Anisotropic Growth and Magnetic Properties of α″-Fe₁₆N₂@C Nanocones

Abstract

α″-Fe₁₆N₂ nanomaterials with a shape anisotropy for high coercivity performance are of interest in potential applications such as rare-earth-free permanent magnets, which are difficult to synthesize in situ anisotropic growth. Here, we develop a new and facile one-pot microemulsion method with Fe(CO)₅ as the iron source and tetraethylenepentamine (TEPA) as the N/C source at low synthesis temperatures to fabricate carbon-coated tetragonal α″-Fe₁₆N₂ nanocones. Magnetocrystalline anisotropy energy is suggested as the driving force for the anisotropic growth of α″-Fe₁₆N₂@C nanocones because the easy magnetization direction of tetragonal α″-Fe₁₆N₂ nanocrystals is along the c axis. The α″-Fe₁₆N₂@C nanocones agglomerate to form a fan-like microstructure, in which the thin ends of nanocones direct to its center, due to the magnetostatic energy. The lengths of α″-Fe₁₆N₂@C nanocones are ~200 nm and the diameters vary from ~10 nm on one end to ~40 nm on the other end. Carbon shells with a thickness of 2–3 nm protect α″-Fe₁₆N₂ nanocones from oxidation in air atmosphere. The α″-Fe₁₆N₂@C nanocones synthesized at 433 K show a room-temperature saturation magnetization of 82.6 emu/g and a coercive force of 320 Oe.

1. Introduction

Permanent magnets can provide a high efficiency and reliability for renewable energy technologies, such as wind turbines and hybrid electric vehicles [1]. Rare-earth-free permanent magnets, expected to be the next generation of permanent magnetic materials, have been paid much attention due to abundant resources, low cost, large coercivity and high Curie temperatures [2]. Body-centered tetragonal (bct) α″-Fe₁₆N₂, being an ordered nitride, is one of the most promising candidates for future applications of rare-earth-free permanent magnets because of the excellent magnetocrystalline anisotropy (7.8 × 10⁵ J/m³) and the highest saturation magnetization reported so far [3,4,5,6,7,8]. In the iron–nitrogen binary system, Jack first reported that the metastable phase α″-Fe₁₆N₂ has a giant magnetic moment [9]. Besides α″-Fe₁₆N₂, ferromagnetic iron nitrides include a hexagonal close packed ε-phase with general formula ε-Fe₃N_1+x showing a wide range of composition range with extreme values of −0.40 < x < 0.48, and the face-centered cubic γ′-Fe₄N. Face-centered cubic γ′-Fe₄N has a ferromagnetic state below its Curie temperature of about 760 K [10], while ε-Fe₃N (x ≤ 0) in a space group P6₃22 with all the N atoms occupying octahedral interstices (2c site) exhibits a well-ordered structure, showing enhancements of the Curie temperature (T_C) and room-temperature saturation magnetization (M_S) with decreasing N content, with a record-high T_C (632 K) and M_S (192 emu/g) at x = 0.88 [11]. The magnetic moment of iron atoms on the 4d site of α″-Fe₁₆N₂ is 3.0 μ_B/Fe [12]. The M_S of partially ordered Fe₁₆N₂ thin films increases monotonically with an increasing volume ratio of the α″-Fe₁₆N₂ and the N site ordering parameter related on the special arrangement of Fe₆N clusters [13]. Anisotropic α″-Fe₁₆N₂ magnets were prepared by starting from pure bulk Fe with urea as the nitrogen source [14]. Because of optimal coercivity performance of single-domain nanoparticles, the lower dimensionality of the enhanced hard magnetic properties of magnetic nanostructures has been attractive in the area of hard magnetic nanomaterials [2,15,16]. Well-controlled nanostructures of hard magnets are a possible future choice for controlling texture and magnetic alignment to support high-density magnetic storage and giant energy density [17].

α″-Fe₁₆N₂ powders and nanoparticles were obtained by mechanical ball milling with a solid nitrogen source of NH₄NO₃ [18] and ammonia nitrification of iron nanoparticles after being reduced from Fe-oxide nanoparticles [19], showing a coercivity (H_C) as high as 3.35 kOe. Core-shell α″-Fe₁₆N₂/SiO₂ [20], α″-Fe₁₆N₂/C [21] and α″-Fe₁₆N₂/Al₂O₃ nanoparticles [22] were obtained through various successive procedures starting from the reduction of Fe-oxides, followed by nitridation. The core-shell structure can restrict the growth of α″-Fe₁₆N₂ nanoparticles during the nitridation process, which improves the coercivity performance of α″-Fe₁₆N₂ nanoparticles. However, the M_S values experimentally determined in α″-Fe₁₆N₂ nanoparticles were in the range 160–225 emu/g, smaller than the value (290 emu/g) for α″-Fe₁₆N₂ thin film [9,13,23,24]. In nanowire arrays, the dipole–dipole interaction between wires becomes important [25]. By combination of the shape anisotropy and the magneto-crystalline anisotropy, ferromagnetic materials with a high magnetic anisotropy exhibit a high coercivity [15].

The most common synthesis route for iron nitrides is nitridation of iron films, powders and nanoparticles in an NH₃ atmosphere at high temperatures (≥573 K) [26,27,28,29]. Such high-temperature nitridation routes inevitably create disordered nitrogen atoms in iron nitrides. The chemical solution method is mild and highly effective to synthesize high-purity magnetic nanocrystals with controlled composition, size and microstructure by varying the reaction conditions [16,30]. Recently, nanoscale and stable ε-Fe₃N_1+x (−0.12 ≤ x ≤ −0.01) nanoparticles with highly ordered N atoms were obtained by using iron(II) acetylacetonate and tetraethylenepentamine (TEPA) as Fe and N/C sources under a lower temperature (533 K) [11]. Such a novel chemical synthesis route is favorable for preparation of monophasic ε-Fe₃N_1+x (x < 0) nanoparticles that were never synthesized previously because the ε-Fe₃N_1+x (x < 0) phases are unstable at a temperature lower than their synthesis temperatures [31,32]. Due to thermal instability at ~523 K, the synthesis of α″-Fe₁₆N₂ with highly ordered N atom occupancies is usually at a nitridation temperature lower than 453 K [33,34,35]. Improvements thereof are expected in the chemical synthesis for α″-Fe₁₆N₂ and its magnetic properties. Herein, carbon-coated one-dimensional (1D) α″-Fe₁₆N₂ nanocones with an average length of 200 nm were synthesized by a new simple low-temperature chemical solution method. Anisotropic growth mechanism, structure and magnetic performance of the α″-Fe₁₆N₂@C nanocones were studied.

2. Materials and Methods

2.1. Chemicals

Iron pentacarbonyl [Fe(CO)₅, 98%], 1-octadecene (ODE, 90%) and oleylamine (OLA, 90%) were purchased from Aladdin reagent company (Shanghai, China). Absolute ethanol (99.7%) and tetraethylenepentamine (TEPA, 90%) were purchased from Sinopharm Chemical Reagent Co. Ltd. (Shenyang, China). Prior to synthesis, the original TEPA was heated to 588 K in a nitrogen atmosphere and kept at this temperature for 5 h. Other chemicals were used without further purification.

2.2. Synthesis of α″-Fe₁₆N₂@C Nanocones

The solution syntheses were carried out at a reaction temperature range from 393 to 553 K with a varied reaction time of 1 and 6 days. When the reaction temperature and reaction time were less than 433 K and 3 days, both Fe₃O₄ and α-Fe were usually found in the products because the nitridation reactions were incomplete. In a typical synthesis of α″-Fe₁₆N₂@C nanocones, 10 mL of the treated TEPA, 30 mL of ODE and 10 mL of OLA were mixed into a 250 mL four-neck flask using a Schlenk line under air-free conditions and magnetically stirred throughout the entire reaction process. Under a mixed Ar/H₂ (95:5) flow, the mixture was first heated to 383 K. The solution was kept at this temperature for 60 min. to remove low boiling-point solvent and oxygen. Then, the temperature was increased to 473 K at 5 K·min⁻¹ and a translucent yellow microemulsion was obtained. At 473 K, a mixed solution consisted of 2 mL of Fe(CO)₅, 6 mL of ODE and 2 mL of OLA in a syringe which was injected into the microemulsion at a rate of 30 mL/h and maintained at this temperature for 30 min. for the formation of Fe nanocrystals (NCs). Subsequently, the reaction system was cooled to 433 K and maintained at this temperature for 6 days, at which temperature α″-Fe₁₆N₂ was synthesized. At room temperature, the product was precipitated by centrifugation at 3000× g rpm for 5 min. The precipitate was rewashed in 20 mL of ethanol three times and dried in a vacuum for further characterization.

2.3. Characterization

The size, morphology and microstructure of the α″-Fe₁₆N₂@C nanocones were observed using a Tecnai G2 F20 transmission electron microscope (TEM) (FEI Inc., Hillsboro, OR, USA) at 200 kV. The scanning electron microscopy (SEM) images were obtained with a JSM 6301F field-emission scanning electron microscope (FESEM) system (JEOL Inc., Tokyo, Japan). Powder X-ray diffraction (XRD) was recorded on a D/Max-2400 diffractometer (Rigaku Inc., Tokyo, Japan) equipped with a Cu K_α radiation source (λ = 0.154056 nm). The composition and surface information of the nanocones were determined by X-ray photoelectron spectroscopy (XPS) (Thermo Fisher Inc., Waltham, MA, USA). The surface of compacted samples was cleaned by argon–ion bombardment. Both the as-prepared and the surface-cleaned samples were analyzed by the XPS to determine the species covering on the surface of samples. Thermal analysis was performed on an STA6000 thermal analyzer (PerkinElmer Inc., Waltham, MA, USA) under N₂ flow with heating rate of 10 K·min⁻¹ between 300 and 900 K. Room-temperature magnetic hysteresis loop and temperature-dependent magnetization in the rising/cooling processes (at H = 1 kOe) were carried out using a vibrating sample magnetometer (VSM) in a physical property measurement system (PPMS) (Quantum Design Inc., San Diego, CA, USA) equipped with a superconducting magnet with a maximum magnetic field of 140 kOe.

3. Results

3.1. Structural Properties

The TEM image in Figure 1a shows that the product synthesized at 433 K for six days consists of fan-shaped particles with a diameter size of about 500 nm, which were formed by an assembly of nanocones (Figure 1b). The microstructure at the edges of the fan-shaped particles shown in Figure 1b and an HADDF image in Figure 1c reveal a relatively flat surface of the fan-shaped particles assembled by two or three layers of nanocones. Such an intricate microstructure forms when the as-synthesized nanocones agglomerate, in order to minimize the magnetostatic energy [29]. The nanocones have different diameters at two ends, which gradually change from ~40 to ~5 nm and the length of the nanocones is about 200 nm (Figure 1d). A small amount of Fe₃O₄ nanocrystals with an average particle size of ~4 nm can be observed on the surface of the nanocones. The HRTEM image of a cubic nanocrystal (Figure 1e) shows a lattice fringe with spacing of approximately 0.26 nm, which represents the (311) planes (0.253 nm) of the cubic Fe₃O₄ according to the Joint Committee on Powder Diffraction Standards (JCPDS) XRD card (19−0629). The Fe₃O₄ nanocrystals should come from the oxidation of excess of Fe nanocrystals (NCs) without carbon coating. The HRTEM image in Figure 1f demonstrates a typical core-shell structure of an individual nanocone with the thickness of the carbon shell being about 2–3 nm. However, there is no evident lattice fringe spacing for the nanocone, indicating the poor crystallization or crystallite constituents with very small grain size.

Figure 2 shows the phase evolution of final products synthesized at the different reaction conditions, such as the reaction temperature and time. The powder XRD pattern in Figure 2a indicates that the sample synthesized at 393 K for three days consisted of Fe₃O₄ and α-Fe. The Fe₃O₄ and α-Fe NCs agglomerated into spherical particles under the magnetostatic energy (not shown here). When the nitridation reaction was performed at 403 K for three days, the powder XRD pattern of the as-prepared sample in Figure 2b reveals that the XRD peaks can be indexed to α-Fe with the JCPDS XRD card (06−696), α″-Fe₁₆N₂ with the JCPDS XRD card (78−1865) and a small amount of Fe₃O₄ characterized by a weak (311) XRD peak at 2θ = 35.4°. When the reaction temperature and time were increased to 433 K and six days, respectively, Figure 2c shows that the product consisted of the main phase α″-Fe₁₆N₂ and a trace amount of Fe₃O₄. There is no α-Fe in the product due to the oxidation of excess of α-Fe NCs without carbon coating, as presented in Figure 1d,e. The XRD peaks of the as-synthesized α″-Fe₁₆N₂ nanocones in Figure 2c match well with the standard XRD pattern of bulk α″-Fe₁₆N₂ (JCPDS card no. 78−1865) that has a tetragonal structure with a space group of I4/mmm (139) and a = b = 5.72 Å, and c = 6.29 Å (Figure 2e). The main peak position of the product with a nanocone shape is consistent with that of (202) for α″-Fe₁₆N₂, indicating that the properly increased reaction temperature and time were helpful to the diffusion of N atoms into α-Fe nanocrystals for the formation of α″-Fe₁₆N₂. However, the (202) XRD peak around 2θ ≈ 42.7° and (220) XRD peak around 2θ ≈ 44.8° for α″-Fe₁₆N₂ emerge as a broad peak, revealing the poor crystallization or α″-Fe₁₆N₂ crystallite with small grain sizes in the nanocones. The average grain size of α″-Fe₁₆N₂ nanocrystals calculated by the Scherrer formula was about 3 nm, in good agreement with the HRTEM result shown in Figure 1d. After raising the reaction temperature and time to 453 K and six days (Figure 2d), the main peak position (2θ) of the product shifted from 42.7° for α″-Fe₁₆N₂ to 43.05° and a new XRD peak at 41.3° emerged for the (111) of γ-Fe₄N (JCPDS card no. 06−627). This reveals that extra N atoms will transform the α″-Fe₁₆N₂ to a stable phase of γ-Fe₄N. There was no carbon detected in the XRD patterns due to the slight amount. XPS was used to detect the species covering on the surface of the nanocones. Figure 3a represents the XPS spectra of the as-prepared and surface-cleaned α″-Fe₁₆N₂ nanocones synthesized at 433 K for six days. The binding energies at 284.6 and 285.7 eV in the XPS spectra in Figure 3b reveal ordered and disordered carbon, respectively, on the surfaces of α″-Fe₁₆N₂ nanocones, which is consistent with the Raman data shown as the inset in Figure 3b. Figure 3c shows the XPS spectra of Fe2p. The binding energies of Fe2p_3/2 at about 706.7 and 710.4 eV should be assigned to α″-Fe₁₆N₂ and Fe₃O₄, respectively. The XPS peak area corresponding to Fe₃O₄ decreased significantly when we extended the Ar ion sputtering time from 0 s to 60 s, confirming the slight Fe₃O₄ existing on the surfaces of α″-Fe₁₆N₂ nanocones. The weak binding energy peak at 397.3 eV featured by the blue line in Figure 3d was assigned to the N element in α″-Fe₁₆N₂, while that at 399.0 eV featured by the red line was assigned to the C-N from TEPA absorbed on the surface of the nanocones.

3.2. Anisotropic Growth

A previous study of the ε-Fe₃N_1+x@C nanoparticles revealed the nitridation and growth mechanism of core/shell structured ε-Fe₃N_1+x@C nanoparticles in a low-temperature wet chemical route [31]. In contrast to the nearly spherical shape of ε-Fe₃N_1+x@C nanoparticles, the anisotropic α″-Fe₁₆N₂ nanocones may possess a different growth mechanism. Scheme 1 presents the anisotropic growth mechanism of the α″-Fe₁₆N₂ nanocones in the ODE-TEPA solution reaction system. Unlike a previous oil-in-water (o/w) microemulsion that was stable below 333 K [36], a stable TEPA-in-ODE (TEPA/ODE) microemulsion was formed by simply heating a mixture of TEPA, ODE and OLA at a temperature higher than 393 K. It should be noted that the OLA served as a surfactant for dispersing the α-Fe nanocrystals, which is not necessary for the formation of a stable TEPA/ODE microemulsion. When a mixed solution of 2 mL of Fe(CO)₅, 6 mL of ODE and 2 mL of OLA in a syringe was injected into the TEPA/ODE microemulsion at 473 K, body-centered cubic (bcc) α-Fe nanocrystals (NCs) were first formed in the ODE phase by thermal decomposition of Fe(CO)₅. Under magnetic stirring, the bcc α-Fe NCs were transferred into the TEPA micelles. Based on the reaction mechanism [31], the α-Fe NCs catalytically decomposed TEPA to form N and C atoms. It should be noted that the reaction rate for the formation of C and N atoms may have been slow due to a reaction temperature of 433 K, which is much lower than the reaction temperature of 533 K for the ε-Fe₃N_1+x@C nanoparticles [31]. At the initial stage, a small amount of N atoms diffused into α-Fe NCs to form Fe(N) NCs. With increasing reaction time, more N atoms participated in the nitridation reaction and tetragonal α″-Fe₁₆N₂ NCs would be formed. In the case of synthesis of ε-Fe₃N@C nanoparticles in TEPA, Fe NCs easily agglomerated to form nearly spherical nanoparticles with a wide particle size distribution of 100–500 nm in diameter [11]. However, in the present case, the interface between the nonpolar ODE phase and polar TEPA micelles controlled the Fe NCs transferring from the ODE phase to the TEPA phase. Once the tetragonal α″-Fe₁₆N₂ NCs were formed in the TEPA micelles, the magnetic α″-Fe₁₆N₂ NCs could assemble along the crystallographic c axis of α″-Fe₁₆N₂ NCs to form α″-Fe₁₆N₂ nanocones. Magnetocrystalline anisotropy energy is suggested as the driving force because the easy magnetization direction of tetragonal α″-Fe₁₆N₂ NCs was along the c axis. With the reaction going on, more and more C atoms were left on the surfaces of α″-Fe₁₆N₂ nanocones, resulting in a core/shell structure. The reaction would be cut off when the α″-Fe₁₆N₂ nanocones were completely covered by the carbon shell. The overall reaction processes can be speculated as following the reaction Equations (1)–(4):

Fe(CO)₅ = Fe + 5CO

(1)

2C₈H₂₃N₅ = 10N + 16C +23H₂

(2)

2N = N₂

(3)

16Fe + 2N = Fe₁₆N₂

(4)

The catalytic decomposition of C₈H₂₃N₅ (TEPA) should occur on the surface of α-Fe NCs due to the large surface energy. Smaller α-Fe NCs allowed easier production and diffusion of N atoms into them for the synthesis of the iron nitride phase [31]. When the reaction was stopped by cooling the reaction temperature from 433 K to room temperature, the TEPA/ODE microemulsion broke and the reaction system separated into a TEPA layer and an ODE layer at about 393 K. Meanwhile, the α-Fe NCs left in the ODE phase were magnetically attracted to the thinner end of the α″-Fe₁₆N₂@C nanocones due to higher magnetic flux density than that on the thicker end (see Figure 1d). The formation of fan structures may have been driven by the reduction of high surface energy and/or by the magnetostatic energy. However, the thinner ends of the α″-Fe₁₆N₂ nanocones directed the centers of the fan-shaped particles due to a higher magnetic flux density of the thinner end than the thicker end, indicating that magnetostatic energy was the dominant driving force in the formation of fan-shaped particles. The α-Fe NCs were oxidized to Fe₃O₄, while α″-Fe₁₆N₂ nanocones were protected by carbon shells from oxidation in an air atmosphere, in good agreement with the XPS and HRTEM results. As a result, anisotropic α″-Fe₁₆N₂@C nanocones with a core-shell structure were successfully synthesized in a new TEPA/ODE microemulsion system at a low temperature of 433 K and six days.

3.3. Thermal Property

Thermal analyses of the α″-Fe₁₆N₂@C nanocones synthesized at 433 K for six days are shown in Figure 4. The TGA curve in Figure 4a exhibits two steps of weight (W) loss in a temperature between 310 K and 800 K. The first step is a slight weight loss (~1 wt.%) below an onset temperature (T_onset = ~470 K) due to the loss of organic molecules adsorbed on the surface of nanocones. The T_onset corresponds to an onset temperature for the decomposition of α″-Fe₁₆N₂ due to high temperature. The second step covers a broad temperature range between T_onset and 800 K. The temperature dependence of dW/dT clearly displayed that the stage between 470 K and 600 K corresponded to release of the N atoms from α″-Fe₁₆N₂, while the stage between 600 K and 800 K may have originated from the reduction of iron oxide by the carbon shells of α″-Fe₁₆N₂@C nanocones following a possible reaction of Fe₃O₄ + 4C = 3Fe + 4CO. As a result, the weight loss of ~8 wt.% in the second step should be ascribed to the losses of ~3 wt.% nitrogen atoms from α″-Fe₁₆N₂ and ~5 wt.% oxygen/carbon atoms due to the losses of oxygen atoms and carbon atoms. Our previous work [37] indicates that amorphous iron oxides on the surfaces of ε-Fe₃N nanoparticles become Fe₃O₄ and Fe₂O₃ crystals at a temperature near 623 K, but loss of oxygen at a temperature between 623 and 793 K. This suggests that the α″-Fe₁₆N₂@C nanocones were metastable below ~470 K and the weight loss of ~3 wt.% from 470 to 600 K is in good agreement with the weight percentage of N atoms in α″-Fe₁₆N₂. Besides the reactions for the decomposition of α″-Fe₁₆N₂ and the redox reduction of Fe₃O₄, an endothermic reaction also happened at about 600 K, as revealed by the DSC curve (Figure 4b). The possible reaction was 5Fe₁₆N₂ + 32C = 16Fe₅C₂ + 5N₂. Compared with Figure 1b, the XRD pattern in Figure 5a reveals almost no change of the α″-Fe₁₆N₂@C nanocones heat-treated at 483 K. However, the main phase became Fe₅C₂ (JCPDS card no. 51-0529) when the α″-Fe₁₆N₂@C nanocones were heated to 600 K (Figure 5b). The main phase of Fe₅C₂ heated at 600 K proves the loss of N atoms from the α″-Fe₁₆N₂@C nanocones at a temperature lower than 600 K.

3.4. Magnetic Properties

Figure 6a represents the room-temperature magnetic hysteresis loops of the α″-Fe₁₆N₂@C nanocones synthesized at 433 K for six days and at 413 K for six days, respectively, with characteristics of ferromagnetic materials. The saturation magnetization (M_S) and the coercivity (H_C) were respectively 83 emu/g and 320 Oe at 300 K for the α″-Fe₁₆N₂@C nanocones synthesized at 433 K for six days, while the M_S and the H_C became 81.6 emu/g and 250 Oe at 300 K for the α″-Fe₁₆N₂@C nanocones synthesized at 413 K for six days. The decrease in the magnetic properties of the α″-Fe₁₆N₂@C nanocones may be ascribed to a slight composition change of α″-Fe₁₆N₂ due to the reduced reaction temperature. Nevertheless, the saturation magnetizations were much lower than the previously reported M_S (290 emu/g) for α″-Fe₁₆N₂ thin film [9]. According to our previous work on the carbon-coated ε-Fe₃N_0.88 nanoparticles, the carbon shells reduce the M_S of ε-Fe₃N_0.88 nanoparticles by a value of 11.8% [11]. Thus, the M_S value of the α″-Fe₁₆N₂ nanocones would slightly increase if removing the negative effect of diamagnetic carbon. Considering the XRD pattern in Figure 2c, the fraction of Fe₃O₄ nanoparticles was very small and can be neglected, while the minimum fraction of carbon was ~8 wt.% estimated based on the Fe₅C₂ phase in Figure 5b. Taking the diamagnetic susceptibility of about −1.4 × 10⁻⁵ emu/g · Oe at 300 K for carbon [38], the M_S of α″-Fe₁₆N₂ nanocones increased at least to 90.24 emu/g at 300 K by removing the C shells from the α″-Fe₁₆N₂@C nanocones. Figure 6b plots the temperature dependence of magnetization (M-T) of the α″-Fe₁₆N₂@C nanocones in the warming and the cooling processes in a temperature range between 300 and 900 K. The M-T curves in the warming and the cooling processes were irreversible. The M-T curve in the warming process A shows a decrease of magnetization with an increasing temperature with a knee point at about 510 K. It reveals the thermal decomposition of α″-Fe₁₆N₂, which was a little higher than the temperature determined by thermal analysis (Figure 4). Below the thermal decomposition temperature of 510 K, the α″-Fe₁₆N₂ was metastable (Figure 5a). Above 600 K, the magnetization in the warming process became almost a constant. The M-T curve recorded in the cooling process B (Figure 6b) presents a magnetic transition temperature at 528 K, which was determined by the point of intersection of the two tangents around the inflection point of the cooling M-T curve. The ferromagnetic–paramagnetic transition at 528 K should be ascribed to the Fe₅C₂ phase determined by the XRD pattern (Figure 5b). Such a T_C of 528 K for Fe₅C₂ agrees with the T_C value previously reported for Fe₅C₂ [39]. There were no magnetic transitions in the cooling process B between 550 and 600 K and in the warming process A between 600 and 900 K. The absence of magnetic phases of ε-Fe₃N (T_C = 575 K), γ-Fe₄N (T_C = 760 K) and Fe₃O₄ (T_C = 864 K) reveals that all the N atoms had been released from the α″-Fe₁₆N₂ nanocones and Fe₃O₄ had been reduced by carbon in the warming process A. Figure 4a suggests that the reduction of Fe₃O₄ by carbon happened at a temperature range of 600–800 K, revealing a mass loss about 5 wt.% due to the loss of oxygen. The magnetic product of Fe₃O₄ reduced by carbon is α-Fe. The M-T curve recorded in the cooling process C (Figure 6b) shows two different magnetic transitions at 483 and above 900 K, respectively. The magnetic transition temperature of 483 K is in good agreement with the T_C of Fe₃C, while the magnetic transition above 900 K may be ascribed to α-Fe. However, the change of the magnetization after raising the temperature from 700 to 900 K in process C was not obvious due to a small amount of α-Fe and a higher magnetic transition temperature of α-Fe (T_C ≈1043 K). It is reasonable that the magnetizations in the process B at the temperature range of 500–600 K (Figure 6b) were smaller than those in the process A due to existence of α″-Fe₁₆N₂ and in process C due to existence of a small amount of Fe. It is possible that the α-Fe partially diffused into the Fe₅C₂ to form Fe₃C following the reaction equation of Fe₅C₂ + Fe = 2Fe₃C, but leaving a small amount of α-Fe in the product heated up to 900 K.

Despite the low saturation magnetization and coercivity at room temperature, anisotropic 1D α″-Fe₁₆N₂@C nanocones were synthesized for the first time in a one-pot chemical solution method. It has been found that diamagnetic carbon shells weaken the saturation magnetization of α″-Fe₁₆N₂ nanocones although the carbon shells can protect the α″-Fe₁₆N₂@C nanocones from oxidation in an air atmosphere. The Fe NCs with small particle sizes were easily nitridized to form α″-Fe₁₆N₂, but small grain sizes of α″-Fe₁₆N₂ NCs may be detrimental to their high coercivity performance. The theoretical shape anisotropy field was 2 πM_S for the nanowire with very high aspect ratio. Moreover, in nanowire arrays, the dipolar interaction between wires becomes important. The maximum field originating from dipole–dipole interactions was −2πPM_S when the magnetization was saturated perpendicular to the wire axis, and 4πPM_S when the magnetization was saturated along the wire axis, where P was a packing density of the nanowire array [40]. The low coercivity at room temperature for the present anisotropic 1D α″-Fe₁₆N₂@C nanocones may have resulted from the fan-like agglomerated structure, which can form a magnetic vortex different from the magnetic state produced by separate elongated particles. Therefore, to enhance the hard magnetic properties of the α″-Fe₁₆N₂@C nanocones, it is worth studying the syntheses of α″-Fe₁₆N₂ nanocones/nanorods with controllable grain/particle sizes and alignments of the separate 1D nanorods/nanocones instead of a fan-like structure. Moreover, as has been demonstrated in the ε-Fe₃N_1+x (x ≤ 0) [11], finely engineering the N content in the iron nitride nanostructures and removing the impurity are also critical to increase the saturation magnetization of α″-Fe₁₆N₂@C material. Our low-temperature chemical synthesis for 1D α″-Fe₁₆N₂ nanostructures has potential to obtain high permanent magnetic performance. Efforts to synthesis a magnetically oriented 1D α″-Fe₁₆N₂ hard magnet are still in progress.

4. Conclusions

We developed a facile protocol for synthesizing carbon-coated tetragonal α″-Fe₁₆N₂ nanocones by a new one-pot microemulsion method with Fe(CO)₅ serving as the iron source and TEPA as the N/C source. Anisotropic growth of the tetragonal α″-Fe₁₆N₂@C nanocones is ascribed to the strong magnetocrystalline anisotropy energy of α″-Fe₁₆N₂ nanocrystals. α″-Fe₁₆N₂@C nanocones agglomerated to form a fan-shaped microstructure due to the magnetostatic energy. 1D α″-Fe₁₆N₂@C nanocones have excellent oxidation resistance. Despite their hard magnetic characteristics with room-temperature saturation magnetization and coercivity of 82.6 emu/g and 320 Oe, respectively, the energy produced is low due to the low saturation magnetization and coercivity of the material. However, as has been demonstrated, the magnetic properties are easily tunable by the content of ordering N atoms and grain sizes. Systematic experimental work is under way to enhance the saturation magnetization and coercivity of 1D α″-Fe₁₆N₂@C nanocones. If these properties are enhanced, 1D α″-Fe₁₆N₂ aligned hard-magnetic materials can be a low-cost, nontoxic alternative to noble-metal- or rare-earth-element-based advanced magnets.