo_zNi_w)C and observed distinct ferromagnetic behavior, tunable through the multielement A-layer composition. Developing A-site solid solution MAX phases is crucial in this field of research, as it can open pathways to novel electronic and magnetic properties. Mn has been shown to form in A-site solid solution MAX phases, such as V₂(Sn_0.67Mn_0.33)C and related compounds.¹⁰ However, Li and colleagues limited the investigation to Mn, Fe, Co, and Ni as A-site elements. Early 3d metals such as Ti, V and Cr have never been considered as viable A-site substitution elements.

Nb₂SnC, a MAX phase with the same structure as the aforementioned V₂SnC, has garnered attention due to its bulk type-II superconductive behavior.⁶ It was first synthesized by Jeitschko and Nowotny in the 1960s and 1970s using the powder metallurgy route.^12,13 Notably, superconductivity has only been observed in ten MAX phases, with the highest critical temperature (T_c) reported at 10 K.^14,15 These include Ti₂InC,¹⁶ Ti₂InN,¹⁷ Ti₂GeC,¹⁸ Lu₂SnC,¹⁹ Nb₂SC,²⁰ Nb₂SnC,⁶ Nb₂AsC,²¹ Nb₂InC,²² Nb₂GeC,²³ and Mo₂GaC.²⁴ Among these superconducting MAX phases, five have Nb as an M-site element, and two have Sn as an A-site element.¹⁴ Nb₂SnC's superconductivity and stability are attributed to a local minimum in the density of states at the Fermi level, serving as a barrier for electrons just below the Fermi level.^6,15 Mechanically, Nb₂SnC is predicted to be ductile and damage tolerant, with Poisson's ratio between 0.1 (purely covalent) and 0.33 (purely metallic).²⁵ Its Poisson's ratio is anisotropic in the xz and yz planes but isotropic in the xy plane. This contrasts with Ti₂SnC, whose Poisson's ratio remains nearly isotropic in all directions.²⁵ Additionally, Nb₂SnC has shown promise for applications in supercapacitors and lithium-ion batteries, exhibiting an increase in discharge capacity over extended cycling.^5,26

Nb₂SnC's inherent superconductivity, machinability, and the potential for property tuning using the methods demonstrated by Li et al.,¹⁰ make it an attractive material to study. To date, the only reported A-site solid solutions of Nb₂SnC are Nb₂(Rh_0.2Sn_0.4Al_0.4)C and Nb₂(Pd_0.5Sn_0.5)C.²⁷ However, that study did not provide a detailed investigation of how A-site alloying influences the properties of the MAX phase.

In this study, buliding upon the work of Li et al.,¹⁰ we report the screening, synthesis, and comprehensive characterization of a new series of A-site solid solutions of Nb₂(Sn_1−xA_x)C where A represents the 3d transition metals Ti, V, Cr, Mn, Fe, Co, Ni, Cu, and Zn, with x = 0, 0.1, 0.2, 0.3, and 0.4. This is the first time traditional M-site elements such as V and Cr are introduced at the A-site. While prior work focused primarily on A-site substitutions in V₂SnC, the systematic investigation of 3d transition metal substitution in the Nb₂SnC MAX phase has not previously been reported. The comparative approach enables direct evaluation of how the nature of the A-site 3d metal influences the emergence and tunability of magnetism in Nb-based MAX phase solid solutions. The unknown A-site solid solution MAX phases were synthesized in sealed quartz ampules, and their structures were characterized using synchrotron-based powder X-ray diffraction (XRD), scanning electron microscopy (SEM), high-resolution transmission electron microscopy (HRTEM), secondary ion mass spectrometry (SIMS), and hard X-ray photoelectron spectroscopy (HAXPES). Phase stability was assessed through first-principles calculations, and magnetic properties were evaluated using vibrating sample magnetometry (VSM). To the best of our knowledge, this work represents the first comprehensive study of a systematic A-site substitution in Nb₂SnC MAX phases and provides insights into the interplay between A-site compositions, structure and magnetism in this class of materials.

Experimental

Preparation of Nb₂(Sn_1−xA_x)C

The target MAX phases were prepared using powdered elemental metals, with amounts targeting approx. 1.33 g of the desired MAX phase product. Experimental details can be found in the ESI,† Table SI-4. Briefly, the precursor powders were mixed in the desired ratios of Nb₂(Sn_1−xA_x)C, where x = 0, 0.1, 0.2, 0.3, and 0.4, and A represents Ti, V, Cr, Mn, Fe, Ni, Co, Cu, and Zn, using an agate mortar and pestle. Notably, for the parent MAX phase, Nb₂SnC, 10% excess Sn accounted for the loss due to evaporation. The mixed powders were pelletized using a 13 mm diameter cylindrical die and a manual hydraulic press at 10 tons. The obtained pellets were transferred to a fused silica ampule and sealed in a vacuum outside the glovebox on a Schlenk line. The ampule containing the parent MAX phase, Nb₂SnC, was heated to 1200 °C for 48 hours with a ramp rate of 5 °C per minute in a chamber furnace (Carbolite Gero CWF 1600). The precursor mixtures for the A-site solid solution MAX phases were heated to 1150 °C for 24 hours with a ramp rate of 5 °C per minute in a separate chamber furnace (Carbolite Gero CWF 1200). All samples were allowed to cool passively to the furnace's ambient temperature of 60–70 °C. The product pellets were homogenized in air using a tungsten carbide mortar and pestle.

Acid washing of Nb₂(Sn_1−xA_x)C

To remove unwanted side phases, such as niobium carbide and niobium oxide, the powdered samples were acid-washed by stirring them gently overnight in 20 mL of glacial hydrochloric acid. The samples were then thoroughly washed and double-filtered in a Büchner funnel with 500 mL of deionized water. The washed powder was dried on filter paper in a 100 °C drying furnace (Carbolite AX 60). The black powders were collected, weighed, and used for further analysis. Please note that while the side phases were removed entirely for some samples and reduced for others, acid washing was done on all samples to keep the synthesis process consistent.

Characterization. X-ray powder diffraction data were obtained from the high-energy beamline ID31 at the European Synchrotron Radiation Facility (ESRF) operated by Momentum Transfer. Powdered samples were carefully loaded into thin cylindrical slots (approx. 1 mm thickness), sandwiched between Kapton windows in a specialized high-throughput holder. The transmission measurements were conducted using high-energy X-rays at 75.00 keV, corresponding to a wavelength of 0.1653 Å. A Pilatus CdTe 2 M detector (1679 × 1475 pixels, 172 × 172 μm² each) was used to collect the diffracted intensities. The sample-to-detector distance was approximately 1.5 m for the high-resolution measurements. Background measurements from empty windows were subtracted to ensure data accuracy, and the NIST SRM 660b (LaB₆) reference sample was utilized for precise geometry calibration. The data processing pipeline included pyFAI software for calibration and subsequent image integration, incorporating corrections for flat-field, geometry, solid-angle, and polarization effects, thus yielding comprehensive and refined diffraction profiles. Rietveld refinement was performed using TOPAS V6 (Bruker).^28,29

SEM images were taken at the Eyring Materials Center at Arizona State University, using a Helios 5 UX (ThermoScientific) with an acceleration voltage of 25 kV, adapted with an Octane Elect detector (EDAX) for collecting EDS data. The EDS data were evaluated using the software EDAX Teams V4.1.

For HRTEM, tiny amounts of powder were dispersed in ethanol by ultrasound. A short waiting time afterward allowed larger grains to precipitate, and then 2 × 5 μL of this dispersion was drop-cast on a C-coated Cu grid for HRTEM, placed on a lab tissue. Bright-field (BF) HRTEM and scanning TEM (STEM) images were acquired with a JEOL 2200FS transmission electron microscope at an acceleration voltage of 200 kV using a 2k × 2k GATAN UltraScan1000XP CCD camera. The local chemical composition was determined using EDS in STEM mode with an Oxford windowless 80 mm² SDD X-MaxN 80 TLE detector with a 0.21 sr solid angle. HRTEM and EDX data were analyzed using Gatan Micrograph Suite and Oxford's Aztec software.

The magnetic properties of all samples were studied by vibrating sample magnetometry (VSM) using a PPMS DynaCool system (Quantum Design). Powder samples (10–30 mg) were weighed and placed into standard capsules. The measurements were taken within the field range of ±9 T at various temperatures ranging from 5 K to 400 K and normalized to the total sample mass.

All SIMS measurements were carried out using a CAMECA IMS SC Ultra instrument with cesium as the primary ion source. Specific details of the measurement procedure were described in previous publications.^30,31 Briefly, to achieve atomic-scale depth resolution, several adjustments to the measurement protocol were introduced, including high-angle ion incidence (75°), ultra-low impact energy (100 eV), in situ ion polishing, refined extraction settings, super cycle operation, and precise beam alignment. Deconvolution and calibration routines were then employed to enable quantitative analysis and determine the precise composition of each atomic layer with an accuracy of ±1%. Measurements with atomic depth resolution were obtained after sputtering ∼200 nm of the material first.

Synchrotron-based hard X-ray photoelectron spectroscopy (HAXPES) was conducted on the Nb₂(Sn_1−xV_x)C samples. Experiments were performed at beamline P22 at PETRAIII, German Electron Synchrotron DESY in Hamburg, Germany.³² A photon energy of 6 keV was used for all experiments, with the energy selected using a Si(111) double-crystal monochromator. A Phoibos 225HV analyzer (SPECS, Berlin, Germany) was used with a small area lens mode and a slit size of 3 mm. Spectra were collected using a pass energy of 30 eV. The total energy resolution in this setup was determined to be 328 meV, the 16/84% Fermi edge (E_F) width of a polycrystalline gold foil, and all spectra were aligned to the Au E_F. All samples were mounted as received on conductive carbon tape. The CasaXPS software package was used to determine the spectral areas for data normalization.

Computational details

All calculations were performed within the framework of density functional theory (DFT) as implemented in Vienna Ab initio Software Package (VASP) version 5.4.4^33–35 with the projector augmented wave (PAW) method^36–38 and the planewave basis set expanded to a kinetic energy cutoff of 520 eV. The Perdew–Burke–Ernzerhof (PBE) functional³⁸ was used to describe the electron exchange–correlation effects. We used a Γ-centered k-point sampling, with a density of 0.1 Å⁻¹.³⁹ All calculations were spin-polarized with an initial magnetic moment of three for atoms at the M- and A-sites in a ferromagnetic (FM) spin configuration. These settings ensure compatibility with data from the Materials Project database (v2021.05.13).⁴⁰ Structures were fully relaxed in terms of volume, shape, and atomic positions. The convergence criterion for self-consistency of relaxed structures is an energy convergence of 10⁻⁷ eV per atom and a force convergence of 0.001 eV Å⁻¹.

The special quasi-random structures (SQS) method⁴¹ as implemented in the Alloy Theoretic Automated Toolkit (ATAT) package,⁴² was used to generate representative supercell structures that approximate a solid solution alloy of Sn and the incorporated 3d transition metal on the A-sites.

The thermodynamic stability of Nb₂(Sn_1−xA_x)C phases was investigated with respect to decomposition into any combination of competing phases. The set of most competing phases was identified using a linear optimization procedure based on the simplex method under the constraint of a fixed MAX stoichiometry.^43,44 The stability is quantified in terms of formation enthalpy ΔH_cp by comparing its energy to the energy of the equilibrium simplex,

ΔH_cp = E(Nb₂(Sn_1−xA_x)C) − E(competing phases). (1)

where ΔH_cp < 0 indicates a stable phase, while ΔH > 0 is not stable or, at best, metastable. The selection of competing phases includes those found in the Materials Project database,⁴⁰ ternary M_n+1AX_n phases (n = 1–3),⁷ M_n+1AX_n phases with in-plane order (n = 1), out-of-plane order (n = 2 and 3) and solid solution disorder (n = 1, 2, and 3).^45,46

However, since we are investigating a solid solution of Sn and another element on the A-site, approximated through modeled disorder (SQS), the contribution from configurational entropy to the Gibbs free energy of formation, G, is approximated using

G = H − TΔS, (2)

where T is the temperature and ΔS is the configurational entropic contribution per A-site, assuming an ideal solution of Sn and A on the A-site, as given by

ΔS = −k_B[x ln(x) + (1 − x)ln(1 − x)], (3)

where x is the concentration of A.

Contributions from other temperature-dependent effects, e.g., lattice vibrations and electronic entropy, to the formation enthalpy are approximated to be negligible, as these tend to be canceled out in the Gibbs free energy of formation term.⁴⁷ This approach has been proven to work exceptionally well for previous theoretical studies of both ternary and quaternary MAX phases.^7,45,46

Results and discussion

The results presented in the following sections elucidate the structural and compositional analysis of solid solutions of Nb₂(Sn_1−xA_x)C derived from synchrotron powder X-ray diffraction, SEM, EDS, HAXPES, HRTEM, and SIMS. Additionally, magnetic properties are discussed.

Structural analysis

Based on synchrotron powder X-ray diffraction (XRD) data (Fig. 1(a)) the parent MAX phase Nb₂SnC was synthesized in high quality with only a minor side phase of NbO. Rietveld refinements of the XRD data deliver the lattice parameters of a = 3.227 Å and c = 13.874 Å for Nb₂SnC, which are in excellent agreement with published data, a = 3.2408 Å and c = 13.802 Å.^6,15,48 The micromorphology of particles with terraces is typical of the layered structure of MAX phases (Fig. 1(b)). The soft edges can be attributed to the acid washing in glacial hydrochloric acid.

Fig. 1 (a) Rietveld refinement (orange line) and residuum curves (grey line) of the synchrotron powder X-ray diffraction data (black dots) of Nb₂SnC. (b) SEM electron micrograph showing the morphology of the MAX phase Nb₂SnC.

After confirmation of the successful synthesis of the parent MAX phase, solid solutions with nine different A elements were synthesized. The A elements span the 3d transition metals from Ti to Zn, and various amounts of these A elements were nominally incorporated into the parent phase Nb₂SnC. This led to a total of 36 samples with the general composition Nb₂(Sn_1−xA_x)C with A = Ti, V, Cr, Mn, Fe, Co, Ni, Cu, and Zn, and x = 0.1, 0.2, 0.3, and 0.4. Higher x values resulted in excessive amounts of side phases.

Please note that we chose one of the solid solution phases to be discussed in more detail here, while the others are discussed as a group. The earliest transition metals, like Ti and V, have never been reported to occupy the A-site of a MAX phase structure. Here, we find that only small amounts of Ti substitute Sn in Nb₂SnC, while surprisingly large quantities of V do, under identical synthesis conditions. For this reason, we will discuss the solid solution phases with V on the A-site in greater detail as it challenges existing paradigms about transition metal solubility at A-sites in MAX phases. The data of the solid solutions with all other 3d metals are compiled in Tables SI-5–14 and Fig. SI-1–9 (ESI†). Rietveld refinements of the synchrotron powder X-ray diffraction data of the solid solution series Nb₂Sn_1−xV_xC with x = 0.1, …, 0.4 (Fig. 2) show that all products contain multiple phases, with the MAX phase being the primary phase. Side phases are different niobium-containing binary compounds, such as niobium oxide and carbide(s).

Fig. 2 Results of Rietveld refinements (orange lines) and residuum curves (grey) of the synchrotron powder X-ray diffraction data (black dots) of the solid-solution Nb₂(Sn_1−xV_x)C. (a) Nb₂Sn_0.9V_0.1C, (b) Nb₂Sn_0.8V_0.2C, (c) Nb₂Sn_0.7V_0.3C and (d) Nb₂Sn_0.6V_0.4C.

In general, the amount of side phases increases with the nominal V amount, and the MAX phase purity decreases in weight percent from 95% (x = 0.1) to 92% (x = 0.2), 79% (x = 0.3), and eventually 64% (x = 0.4). The trend suggests a limit of solid solubility of vanadium in Nb₂SnC under the current synthesis conditions. This can be attributed to the increased lattice strain or formation of competing side phases at higher vanadium concentrations, especially Nb₃Sn and Nb₆C₅. This should be considered in the following integral methods, such as HAXPES and magnetometry, as we show below. Aside from the composition of the products, analysis of the XRD data also allows the determination of the lattice parameters of the solid solution series. The lattice parameters change upon the incorporation of V on the Sn-site (Fig. 3(a)). Although the change is not linear, the c-lattice parameter increases while the a-lattice parameter decreases, leading to a reduction in cell volume with higher x values (Fig. 3(a)).^49,50 This shows that incorporating V on the A-site causes the crystal structure to contract laterally while expanding along the c-axis due to differences in metallic radii and bonding characteristics between Nb and V. It is important to note that this trend is decoupled from the existence of side phases in the product because only the peaks that belong to the MAX phase structure are analyzed. The trend observed for the a-lattice and c-lattice parameters is consistent with the A-site substitution of MAX phases, as it is typically associated with changes in the A-layer bonding.⁴ It is also confirmed with calculated cell volume and lattice parameters, as shown in Fig. SI-32 (ESI†).

Fig. 3 (a) Change in lattice parameters and volume with increasing x in Nb₂Sn_1−xV_xC. (b) Correlation between nominal V amount and actual V amount based on multiple EDS spot measurements. (c) EDS spectrum of the Nb₂(Sn_0.7V_0.3)C MAX phase. EDS measurement error ±2.5%.

Unlike the anisotropic change of lattice parameters observed here, M-site substitution typically changes both lattice parameters in the same direction.⁵¹ Synchrotron X-ray powder diffraction data also confirm the behavior. Fig. SI-9 (ESI†) shows the (006) peak shift to lower angles as we increase the content of the 3d element in Nb₂SnC, which is typically associated with changes in A-layer bonding and spacing.⁵⁰ In the case of Nb₂Sn_1−xV_xC series, the (006) peak shifts from 2θ = 4.098° for Nb₂SnC to 2θ = 4.047° for x = 0.4. This is expected as the V has a smaller metallic radius (134 pm) than Sn (140 pm). The same trend is observed for all 3d metals (except for Ti), as studied here (full data analysis is presented in the ESI,† Tables SI-5–14 and Fig. SI-1–9).

SEM/EDS was conducted on the solid solution series Nb₂Sn_1−xV_xC with x = 0.1, …, 0.4 to further confirm the V content and distribution within the samples. Ten random spots were measured per sample to guarantee sufficient statistics for each composition. We find a good correlation between the nominal amount of V and the measured amount (by EDS), with the most significant error bar for the sample with x = 0.2 (Fig. 3(b)). Fig. 3(c) shows the EDS spectrum of the Nb₂(Sn_0.7V_0.3)C sample (as an example), indicating 26%, close to the nominal 30% V content. The same data were collected for the solid solutions with all other 3d metals (ESI,† Fig. SI-19–26), which are discussed in greater detail below.

HAXPES was used further to investigate the chemical nature of the Nb₂Sn_1−xV_xC series. Survey spectra (Fig. SI-30, ESI†) show all expected elements, including oxygen from partial oxidation, as discussed above. Core-level spectra (shown in Fig. 4) were collected and show a complex picture of chemical states, particularly for Nb. The Nb 3d core level (Fig. 4(a)) shows four distinguishable chemical state contributions. The MAX phase results in a sharp doublet with the Nb 3d_5/2 peak (Nb_MAX) at 203.3 eV with a spin–orbit–split (SOS) of 2.75 eV. A broader, intense doublet is found with its Nb 3d_5/2 peak at 207.6 eV, commensurate with Nb₂O₅. Two additional, lower-intensity features are visible on either side of the main Nb_MAX 3d_5/2 feature: a peak at a lower binding energy (BE) of 202.4 eV and a peak at a higher BE of 204.3 eV. Due to the lack of literature references and the difficulty in obtaining reliable references, the origin of these peaks can only be inferred through qualitative arguments based on general trends in chemical state shifts and the information gained from XRD. The lower BE peak most likely stems from the NbC and Nb₃Sn phases, whilst the higher BE peak could originate from the NbO phase observed.

Fig. 4 HAXPES core level spectra of the Nb₂(Sn_1−xV_x)C sample series, including (a) Nb 3d, (b) Sn 3d, (c) C 1s, and (d) V 2p. Spectra for each sample are normalized to their respective total Nb 3p_3/2 area after subtracting a Shirley-type background. The brackets in (a) indicate the doublet peaks associated with Nb in the MAX phase (Nb_MAX) and Nb₂O₅.

Qualitatively, the trends in relative intensities of these features relative to the main Nb_MAX peak agree with the ratios found from XRD (although, of course, it should be noted that the probing depth of the two techniques differs greatly). The Nb 3p core level was also collected for completeness (Fig. SI-31(a), ESI†). Due to its larger lifetime broadening, it only resolves a Nb_MAX (BE(3d_5/2) = 361.8 eV) and a Nb₂O₅ (BE(3d_5/2) = 365.8 eV) feature (Fig. SI-30, ESI†). The Sn 3d core level spectra (Fig. 4(b)) show two doublets, one from the MAX phase (Sn 3d_5/2 at 485 eV) and a second from Sn oxide (Sn 3d_5/2 at 486.9 eV). Again, the feature arising from the MAX phase is considerably sharper than that arising from the oxide. The C 1s spectra (Fig. 4(c)) show a clear contribution from the MAX phase (C–M) at 283.1 ± 0.1 eV. In addition, adventitious C⁰ produces a strong feature at 284.7 eV for all samples. Finally, only very weak signatures of the V 2p core level could be detected in the HAXPES experiments (Fig. 4(d)), with its analysis further hampered by its position between the strong Sn 3d and O 1s core lines (Fig. SI-31(b), ESI†). The V 2p_3/2 peak appears at approx. 512.9 eV and a faint signature from the 2p_1/2 peak can be seen for the higher nominal amount V samples at approx. 520.3 eV. These BE positions and the SOS value are commensurate with V in oxidation state zero.

The low relative intensity of the V core level cannot be explained by photoionization cross sections alone. Whilst the photoionization cross section for V 2p_3/2 is approximately half that of Sn 3d_5/2,^52–54 it is clear that even for the samples with higher nominal V concentrations, the V 2p_3/2 signal intensity does not reach levels anywhere near expected. This is in stark contrast to the XRD and EDS values reported, which suggests a considerable difference in elemental distribution at the sample surface in comparison to the bulk, as further confirmed by SIMS depth profiling (Fig. SI-29, ESI†) performed from the sample surface in conventional mode (without atomic resolution), which reveals that the first ∼20 nm contain significant amounts of oxygen, tin, niobium, and carbon, but no detectable vanadium. Additionally, the signals are noisy, suggesting an amorphous and disordered structure. Only deeper into the sample do the signals stabilize, and vanadium becomes detectable, indicating that vanadium's surface layer is strongly oxidized and depleted.

As mentioned above, V is the most unusual and unexpected element to be incorporated on the A-site of Nb₂SnC in surprisingly high amounts. Therefore, we collected HRTEM, STEM/EDS, and SIMS data on the Nb₂Sn_1−xV_xC sample with ca. 30% V on the Sn-site. Fig. 5(a) presents a bright-field HRTEM image of the surface of a micron-sized grain (see inset of Fig. 5(a) for the whole grain), showing the layered structure where the grain is thin enough to undergo TEM. The red-squared area is directionally FFT-filtered perpendicularly to the visible layers, and the result is displayed in Fig. 5(a). The distance of 10 bright layers, i.e., presumably Nb layers, is 7.06 nm and is equivalent to five times the c-axis lattice parameter. TEM delivers c = 14.12 Å, which deviates from the more precise XRD determination in Fig. 3(a) by less than 2%. STEM/EDS of the identical grain is shown in Fig. 5(b), which includes an electron image and elemental distribution maps of Nb, V, and Sn. The composition of this grain is Nb₂Sn_0.83V_0.17C, as determined from the corresponding EDX spectrum in ESI,† Fig. SI-27. EDS of another grain delivered Nb₂Sn_0.8V_0.2C, as shown in ESI,† Fig. SI-28. It is striking that the M₂A₁X₁ structure suggests that V occupies A-sites. Compared to SEM/EDS analysis of larger sample volumes, the lowered V content in the MAX phase is ascribed to grain-to-grain variations or some V not incorporated in the MAX phase.

Fig. 5 (a) Bright-field HRTEM image of a nominally Nb₂Sn_0.7V_0.3C MAX phase grain with a zoomed-in view of the directional FFT-filtered image of the red square with c-axis periodicity. The inset shows the whole grain. (b) STEM/EDS electron image of grain, Nb, V, and Sn elemental distribution. (c) Layer stacking is measured using SIMS. Nb is found elementally pure in M-layers, while Sn and V form the A-layer, and C, enriched with about 9% O, coexists in X-layers.

The SIMS technique used in this study, when operated under optimized conditions, provides accurate atomic-level resolution. In this mode, the composition of each atomic layer is determined independently, eliminating any mixing of signals between layers. Consequently, the elemental concentrations can be extracted with high confidence, enabling precise quantification of the chemical composition of individual layers (Fig. 5(c)). The M-type layer was found to consist solely of Nb. The X-type layers are composed of carbon atoms, with an additional oxygen contribution of about 9%. This aligns with previous findings that suggest a strong tendency for oxygen incorporation into the X layer.^30,31 The most intriguing results were obtained from the A-type layer, where Sn accounts for 70% of the composition, while the remaining 30% consists of V. This provides direct and clear evidence that vanadium has been incorporated into the A layer and is entirely absent from the M layers. Both studies of the nominally Nb₂Sn_0.7V_0.3C sample, HRTEM, and particularly the SIMS profile, unequivocally prove that V is indeed sitting on the A-site and – not even partially – on the M-site, which would appear very likely because the ternary MAX phase V₂SnC also exists.⁵⁵

To visualize a broader picture along the 3d series from Ti to Zn, the correlation between the nominal and experimentally determined amount from SEM/EDS is presented in Fig. 6 (bubble size in Fig. 6) is shown for all 3d elements. As the synthesized MAX phases are treated with glacial hydrochloric acid, the actual amount of incorporated A-element also depends on the stability of the A-element against the acid. Therefore, we find a complex relationship between nominal and measured concentrations for various transition metals (Ti, V, Cr, Mn, Fe, Co, Ni, Cu, and Zn). A detailed table of the relative occupancy deviation from synchrotron powder X-ray diffraction and average A amount from EDS is shown in ESI,† Tables SI-5–14.

Fig. 6 Correlation between average A element amounts determined by EDS analysis (bubble size) and nominal A element amount in Nb₂(Sn_1−xA_x)C for all investigated A elements. The bubble sizes represent the average A element amounts, with the smallest bubble corresponding to 0 at% and the largest bubble corresponding to 40 at% as measured by EDS. EDS measurement error ±2.5%.

Interestingly, the non-neighbors V and Zn demonstrate close agreement between the nominal and measured concentrations, suggesting superior stability and incorporation independent of their position in the 3d metal series. Most elements exhibit an overall increasing trend with nominal concentration, albeit non-linearly. This stems from differential reactivity with hydrochloric acid and the formation of potentially stable secondary phases. This is especially true for Cr, Mn, and Fe, as they generally have lower-than-intended levels of incorporation in the structure. On the other hand, Cu aligns well with the trend at lower concentrations but deviates significantly at x = 0.4, suggesting a possible solubility limit below 40% atomic concentration. This agreement between average values from EDS and occupancy from refinements highlights the need for accurate post-synthesis characterization of such a complex materials system. The comprehensive analysis combining synchrotron X-ray powder diffraction, EDS in SEM and TEM, and SIMS confirms the successful incorporation of A elements in the Nb₂(Sn_1−xA_x)C solid solution MAX phase. At the same time, the EDS results reveal the complex relationship between nominal and actual compositions.

Stability predictions

To gain further insights into the extent and feasibility of incorporating 3d transition metals into the A-site of Nb₂SnC, DFT calculations were performed. Thermodynamic stability was evaluated for Nb₂(Sn_1−xA_x)C with x = 0, 1/6, 1/3, and 0.5, where A spans the 3d transition metals from Ti to Zn. Configurational entropy contributions were accounted for T > 0 K, for both Nb₂(Sn_1−xA_x)C and considered competing phases with solid solution disordered configurations. Note that ternary M₂SnC MAX phases, including Nb₂SnC, were used as competing phases, as the goal is to evaluate the impact of A-site alloying. This ensures a reference stability of zero at x = 0, even though Nb₂SnC alone is stable at −43 meV per atom relative to NbSn₂, Nb₆C₅, and Nb₂C.

Fig. 7 illustrates the stability (ΔG_cp) for Nb₂(Sn_1−xA_x)C evaluated at temperatures 0 to 2000 K. Five of the nine studied systems, that is, A = Mn, Fe, Ni, Cu, and Zn, are thermodynamically stable (ΔG_cp < 0) at various T and x. Notably, Zn is stable across all x and T, while Mn, Fe, Ni, and Cu become stable upon increasing temperature due to the configurational entropy contribution from the A-site solid solution.

Fig. 7 Stability evaluated at different temperatures for Nb₂(Sn_1−xA_x)C as a function of x with A from Ti to Zn. Note that Nb₂SnC has been considered as a competing phase.

Two additional systems (A = V, and Co) approach stability at x ≤ 1/3 with ΔG_cp < +30 meV per atom at 1500 K. Increasing the temperature to 2000 K further stabilizes these systems, reducing ΔG_cp < +25 meV per atom for x ≤ 1/3. Such small positive values may indicate metastability, meaning that while the system is not the lowest in energy, it does not imply that synthesis of the phase can be ruled out, and synthesis may still be feasible. It should be noted that Nb₂(Sn_1−xA_x)C have been modeled assuming ideal occupancy of the M-, A-, and X-sites without consideration of vacancies or defects, which are known to sometimes enhance phase stability.

An interesting trend in Fig. 7 is that all phases, except for A = Ti, become more stable with increasing temperature, as shown by decreasing ΔG_cp. This behavior can be traced back to the set of most competing phases, see ESI,† Table SI-16, where disordered MAX phases, like (Nb_0.67Ti_0.33)₂SnC and (Ti_0.67Nb_0.33)₃SnC₂, are consistently among the set of most competing phases for all values of x and T. Since both Nb₂(Sn_1−xTi_x)C and competing phases (Nb_0.67Ti_0.33)₂SnC and (Ti_0.67Nb_0.33)₃SnC₂ exhibit configurational entropy from disorder on either A- or M-site, the entropic stabilization cancels out, leading to minimal temperature dependence for A = Ti. In contrast, for A = V to Zn, similar competing MAX phases (e.g., (Nb_0.67A_0.33)₂SnC or (A_0.67Nb_0.33)₃SnC₂) exist but are not identified as the most competitive ones, enabling the entropy from A-site disorder to effectively stabilize Nb₂(Sn_1−xA_x)C with increasing temperature. This may explain why our approach does not work for incorporating Ti on the A-site.

Magnetic analysis

As a first step towards the application of the new Nb₂(Sn_1−xA_x)C compounds, we screened their magnetic response. In this way, cooperative effects like ferromagnetism^10,56 or superconductivity⁶ can be identified. Most MAX phases, however, are Pauli paramagnets, and element substitution can significantly increase the density of states at the Fermi energy.⁵⁷ Since magnetometry delivers integral data of the samples, special attention must be paid to the contribution of side phases. The XRD refinement and SEM-EDS analysis revealed MAX phase purities of the V series of 95% for x = 0.1 and decreasing purity toward 64% for x = 0.4. So, we expect the most significant results for the samples at a lower substitution level. In all samples of the V series, the side phases Nb₆C₅, NbO, and Nb₃Sn have been identified. While the first two compounds are paramagnetic, it is well known that the Nb₃Sn phase is a superconductor with T_C = 18 K and paramagnetic above.

Fig. 8 presents the field-dependent magnetization of the V series at 300 K. Overall, all samples show a linear response at high fields, indicating paramagnetism. The paramagnetic slopes increase with the V substitution level. The Nb₂Sn_0.7V_0.3C MAX phase additionally has a tiny ferromagnetic contribution of 0.2 × 10⁻² A m² kg⁻¹. Such behavior has previously been observed for the NbC phase with M = 6.0 × 10⁻² A m² kg⁻¹.⁵⁸ Scaling this data to our results suggests that about 3.5 mass% of the sample is NbC, which matches well the result of the XRD refinement (3.1 mass%). Note that the authors⁵⁸ ascribed this feature to a faint Fe contamination from preparation, which we cannot exclude here.

Fig. 8 Magnetization of the Nb₂(Sn_1−xV_x)C series with x = 0.1–0.4 at T = 300 K. The inset shows the temperature-dependent magnetization of the x = 0.3 sample in B = 10 mT after zero-field cooling. T_C denotes the critical temperature.

More interesting is the behavior at low temperatures. The inset in Fig. 8 presents the temperature-dependent magnetization of the Nb₂Sn_0.7V_0.3C MAX phase at B = 10 mT in the interval 2–30 K. We obtain the typical diamagnetic behavior of a superconductor below T_C = 13 K with a rather broad transition region and paramagnetism above. All other samples of the V series also show superconducting response with T_C = 12–14 K at strongly varying negative magnetization values at T = 2 K after cooling in zero field. This clearly points to a side phase producing this feature. Although the critical temperature T_C = 18 K of the XRD-identified Nb₃Sn phase is not reached, it is interesting to compare with non-MAX phase compounds with similar composition of elements like in the MAX phase, except C. NbSnV compounds were investigated in the 1970s.⁵⁹ The researchers found that T_C varies strongly with composition. Close to the MAX phase stoichiometry, omitting C, e.g., in Nb_0.625Sn_0.25V_0.125 or Nb_0.5Sn_0.25V_0.25, the ordering temperatures were determined to be T_C = 14.2 K and 9.8 K, respectively, while the A15 crystal structure of Nb₃Sn is retained. This explains both the XRD observations and magnetometry results for the superconducting side phase.

Overall, the magnetic response of the Nb₂(Sn_1−xV_x)C MAX phase series is paramagnetic, and V substitution increases the paramagnetic susceptibility. Observed deviations from this behavior can solely be ascribed to side phases. MAX phases with other elements behave similarly. Thus, the best choice of field and temperature for a magnetic screening of the entire Nb₂(Sn_1−xA_x)C series is a temperature far above possible superconducting transitions (e.g., 300 K) and high magnetic fields to gain solid data for the high-field susceptibility.

In Fig. 9, the results of the high-field susceptibility are presented for all A-site elements of the 3d series for different substitution levels. Here, we use the amounts of substitution achieved on A-sites as experimentally determined by SEM/EDS (Fig. 6) after chemical etching. Circle areas reflect the magnitude of the susceptibility, while blue color accounts for paramagnetic and red color diamagnetic responses. For individual substituting elements, χ has the tendency to rise as expected.

Fig. 9 Magnetic high-field susceptibility χ as a function of substituting element A in the Nb₂(Sn_1−xA_x)C MAX phases and EDX-measured substitution level at T = 300 K. The circle area reflects the magnitude of χ. Blue and red circles indicate paramagnetic and diamagnetic responses, respectively.

The susceptibility χ strongly increases with the rising number of 3d electrons up to Mn in the 3d series, while it continuously decreases towards Ni substitution, and for Cu and Zn, the samples become diamagnetic. More quantitatively, crossing the elements at a constant substitution level, e.g., around 10 at%, reveals a strongly non-linear behavior, as plotted in ESI,† Fig. SI-33. The magnetic susceptibility peaks for Mn at χ = 17 × 10⁻³ A m² kg⁻¹ T⁻¹, followed by a sharp drop over Fe to Co, before the late 3d elements exhibit diamagnetism. The paramagnetic base level is only χ = 0.3 × 10⁻³ A m² kg⁻¹ T⁻¹ for Ti and χ = 0.4 × 10⁻³ A m² kg⁻¹ T⁻¹ for Co. Such behavior across the 3d series suggests that the enrichment with electrons of the Nb₂(Sn_1−xA_x)C MAX phases is equivalent with one or more 3d-peak(s) in the electronic density of states successively shifting through the Fermi energy as we observed before for M-site substitution in the Pauli paramagnetic (V_1−xCr_x)₂GaC system.⁵⁷ Moreover, the range of paramagnetic susceptibilities matches well with the varieties in ternary MAX phase systems like V₂AlC (χ = 0.14 × 10⁻³ A m² kg⁻¹ T⁻¹) and Cr₂GaC (11.1 × 10⁻³ A m² kg⁻¹ T⁻¹).^60,61

Conclusions

In this study, we investigated a series of new A-site solid solutions of MAX phases based on Nb₂SnC with partial substitution of Sn by 3d transition metals. Powder samples with nominal chemical compositions Nb₂(Sn_1−xA_x)C (A = V, Cr, Mn, Fe, Co, Ni, Cu, and Zn; x = 0.1 to 0.4) were synthesized by high-temperature solid-state reactions.

Near single-phase products were obtained at x = 0.1, while increasing substitution levels led to the appearance of secondary phases, primarily niobium oxide and carbide. Interestingly, Ti is not incorporated at the A-site to the same extent as the other 3d metals do, a result corroborated by DFT stability calculations across different temperatures. A key finding is the successful substitution of Sn by V and Cr up to 40 at% and 22 at%, respectively, representing the first known incorporation of these elements into the A-site of a MAX phase. Synchrotron powder X-ray diffraction, electron microscopy, and element-specific spectroscopy techniques thoroughly characterized the new solid solution phases. Substitution with 3d metals on the A-site led to systematic changes in lattice parameters, where the a-axis generally decreased while the c-axis increased, leading to an overall reduction in unit cell volume. The nominal 3d metal amount is compared to the actual amount determined by multiple elemental analyses. EDS point measurements showed that the incorporated amount of A-site 3d transition metal increases with nominal x, though the maximum achievable incorporation depends on the specific element. One of the earliest 3d metals, V, substitutes Sn in surprisingly high amounts, and therefore, the sample with nominal composition Nb₂Sn_0.7V_0.3C was selected for detailed study. While HAXPES data do not show a significant amount of Sn in the V-samples in the probed near-surface region, both TEM-EDS and SIMS analysis (both in deeper probing regions) confirm that vanadium is incorporated on the A-site, most likely with some grain-to-grain variations. Furthermore, magnetic measurements of all new solid solution phases showed that high-field susceptibility (χ) strongly increases with the number of 3d electrons up to Mn, then continuously decreases towards Ni incorporation, and eventually becomes diamagnetic for Cu and Zn. Despite the reduced phase purity at higher substitution levels, the successful synthesis of these novel A-site solid solution MAX phases demonstrates a viable strategy that opens new avenues for tuning the structure and properties, thus offering exciting opportunities for future applications.

Author contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.

Conflicts of interest

There are no conflicts to declare.

Data availability

All data supporting this article have been included as part of the ESI.†

Acknowledgements

This work has been supported by the Deutsche Forschungsgemeinschaft (DFG, German Research Foundation) within CRC/TRR 270, projects B03 and B02 (Project-ID 405553726) and the Swedish Government Strategic Research Area in Materials Science on Advanced Functional Materials at Linköping University, Faculty Grant SFOMat-LiU 2009-00971. This material is based upon work supported by the National Science Foundation under Grant No. 2143982. We acknowledge the use of facilities within the Eyring Materials Center at Arizona State University, supported in part by NNCI-ECCS-1542160. We acknowledge the European Synchrotron Radiation Facility (ESRF) for providing synchrotron radiation facilities and Momentum Transfer to facilitate the measurements. Jakub Drnec is thanked for their assistance and support in using beamline ID31. The measurement setup was developed with funding from the European Union's Horizon 2020 research and innovation program under the STREAMLINE project (Grant Agreement ID: 870313). The computations were enabled by resources provided by the National Academic Infrastructure for Supercomputing in Sweden (NAISS) at the National Supercomputer Centre (NSC) and the PDC Center for High-Performance Computing, partially funded by the Swedish Research Council through grant agreement no. 2022-06725. We acknowledge DESY (Hamburg, Germany), a member of the Helmholtz Association HGF, for the provision of experimental facilities. Parts of this research were carried out at PETRA III using beamline P22. Beamtime was allocated for proposals I-20230848 and H-20010087. PPM was supported by the National Science Centre, Poland, within SONATA BIS 14 2024/54/E/ST11/00171 and the National Centre for Research and Development, Poland, within LIDER XII LIDER/8/0055/L-12/20/NCBR/2021 projects.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d5tc02347e

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