Multielemental single–atom-thick A layers in nanolaminated V₂(Sn, A) C (A = Fe, Co, Ni, Mn) for tailoring magnetic properties

Significance

M_n+1AX_n phases are a family of inherently nanolaminated ternary compounds with hexagonal crystal structure (space group P₆₃/mmc, 194). Here, M is vanadium element, and A is Fe, Co, Ni, Mn, or their binary/ternary/quaternary mixtures. Due to the elemental flexibility at A site, 15 nanolaminated V₂(A_xSn_1-x)C MAX phases are synthesized, including 1 high-entropy MAX phase that all Fe, Co, Ni, Mn, and Sn elements simultaneously occupied A site. Tailoring of individual single–atom-thick layers in nanolaminated MAX phases offers atomic-level control of material properties, such as their distinct magnetic behaviors. The alloying in 2-dimensional A layer of MAX phases provides a unique route to design their crystal structure and to discover unexploited properties, which would develop promising functional materials for microelectronic device.

Abstract

Tailoring of individual single–atom-thick layers in nanolaminated materials offers atomic-level control over material properties. Nonetheless, multielement alloying in individual atomic layers in nanolaminates is largely unexplored. Here, we report 15 inherently nanolaminated V₂(A_xSn_1-x)C (A = Fe, Co, Ni, Mn, and combinations thereof, with x ∼ 1/3) MAX phases synthesized by an alloy-guided reaction. The simultaneous occupancy of the 4 magnetic elements and Sn in the individual single–atom-thick A layers constitutes high-entropy MAX phase in which multielemental alloying exclusively occurs in the 2-dimensional (2D) A layers. V₂(A_xSn_1-x)C exhibit distinct ferromagnetic behavior that can be compositionally tailored from the multielement A-layer alloying. Density functional theory and phase diagram calculations are performed to understand the structure stability of these MAX phases. This 2D multielemental alloying approach provides a structural design route to discover nanolaminated materials and expand their chemical and physical properties. In fact, the magnetic behavior of these multielemental MAX phases shows strong dependency on the combination of various elements.

Tailoring individual single–atom-thick layers in nanolaminated materials offers atomic-level control over modifying a desired property of a material. For example, artificially nanolaminated magnetic materials are widely used in storage media and devices. Notably, the demonstration of the giant magnetoresistance (GMR) effect has enabled hard drives and other storage media (1). The GMR sensitivity is expected to be highest when only single-atom layers of ferromagnetic materials are sandwiched since alignment of magnetization vectors of neighboring ferromagnetic atoms has the lowest energy cost.

Conceptually, this type of structure with single–atom-thick layers can be correlated to M_n+1AX_n phases (or MAX phases), which are a family of inherently nanolaminated ternary compounds with hexagonal crystal structure (space group P₆₃/mmc, 194), where M is an early transition metal, A is mainly from A group elements, X is carbon and/or nitrogen, and n = 1 to 3 (2, 3). Their crystal structure can be depicted by alternative stacking of M_n+1X_n sublayer and a single atomic layer of A, usually referred to as 211, 312, and 413 phases according to the value of n. If alloying or replacement of ferromagnetic elements at specific crystal sites in MAX phases could be realized, their magnetic properties may also be tailored. The key notion here is that the A layer is just 1-atomic layer thick. However, there are few reports on magnetic MAX phases in which M site contains Mn or Fe, such as (Cr_0.75Mn_0.25)₂GeC (4), Mn₂GaC (5), (V,Mn)₃GaC₂ (6), and (Cr,Fe)₂AlC (7), etc.

Control of the occupancy of magnetic elements on the A rather than M sites would be crucial to tune magnetic properties. The finding of iron in the A site of Mo₂(GaAuFe)C (8) is encouraging; however, the presence of secondary iron-containing impurity phases in the Mo₂(GAuFe)C films impedes the determination of magnetic properties on the A plane. Moreover, theoretical predictions suggest that Ni and Co should tend to occupy the M sites of the MAX phases (9), but this has not been demonstrated experimentally, likely due to the thermodynamically preferred formation of competing binary MA alloys or intermetallic phases. Generally, the late transition elements of Fe, Co, Ni, and Mn have not been considered as possible A-site elements in the definition of MAX phases. It would be of great interest to introduce these magnetic elements at A sites since the overlap between electron clouds of M and A atoms is limited compared with that between M and X atoms, which should aid in preserving the magnetic properties. A possible approach to realize this is by alloying with other main group elements in such a way that the alloying transforms neither the atomic stacking in the crystal nor their bonding with the parent structure. Fe, Co, and Ni have been used as additives for the formation of M′₃SnC₂ (M′ = Ti, Zr, Hf) MAX phases (10), while none of the Fe/Co/Ni were detected in the final target MAX phases. The A_xSn_y (A = Fe, Co, Ni, Mn, or their binary mixtures) phases belong to the same space group (P₆₃/mmc or 194) as MAX phases. This similarity in crystal structures facilitates the nucleation of MAX phases in saturated alloys by a reaction between a binary carbide and an intermediate state of A_xSn_y alloys. During such a reaction, these complex atoms in molten (or solid) alloys may thermodynamically and coherently arrange with [M₆C] octahedral building blocks to form ternary-layer structure. Thus, it would be of great interest to obtain magnetic MAX phases by an alloy-guided reaction.

Here, we demonstrate this approach to A-site alloying of Sn with Fe/Co/Ni/Mn magnetic elements to synthesize a series of magnetic MAX phases of V₂(A_xSn_1-x)C (A = Fe, Co, Ni, Mn, or their binary/ternary/quaternary combination). The regular M₂AX crystal structure of V₂SnC is retained, with a mixture of 2 to 5 elements (Sn plus 1 or more of Fe, Co, Ni, and Mn) randomly occupying the A sites. The magnetic properties of V₂(A_xSn_1-x)C MAX phases exhibit distinct ferromagnetic behavior that can be tailored by their constitutive elements. These results have wide-ranging implications since they not only can be used for tailoring magnetic properties but most importantly, demonstrate the formation of multielement A layers as an analogy to high-entropy alloys (11, 12), 2-dimensional (2D) in the sense that alloying exclusively occurs in the single–atom-thick A layers, in layered transition metal carbides.

Results and Discussion

MAX Phases Containing 1 Magnetic Element.

X-ray diffraction (XRD) pattern of V-Fe-Sn-C system is close to the characteristic crystal structure of V₂AlC MAX phase (13), with characteristic peaks at 2θ ∼ 13°, 2θ ∼ 26°, and 2θ ∼ 40°, indicating formation of a 211 MAX phase (shown in Fig. 1A) and also, with some by-products of Sn metal, intermetallic compounds (FeSn₂), and nonstoichiometric vanadium carbide (VC_x) (SI Appendix, Fig. S1A). The micromorphology of observed particles with terraces is typical layered structure of MAX phases (Fig. 1B). The atomic force microscopy (AFM) measurement also shows the typical layer structure in nanoscale (SI Appendix, Fig. S2), and the height of terrace is about 3 nm, which is close to double length of unit cell of M₂AX in c direction. In previously reported M′₃SnC₂ (M′ = Ti, Zr, Hf) phases (10), Fe was absent in the final products. Here, however, the corresponding energy-dispersive X-ray spectroscopy (EDS) spectra of these particles showed the presence of Fe beside V, Sn, and C (Fig. 1C). The relative atomic ratio of V:(Fe + Sn) is ∼2:1, consistent with the stoichiometry of 211 MAX phases (SI Appendix, Table S1). The high content of Fe (9.4 at.%) and its uniform distribution in elemental mapping (SI Appendix, Fig. S3) further indicate that Fe is incorporated in the as synthesized MAX phase.

Fig. 1.

V₂(Fe_1/3Sn_2/3)C MAX phase: XRD pattern (A) and SEM image (B) after acid treatment. (C) EDS spectrum indicated that particles contain V, Sn, Fe, and C elements. High-resolution STEM images showing atomic positions along [11$\overline{2}$0] (D) and [1$\overline{1}$00] (E) directions, respectively. The corresponding high-resolution image together with EDS mapping (F) and line scanning (G) of V-Kα, Sn-Kα, and Fe-Kα signals are shown. (Scale bars: STEM images, 1 nm.) a.u., arbitrary unit; cps, counts per second.

To determine the position of Fe, we performed high-resolution, high-angle annular dark field (HAADF) scanning transmission electron microscopy (STEM) and lattice-resolved EDS as shown in Fig. 1 D–G. STEM images acquired with the beam along the [11$\overline{2}$0] and [1$\overline{1}$00] zone axes are shown in Fig. 1 D and E, respectively. One layer of brighter spots (the A atomic layers) is interleaved by 2 adjacent layers of darker spots (the M atomic layers). Carbon is typically not visible because of its low contrast compared with the heavier M and A atoms, and the characteristic zigzag stacking of the M_n+1X_n slabs viewed along the [11$\overline{2}$0] zone axes is seen (14–17). STEM-EDS mapping (Fig. 1F) and line scanning (Fig. 1G), which indicate that V is in the M sites following the zigzag stacking, and overlapped Fe (purple) and Sn (green) are proving that both are in the A sites of MAX phase. The compositions of the elements V, Sn, and Fe are 67, 22, and 11 at. % (atomic percentage), respectively (SI Appendix, Table S2). The molar ratio is in good agreement with the above EDS result. Thus, the chemical formula of the resultant MAX phase is close to V₂(Sn_2/3Fe_1/3)C. To prove the generality of this methodology, we used Co, Ni, and Mn instead of Fe in the starting materials and followed the same chemical synthesis process. The comprehensive characterization of MAX phases whose A sites contain Co, Ni, and Mn elements are shown in SI Appendix, Figs. S4–S6 and section S1).

MAX Phases Containing 2 Magnetic Elements.

If 2 magnetic elements can be simultaneously incorporated into the MAX structure, this would offer additional prospects for tuning magnetic properties since strong spin-electron coupling or interaction between different magnetic elements may enhance the magnetic properties as known from, for example, permalloy Ni₈₀Fe₂₀ with high magnetic permeability (18). For coincorporation of Fe and Co, the XRD pattern (Fig. 2A) showed that the main product is a 211 MAX phase with small amounts of by-products (SI Appendix, Fig. S7). SEM demonstrated the typical terraced laminate microstructure (Fig. 2B). From the EDS results in SEM (Fig. 2C), V, Fe, Co, Sn, and C (5 elements) were detected; the molar ratio of V:(Fe + Co + Sn) is very close to 2:1; and the Sn:(Co + Fe) ratio is 2:1 (SI Appendix, Table S3). Elemental mapping further provided evidence that Fe, Co, and Sn have the same distribution (SI Appendix, Fig. S8). The molar ratio of V:(Fe + Co + Sn) is very close to 2:1, and the Sn:(Co + Fe) ratio is 2:1. Therefore, the obtained MAX phase has the formula of V₂(Fe_1/6Co_1/6Sn_2/3)C. Furthermore, STEM images of the V₂(Fe_1/6Co_1/6Sn_2/3)C phase with the beam along the [11$\overline{2}$0] and [1$\overline{1}$00] zone axes are shown in Fig. 2 D and E, respectively. Furthermore, atomically resolved EDS mapping analysis (Fig. 2F) and line scan analysis (Fig. 2G) in STEM mode provided direct evidence that Fe, Co, and Sn elements all only occupy A sites.

Fig. 2.

V₂(Fe_1/6Co_1/6Sn_2/3)C MAX phase: (A) XRD pattern and (B) SEM image after acid treatment. (C) EDS spectrum indicating particles contain V, Sn, Fe, Co, and C elements. High-resolution STEM images showing atomic positions along [11$\overline{2}$0] (D) and [1$\overline{1}$00] (E) directions, respectively. The corresponding high-resolution image together with EDS mapping (F) and line scanning (G) of V-Kα, Sn-Kα, Fe-Kα, and Co-Kα signals are shown. (Scale bars: STEM images, 1 nm.) a.u., arbitrary unit; cps, counts per second.

Following the same synthesis methodology, we also synthesized several other MAX phases with 2 magnetic elements on the A site: that is, V₂(Fe_xNi_ySn_1-x-y)C, V₂(Co_xNi_ySn_1-x-y)C, V₂(Mn_xCo_ySn_1-x-y)C, V₂(Mn_xFe_ySn_1-x-y)C, and V₂(Mn_xNi_ySn_1-x-y)C. Phase composition, micromorphology, and elemental distribution are provided in SI Appendix, Figs. S9–S13.

Multielement A-Site Phases.

A multielemental feature in chemical composition of a single-phase homogeneous material can generally have drastic implications for physical and chemical properties as realized for so-called high-entropy alloys or multiprincipal element alloys (11, 12). In analogy with this materials design strategy, we posed the hypothesis that, since 1 or 2 magnetic elements from these 4 (Fe, Co, Ni, and Mn) can be alloyed with Sn at the A site, it would be logical that also 3 or even 4 magnetic elements could be simultaneously incorporated. We therefore performed the same synthesis for all combinations of 3 of these elements. The XRD and SEM-EDS results showed the formation of MAX phases of V₂(A_xSn_1-x)C, where A is a combination of 3 elements from Fe, Co, Ni, and Mn and x ∼ 1/3. The details are provided in SI Appendix, Figs. S14–S17.

Furthermore, we synthesized an MAX phase with all 4 magnetic elements as well as Sn simultaneously. The XRD pattern shows that the final product is composed of MAX phase and various tin alloys (Fig. 3A and SI Appendix, Fig. S18). The laminated morphology of the particles is similar to the above-mentioned MAX phases (such as in Figs. 1B and 2B) but with more rounded edges (Fig. 3B). EDS in SEM detected all constitutive elements (V, Sn, Fe, Co, Ni, Mn, and C) in these particles (Fig. 3C), and the molar ratio of all 4 magnetic elements to Sn was close to 1:2 (SI Appendix, Table S4). Elemental mapping of Sn, Fe, Co, Ni, and Mn corroborated that all of these 5 elements have the same distribution (Fig. 3D) and certainly occupy the A site. The obtained V₂(A_xSn_1-x)C (where again A is a combination of Fe, Co, Ni, and Mn and x is close to 1/3) thus constitutes a realization of a multielement (analogous to a high-entropy alloy) nanolaminated material, 2D in the sense that the multielement alloying exclusively occurs on the A layers.

Fig. 3.

V₂(A,Sn)C (A = Fe, Co, Ni, Mn) MAX phase: XRD pattern (A) and SEM image (B) after acid treatment. (C) The corresponding EDS spectrum indicated that particles contain V, Sn, Fe, Co, Ni, Mn, and C elements. (D) Elemental mapping on 1 particle clearing proving the uniform distribution of Fe, Co, Ni, Mn, and Sn. (Scale bars: 5 μm.) a.u., arbitrary unit; cps, counts per second.

Phase Diagrams.

Fig. 4 shows the calculated isothermal sections at 1,100 °C of the phase diagrams of V-Sn-C, V-(Sn_2/3Fe_1/3)-C, and V-Fe-C systems. These results indicate that V₂SnC is the only thermodynamically stable ternary phase in the V-Sn-C system (Fig. 4A) at equilibrium with Sn metal and vanadium carbide. In the case of Fe addition (Fig. 4B), the V₂(Fe_1/3Sn_2/3)C phase can also be at equilibrium with V₃C₂ and FeSn₂, consistent with the experimental results (Fig. 1A). However, in the V-Fe-C phase diagram, the hypothetical ternary MAX phase V₂FeC is not stable (Fig. 4C). Instead, vanadium carbides (V₂C and V₃C₂) and an Fe-rich V-Fe intermetallic phase are the most competitive phases, corroborated by the experimental results. In fact, all of our attempts to synthesize V₂AC (A = Fe, Co, Ni, and Mn; i.e., without Sn) phases failed (SI Appendix, Fig. S19).

Fig. 4.

The isothermal sections at 1,100 °C for phase diagrams of (A) V-Sn-C system, (B) V-(Sn,Fe)-C system, and (C) V-Fe-C system. (D) Gibbs free energy of phases during experimental synthesis of V₂(Fe,Sn)C MAX phase.

The Gibbs free energy values of V₂(Fe_1/3Sn_2/3)C as well as intermediate phases during synthesis are shown in Fig. 4D. According to this, V₂(Fe_1/3Sn_2/3)C can appear at temperatures as low as 400 °C. Below this temperature, the V₂C and Fe₅Sn₃ phases are dominant. Above 750 °C, V₂C and Fe₅Sn₃ gradually disappear and transform into V₂(Fe_1/3Sn_2/3)C phase by a peritectic reaction: that is, solid V₂C and an intermediate liquid Fe₅Sn₃ transform into V₂(Fe_1/3Sn_2/3)C. Without the presence of Fe (or Co, Ni, and Mn), the formation of V₅Sn₃ is instead thermodynamically favored. Fe has higher affinity to Sn and thus, a stronger tendency to form Fe-Sn alloys than V does. This should favor nucleation of VC_1-x at low temperature and promote the peritectic reaction between VC_1-x and liquid Fe₅Sn₃ alloy to form the final V₂(Fe_1/3Sn_2/3)C phase.

In general, the stable 5₂(A_xSn_1-x)C (A = Co, Ni, or Mn) MAX phases follow similar reaction paths as V₂(Fe_1/3Sn_2/3)C because of the reduced Gibbs free energy of the phase in the V-A-Sn-C system through the addition of A elements. The same is apparently true for multielement MAX phases V₂(A_xSn_1-x)C, where A is 2, 3, or 4 of Fe, Co, Ni, and Mn. The mixing entropy at the A site must, therefore, account for most of the decrease in Gibbs free energy and the corresponding thermodynamic stability.

Density Functional Theory Calculations.

Here, first-principles density functional theory calculations (19) were performed on the configurations V₂AC (A = Sn, Fe, Co, Ni, Mn) and V₂(A_1/3Sn_2/3)C (A = Fe, Co, Ni, Mn). The corresponding total energies (SI Appendix, Table S5), lattice parameters and bond lengths (SI Appendix, Table S6), and atomic charges (SI Appendix, Table S7) are provided. Side-view images of V₂SnC (Fig. 5A), V₂FeC (SI Appendix, Fig. S20A), and V₂Sn_2/3Fe_1/3C (Fig. 5C) are presented, and the corresponding charge density distributions in real space are also provided. Based on the total energies, the reaction energies for the possible chemical reactions numbered as Reactions S1 to S4 in SI Appendix, section S4 were also calculated. These results imply that A-site alloying in V₂(A_1/3Sn_2/3)C (A = Fe, Co, Ni, Mn) is thermodynamically favored. Compared with V₂AC (A = Fe, Co, Ni, Mn), V₂SnC may show a higher stability due to the lower atomic charge and fewer valence electrons of Sn according to the findings of a previous report (20). In fact, the calculation results of phonon dispersion (SI Appendix, Fig. S21) show that V₂SnC is dynamically stable, while V₂FeC is not and thus, should not exist in the current space group.

Fig. 5.

The side views (A) and PDOS (B) of V₂SnC. The side views (C) and PDOS (D) of V₂Sn_2/3Fe_1/3C. The charge densities in the real space are also shown in this figure, and the isosurface is set as 0.1 e Å⁻³.

The projected densities of states (PDOSs) of V₂SnC (Fig. 5B), V₂FeC (SI Appendix, Fig. S20B), and V₂(Fe_1/3Sn_2/3)C (Fig. 5D) were also investigated. Sn showed a small contribution in the vicinity of Fermi level in the PDOS of V₂SnC, while Fe presented large density of states around the Fermi level of V₂FeC. This is also consistent with a lower stability of V₂FeC than V₂SnC due to the higher electron energy in the PDOS. Regarding the solid solution V₂(Fe_1/3Sn_2/3)C, Fe showed a lower PDOS in the vicinity of Fermi level compared with that of V₂FeC.

Structure Stability.

In earlier work, Fe/Co/Ni have been shown to effectively promote the formation of M₃SnC₂ (M = Ti, Zr, Hf) MAX phases (10) without incorporation of Fe, Co, or Ni in the final MAX phases. In contrast, in these experiments, adding Fe/Co/Ni/Mn element into V/Sn/C raw materials did not promote the formation of V₂SnC but a series of V₂(A_xSn_1-x)C (A = Fe, Co, Ni, Mn, or their combination) MAX phases. This also indicates that the M element (or the M_n+1X_n layer) plays an important role in the incorporation of Fe, Co, Ni, and Mn elements into the A layer. As mentioned by Villars (21), at the microscopic level, it is well known that the structural stability is determined by 3 pertinent factors, namely 1) the difference in atomic radius, 2) the electronegativity difference, and 3) the electron per atom ratio (electron concentration).

Vanadium has 1 more outmost d electron but smaller atomic radius than titanium (also Zr and Hf). Thus, the corresponding M_n+1X_n layer should become more compact. Therefore, the space available for Sn atoms in V₂SnC is less than that in M₃SnC₂ (M = Ti, Zr, Hf) MAX phases. Fe, Co, Ni, and Mn have smaller atomic radii than Sn. Thus, these elements can pack closely to Sn atoms and reduce the final molar volume in V₂(A_xSn_1-x)C MAX phases, meaning that the mixture of A element and Sn would keep the structure stable in a constrained A layer. That is, the M_n+1X_n layers provide less room for the A-element layers in V₂SnC when compared with M₃SnC₂ (M = Ti, Zr, Hf) MAX phases.

However, the criterion of atomic radius cannot be the only factor determining the structure stability; the electronegativity also accounts for the bonding structure and strength in alloys. The electronegativity (Allen scale) difference between the present A elements and Sn is also very small (χ_Mn = 1.75, χ_Fe = 1.8, χ_Co = 1.84, χ_Ni = 1.88, and χ_Sn = 1.82). Obviously, the compatibility of A atoms with Sn atoms will not modify the stacking mode of A element in between M_n+1X_n layers due to similarity in electron donor–acceptor capability.

Moreover, the electron concentration effect is crucial in complex materials, such as ternary MAX phases, because of complex electronic interactions among the constituent elements. As mentioned by Mizutani (22), the Hume–Rothery rules, which are guiding principles in the search of new alloys, use e/a as its electron concentration rule to show a unique e/a-dependent phase stability (here, e is total itinerant electron of all constituent elements in a primitive cell, and a is the corresponding total atom numbers). Different phases can successively exist at a particular e/a range. The itinerant electron of Fe, Co, Ni, and Mn is the same as that of Sn, which is 2. Therefore, the coexistence of these A elements with Sn at A site of MAX phases can stabilize the original crystal structure of V₂SnC.

Magnetic Properties.

Temperature-dependent magnetization M(T) curves under 0 field-cooled (SI Appendix, Fig. S22) and magnetic hysteresis loops (Fig. 6) of the as synthesized series of V₂(A_xSn_1-x)C (A = Fe, Co, Ni, Mn, or their combination) MAX phases were studied. For V₂(Fe_xSn_1-x)C, except at 2 K, all of the magnetic hysteresis loops follow “S-shaped” character (Fig. 6A) with small coercive force and residual magnetization (SI Appendix, Table S8), suggesting that the V₂(Fe_xSn_1-x)C compound is a typical soft magnetic material (above 2 K) with saturation magnetization (M_s) gradually decreasing with increasing temperature. At 2 K, the coercive force (H_c), residual magnetization (M_r), and saturation magnetization (M_s) of V₂(Fe_xSn_1-x)C are 150.88 Oe, 0.00038 emu/g, and 0.0806 emu/g, respectively. In the case of V₂(Fe_xCo_ySn_1-x-y)C, the magnetic hysteresis loops were collected at 2, 50, 100, 200, 300, and 400 K (Fig. 6B and SI Appendix, Table S9). At 2 K, the coercive force (H_c), residual magnetization (M_r), and saturation magnetization (M_s) of V₂(Fe_xCo_ySn_1-x-y)C are 481.85 Oe, 0.0262 emu/g, and 0.1378 emu/g, respectively.

Fig. 6.

Magnetic hysteresis loops of V₂(Fe_xSn_1-x)C (A), V₂(Fe_xCo_ySn_1-x-y)C (B), V₂(Fe_xCo_yNi_zSn_1-x-y-z)C (C), and V₂(Mn_xFe_yCo_zNi_nSn_1-x-y-z-n)C (D) at different temperatures in the range from −10 to 10 kOe.

Compared with Fe on A sites of V₂(Fe_xSn_1-x)C, the magnetization with 2 magnetic elements on A sites of V₂(Fe_xCo_ySn_1-x-y)C is stronger. Therefore, it can be concluded that the magnetic properties can be tuned by adjusting the quantity and type of magnetic elements on the A sites. For instance, for V₂(Fe_xCo_yNi_zSn_1-x-y-z)C (Fig. 6C), the residual magnetization (M_r) and saturation magnetization (M_s) at 2 K are 0.1499 and 0.7400 emu/g, respectively (SI Appendix, Table S10), much higher than for V₂(Fe_xCo_ySn_1-x-y)C and V₂(Fe_xSn_1-x)C MAX phase. In V₂(Mn_xFe_yCo_zNi_nSn_1-x-y-z-n)C (Fig. 6D), which contains the antiferromagnetic element Mn, at 2 K the coercive force (H_c), residual magnetization (M_r), and saturation magnetization (M_s) are 320.4 Oe, 0.0471 emu/g, and 0.5677 emu/g, respectively. Although the ferromagnetic properties of V₂(Mn_xFe_yCo_zNi_nSn_1-x-y-z-n)C are less prominent than those of V₂(Fe_xCo_yNi_zSn_1-x-y-z)C in the range of 2 to 300 K, they are still much stronger than for the phases containing 1 or 2 magnetic elements. Thus, this approach provides a route for altering the magnetic properties of MAX phases by changing the chemical composition and component. The positions of magnetic elements (Fe, Co, Ni, Mn) in A layer are structurally identical but chemically disordered for their close atom masses and radii. If the distribution and proportion of magnetic elements in A layers can be controlled, the variation of magnetic properties of these MAX phases will be well understood. Moreover, one possible way to reduce the chemical disorder of magnetic elements is to take advantage of double-A layer in some MAX phases, such as Mo₂Ga₂C. Double-A layer in these MAX phases may provide additional out-of-plane coupling of spin electrons other than in-plane interaction, which would be promising to further enhance the magnetic properties.

Concluding Remarks and Outlook

We have demonstrated that the series of V₂(A_xSn_1-x)C phases (A = Fe, Co, Ni, Mn, or their binary/ternary/quaternary combinations) can be synthesized by A-site element alloying, providing a generally applicable route to introduce 1 or more magnetic elements in A sites and tune the resulting properties. A multielement phase V₂(A_xSn_1-x)C (where again A is a combination of Fe, Co, Ni, and Mn and x is close to 1/3) was realized. The fact that the magnetic properties are greatly enhanced for multielement A layers lends credence to this concept and offers a rich chemical space for discovering materials and properties using A-site multielement alloying strategy.

Methods

Preparation of V₂(A_xSn_1-x)C (A = Fe, Co, Ni, Mn, and Combinations Thereof).

The starting powders were mixed in a stoichiometric ratio of V, Sn, C, A, NaCl, and KCl; the experimental details are shown in SI Appendix, section S6. After grinding for 10 min, the mixed powders were placed into an aluminum oxide boat. Then, the alumina boat was inserted into a tube furnace and heated to reaction temperature during 3 h with a heating rate of 10 °C/min under an argon atmosphere. After the end of reaction, residual chloride salts were washed out from MAX phases in deionized water. Some metal and alloys impurities were further dissolved in HCl acid in order to obtain pure MAX phase powders. The weight change during various steps of the process is provided in SI Appendix, Fig. S23 and Table S12.

Computational Details.

All the first-principles calculations were performed using the VASP code (23, 24). Based on the projector-augmented wave pseudopential (25) with a plane-wave cutoff energy of 500 eV, the generalized gradient approximation as implemented by Perdew–Burke–Ernzerhof (26) was used for describing the exchange-correlation functional. Regarding atomic charges, a Bader charge analysis (27) based on 180 × 180 × 180 grid was performed. Phonon dispersion was investigated with Phonopy software (28) and the VASP code based on the density functional perturbation theory (29). All of the structures were visualized in the VESTA3 code (30).

The Calculation of Phase Diagrams (CALPHAD) approach was applied to calculate the phase diagrams. Due to the lack of experimental data on the ternary V-Sn-C, V-Fe-C, and V-Sn-Fe-C compounds, first-principles calculations were conducted to support the CALPHAD work (31). The Gibbs free energy function of V₂(Sn,Fe)C was then determined with the Neumann–Kopp rule and added in the CALPHAD-type dataset of the V-Sn-Fe-C system, which included the thermodynamic parameters of the binary V-Sn, V-Fe, V-C, Sn-C, Sn-Fe, and Fe-C systems (32–36). All of the calculations were performed with the Thermo-Calc software.

Characterization.

Phase composition of the samples was analyzed by XRD (D8 Advance; Bruker AXS) with Cu Kα radiation. The microstructure and chemical composition were observed by SEM (QUANTA 250 FEG; FEI) equipped with an energy-dispersive spectrometer. AFM analysis was performed by means of a Dimension 3100 V system (Veeco) under tapping mode. Scanning transmission electron microscopy (STEM) imaging and EDS analyzing were performed in a double-corrected FEI Titan³ 60 to 300 microscope. The magnetic properties were measured on a Quantum Design superconducting quantum interference device magnetometer.

Data Availability.

The data presented in this article are available in SI Appendix. SI Appendix contains experimental conditions, XRD patterns, AFM topography, SEM images, SEM-EDS mapping, STEM images, mapping and line scanning, density functional theory calculation, and magnetic characterization of V₂(A_xSn_1-x)C MAX phases.