Iron Nitride Family at Reduced Dimensions: A Review of Their

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The Iron Nitride Family at Reduced Dimensions: A Review of Their Synthesis Protocols, Structural and Magnetic Properties Sayan Bhattacharyya J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/jp510606z • Publication Date (Web): 19 Dec 2014 Downloaded from http://pubs.acs.org on December 24, 2014

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The Journal of Physical Chemistry

The Iron Nitride Family at Reduced Dimensions: A Review of Their Synthesis Protocols, Structural and Magnetic Properties Sayan Bhattacharyya Department of Chemical Sciences, Indian Institute of Science Education and Research (IISER) Kolkata, Mohanpur - 741246, India Email: [email protected]

Abstract Magnetic nanoparticles (NPs) are prominent in various fields of scientific research and applications such as magnetic energy storage, magnetic fluids, biomedical fields and catalysis. Metallic Fe NPs are difficult to protect from air oxidation and iron oxide NPs suffer from reduced magnetization. Iron nitrides have the advantage of retaining high saturation magnetization (MS) for example, α″-Fe16N2 and γ′-Fe4N have average magnetic moments of 2.9 and 2 µB/Fe, respectively, which are comparable to 2.2 µB/Fe for α-Fe. The iron nitrides are ferromagnetic (FM) up to a maximum lattice dilution of 25% nitrogen for ε-Fe3N. Even though Fe-N NPs are extremely attractive as magnetic materials, the focus has been majorly on the bulk powders and thin films. This review provides a comprehensive overview of the crystal structures, nitriding kinetics, synthesis methodologies of binary, doped and ternary nanostructures, thin films and bulk materials and their magnetism. Substitution of Fe by any other metal atom in the doped and ternary nanostructures breaks the long-range FM ordering, but equally provides interesting low temperature magnetic ordering such as spin glass and exchange bias. The dopant concentration dependence of the magnetic properties of the hybrid systems is discussed. 1 ACS Paragon Plus Environment


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Keywords: Iron nitride; Nanostructures; Crystal Structure; Nitriding; Magnetism 1. Introduction Transition metal nitrides form an interesting class of commercially important compounds because of their versatile magnetic, electrical, mechanical and tribological properties. Lighter elements like hydrogen, boron, carbon, and nitrogen can occupy interstitial lattice positions of the host 3d transition metals, among which carbon and nitrogen can dissolve into the lattice of body-centered cubic (BCC) α-Fe (most stable), face-centered cubic (FCC) γ-Fe, and hexagonal close packed (HCP) ε-Fe to yield a series of solid state materials for practical applications.1 The interest in iron nitrides stems from their potential applications in high-density magnetic recording heads and magnetic recording media due to their excellent magnetic properties,2 catalysis depending on the availability of N-active sites,3-5 biomedical fields since the iron nitrides are relatively less cytotoxic than the iron oxides and as high wear-resistant and corrosion resistant coatings.6, 7 As compared to other iron compounds, the nitrides possess greater degree of magnetization than the iron oxides and are more cost-effective than FM alloys such as FePt.8, 9 These interstitial Fe-N alloys have shiny metallic colors and are hydrolytically stable. Metallurgists find iron nitrides to be attractive because of their superior tribological, corrosion-resistant and elastic properties used for hardening steels.10,

11

However, a majority of the literature on iron nitrides deal with their most attractive magnetic properties because of their high magnetic moment, which is tunable with the concentration of nitrogen in the FexNy lattice. The Fe-N phases are FM at room temperature (RT) only for atomic % of N ≤ 25 in the lattice.12 Most of these iron nitrides were synthesized from the nitridation of the cheapest possible magnetic element, Fe. The

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iron nitrides are very popular as bulk materials and thin films. However stabilizing the ordered phases of these materials at the nanometer size regime is more challenging. In fact, scaling the size of the iron nitride particles down to the nanoscale, destabilizes the spontaneous magnetization due to the superparamagnetic effects. All the available iron nitrides are metallic conductors, and the Fe-N phases are metastable in every form either in the bulk, as thin films, or as NPs, due to the kinetic constraints. The ordered phases exist only at moderate temperatures 350-550oC. Pertaining to the above limitations, the iron nitrides are relatively less studied as compared to their oxide counterparts, inspite of superior magnetic properties of the former. The binary interstitial Fe-N compounds are capable of forming mutual solid solutions or ternary phases with other metals and nonmetals. This review compiles the structural aspects and phase characteristics of different iron nitrides, kinetics of the nitriding process, the progress in synthesis methodologies of Fe-N nanostructures, substitutional lattice doping leading to the pseudo-binary phases, Fe-M-N (M = metal) ternary compounds, composites and major progress of the superlative magnetic properties of nanostructured iron nitrides. 2. Structure and Bonding In general, the interstitial alloys can be divided into two major classes: (1) interstitial solid solutions where the interstitial atoms are soluble in the host crystal structure, e.g. C in austenite, O in Ti and Zr; and (2) interstitial compounds where the interstitial atoms alter the crystal structure of the metal or alloy lattice, e.g. ε-Fe3N, γ′Fe4N. The radius ratio, rx/rm, of the non-metal (x) and metal (m) atoms play an important role in the formation of these interstitial alloys.13 Below a critical magnitude of rx/rm = 0.59, simple structures such as cubic close packed (CCP), BCC or HCP are formed and



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are characteristics of metal atoms themselves, whereas, above 0.59 more complex structures result. In this regard H, C, N, B and Si having atomic radii 0.46, 0.77, 0.71, 0.91 and 1.17 Å, respectively are potential interstitial atoms. Another important factor, which favors interstitial alloy formation, is the metal – non-metal bonding. In this regard, C and N form stronger bonds through electron donation to the incomplete inner orbitals of the metal atoms and are better interstitials. The structure of transition metal nitrides can be classified by Hagg's empirical rule,13 where the structures can be determined by metal atoms. The nitrogen atoms are filled in tetrahedral, octahedral or trigonal prismatic interstitial sites of the original host structure of metal atoms. Calculations using spin density functional theory showed that nitrogen occupies the octahedral interstices in the close packed Fe-lattice, and strong Fe-N bonds exist with σ-type p-d hybridization and charge transfer to N.14 The interstitial site has to be small enough to provide sufficient bonding between the metal and non-metal atoms and large enough to maintain the bonding between metal atoms of the host lattice. The symmetry of the metallic framework consisting of dominant Fe-Fe bonds suffers a limited degree of distortion after incorporating the interstitial N atoms. This can give rise to different Fe-N coordination and various crystal structures, directly correlated to the nitrogen content inside the lattice. Depending on the nitrogen concentration, there exists a series of binary Fe-N compounds: γ″-FeNy (y = 0.9 - 1.0) with ~50 atomic% N, ζ-Fe2N with ~33% N, εFe3N1+y (y = 0 - 0.33) with ~25% N, γ′-Fe4N with ~20% N, and Fe8N (or α″-Fe16N2) with ~11% N.10,

12

The iron nitrides (FeNy) with y < 0.2 are reported to be energetically

unfavorable.10 Nitrogen stoichiometry can be experimentally determined by Kjeldhal analysis, CHN analyzer and/or by combustion analysis which involves flash combustion


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followed by measurement of N2 concentration by gas chromatography. The nitrogen concentration can be synthetically altered to stabilize the Fe-N phases with different crystal structures and thus the electronic and magnetic properties can be tuned. For a target oriented synthesis of the Fe-N nanostructures, knowledge of the Fe-N phase diagram is crucial. The earliest work(s) done on the Fe-N phase diagram were in the 1920s and 1930s.15, 16-19 With the help of X-ray diffraction (XRD) studies, the first Fe-N phase diagram for temperatures 300-800oC was plotted by Jack (Figure 1a).20-23 The phase diagram was updated later with data on magnetic transformations in the Fe-N system (Figure 1b), and with the results of nitridation at atmospheric pressure (Figure 1c).24, 25 Other phase diagrams were plotted on the basis of thermodynamic computations, which allows wider temperature and concentration ranges.26 Below 2.4 atomic% nitrogen inside the BCC lattice of α-Fe, the N atoms randomly occupy the octahedral interstices and above that 1/3rd of the octahedral interstices are occupied resulting in tetragonal straining. When the γ-phase with 10.2 atomic% nitrogen is cooled, it undergoes eutectoid decomposition (γ → α + γ′) or shear transformation (γ → α′) (Figure 1a).26 γ′-Fe4N with ~20 atomic% N transforms to the metastable intermediate α″-Fe16N2 upon ageing or tempering whereby N atoms are removed from some specific interstitial sites. With 15-33 atomic% nitrogen, the stable composition of ε-Fe2-3N with HCP Fe sublattices show the largest homogeneity range.27 Several intermediate phases can also exist. For example, εFe24N10 has a stoichiometry between ε-Fe2N and ε-Fe3N.22 Higher nitrogen contents, 33 and 50 atomic% give rise to higher iron nitrides ζ-Fe2N and γ″-FeNy, respectively. The crystal structures were simultaneously first studied by Jack,20,

21

and are

illustrated in Figure 2. γ″-FeNy (y = 0.9 - 1.0) phase has a cubic zinc-blende type structure 5 ACS Paragon Plus Environment


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(space group F 4 3m) with lattice parameter a = 4.33 Å.28 First-principles TB-LMTOASA (tight-binding linear muffin-tin orbitals within the atomic sphere approximation) calculations show that FeN can exist with NaCl (γ′″-FeN), ZnS (γ″-FeN), and CsCl structures,29 although only the γ″-FeN phase was reported experimentally.30 In the ZnStype structure the nitrogen atoms occupy the tetrahedral interstices, whereas the NaCltype structure has the lowest energy at the theoretical equilibrium volume and the N atoms occupy the middle of the edges of FCC iron lattice. With a slight decrease in nitrogen concentration, ζ-Fe2N exists in HCP arrangement of Fe atoms, and N atoms occupy half of the octahedral interstices in each layer. ζ-Fe2N is orthorhombic and crystallizes in the space group Pbcn, the nitrogen atoms are arranged in zigzag chains parallel to the orthohexagonal b axis.10 The ε-Fe3N phase also crystallizes in the HCP crystal structure with space group P63/mmc, P6322 or P312 with two formula units i.e. Fe6N2 and is described as succession of Fe-(N)-Fe-(N) layers with N as spacer along caxis in the Fe-lattice. The vacant sites are orderly and alternately filled with N atoms. An ideal hexagonal structure of ε-Fe3N has c/a ratio = 1.633, which might decrease to 1.62 in the case of NPs.31 The changes in lattice parameters from bulk to nanosized particles are mostly related to the lattice strain involved due to size reduction. When the N content increases, the N atoms are randomly filled until half of the available sites are occupied, and thus transforms to orthorhombic ζ-Fe2N.32 γ'-Fe4N crystallizes in antiperovskite type FCC structure (space group Pm3m) with N atoms occupying one quarter of the octahedral interstices in an ordered manner. The crystal structure of γ'-Fe4N can be visualized as FCC Fe lattice with N atoms at the center of the unit cell. There are two equivalent Fesites, one occupying the corners (Fe I) and the other at FCC (Fe II) positions of the cube. 6 ACS Paragon Plus Environment

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Due to the introduction of N atoms in the BCC γ-Fe lattice, the lattice parameter increases from 3.45 Å for γ-Fe to 3.797 Å for γ′-Fe4N.33 In fact, the bulk lattice parameter of γ′-Fe4N can be effectively stabilized even in the NP form, without any significant changes in the lattice volume.34 The metastable α″-Fe16N2 phase crystallizes in the tetragonal space group I4/mmm, and can be viewed as eight BCC α-Fe unit cells, with 2 octahedral interstices occupied by the N atoms. There are three Fe positions: (0, ½, ¼), (0, 0, 0.31), and (¼, ¼, 0) with site symmetries 4d( 4 m2), 4e(4mm), and 8h(mm), respectively.35 In α-Fe, the octahedral hole is situated in between the empty octahedron of four next-nearest neighbor Fe atoms in the ab plane, and the two body-centered Fe atoms are placed along the c-axis above and below the ab plane. In the process of making α″-Fe16N2, the N atom sits inside this hole distorting the octahedron. Among a mixture of Fe-N phases in the same sample, Rietveld analysis of the XRD pattern is a prime requisite to quantitatively elucidate the individual phases. The unit cell parameters of the Fe-N phases are collected in Table 1 based on available literature.10,

21, 31, 36-45

The

variations mostly depend on the nitrogen concentration and morphology. The stability of the Fe-N phases was ascertained from the enthalpy of formation (∆Hf0 in KJ/mol) of the nitrides determined within experimental error by a high temperature solution calorimetry in molten sodium molybdate.10 The interstitial N atoms are weakly bonded to the Fe atoms and hence generates less negative ∆Hf0: γ″-FeN0.91: 47.08 ± 3.47, ζ-Fe2N: -34.30 ± 7.84, ε-Fe3N1.33: -43.33 ± 6.50, ε-Fe3N: -40.00 ± 9.87, γ′Fe4N: -12.17 ± 20.26, α″-Fe16N2: 85.2 ± 46.8 as compared to the iron oxides: FeO: -272.0 ± 2.1, Fe3O4: -1115.7 ± 2.1, Fe2O3: -826.2 ± 1.3. Thus the Fe-N compounds have



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dominant Fe-Fe bonding interactions, which enhance their magnetic moment similar to that of BCC α-Fe. 3. Kinetics of Nitridation Iron nitrides are conventionally formed as a result of the interaction of ammonia (NH3) with the solid iron precursor at temperatures ≥ 400oC. The metal precursor and the resulting nitride phases act as the catalyst in the dissociation of NH3 to atomic and subsequently to molecular nitrogen and hydrogen (2NH3 → 2N + 6 H → N2 + 3 H2). The nitrogen in its atomic state diffuses into the metal to convert it to its nitride and the extent of nitridation depends on temperature of the reaction and flow rate (flow volume per unit time) of NH3 gas.46 The NH3 decomposition on Fe and/or Fe-N surface runs parallel to the nitride formation process. The gas phase compositions were measured by thermogravimetry (TG) and kinetics of the parallel reactions were investigated on bulk Fe-alloys.47 Figure 3a shows the TG curves and the mass gain due to transformation from Fe to Fe-N. The curves at each temperature saturate in twice the time than actually shown. The inflection points below the stoichiometric concentration of nitrogen in each phase correspond to the phase transition from Fe to the nitrides. The concentration of NH3 is maximum at lower temperatures when the immediate phase transition takes place. Also, the rate of NH3 decomposition is higher at the initial stages of the nitriding process, since as Fe-N phase forms, the available surface of Fe decreases in parallel to the decrease of the N2 concentration at its surface.47 The activation energies of NH3 decomposition over Fe and Fe-N surfaces were determined to be 68 and 143 kJ/mol, respectively.48 To understand the nitriding process, the individual gas contents within the furnace was measured.47 Figure 3b shows the gas phase composition changes at 450oC.


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As expected in the catalytic ammonia decomposition, the amount of N2 is 3-fold lesser than the amount of H2. With an increase of reaction temperature, the Fe-N phases (α-Fe → γ′-Fe4N → ε-Fe3N → ζ-Fe2N) do not appear sequentially but coexist within certain temperature ranges and the reaction rates depend on the partial pressure of NH3.49 To stabilize only one Fe-N phase and eliminate the others, the nitriding reactions should be optimized at a certain temperature according to the Fe-N phase diagram, since the interstitial nitrogen content inside the lattice depends on the decomposition of NH3 over the Fe-precursor followed by diffusion of nitrogen inside the Fe/Fe-O NPs which is controlled by the nitriding temperature and time. Nitrogen is highly mobile within the Fe-N lattice owing to its interstitial character and the presence of a large number of constitutional vacancies on the octahedral sites.50 The high mobility of nitrogen inside the Fe-N lattice results in its diffusion between the different crystallographic sites by virtue of which the N-stoichiometry is altered. ε-Fe3Ny (y = 1.10-1.39) phases can be tuned by changing the temperature and time of reaction between Fe powder and NH3 / NH3 + H2 mixtures i.e. ε-Fe3N1.10 was formed at 407oC after 8 cycles of 36 h, and ε-Fe3N1.39 was formed at 367oC after 48 h of nitridation.51 As a result of diffusion of more nitrogen, the lattice cell volume increased (ε-Fe3N1.10: a = 4.718(1) Å, c = 4.388(1) Å and ε-Fe3N1.39: a = 4.791(1) Å, c = 4.419(1) Å). The interstitial nitrogen can diffuse out of the Fe-N crystal lattice by annealing at higher temperatures. When a compound consisting of both γ′-Fe4N and ε-Fe3N was treated at 357oC lower than the ε-Fe3N nitriding temperature of 550oC, new γ′-Fe4N layer forms at the cost of the existing ε-Fe3N.52 With the help of macroscopic diffusion experiments and mechanical spectroscopy, the N2 mobility has been extensively studied.24 The faster 9 ACS Paragon Plus Environment


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diffusion of N than Fe is obvious from their relative atomic sizes (rFe/rN ~ 1.6) and their self-diffusion can be probed by alternate isotopic labeling. However, the atomic size dependence was not obeyed in the neutron reflectometry measurements on [FeN/57FeN] and [FeN/Fe15N] multilayers and Fe diffusion was observed to be faster than N.53 The nitrogen mobility can also be investigated by studying the changes of the occupational order, usually associated with the long range diffusion processes,50, 54, 55 and ε-Fe3N is the best studied model since it has a wide range of ordering in its N-atom superstructure.51, 56, 57

In the well-ordered ε-Fe3N, nitrogen occupies only the Wyckoff 2c sites. From the

intensities of N-superstructure reflections of in situ neutron powder diffraction of previously quenched ε-Fe3N powder, the nitrogen long range order was found to decrease at high-temperatures due to the partial transfer of N from the 2c to 2b sites leaving the 2d site virtually empty (Figure 4a-d).58 The partially disordered quenched states exhibit weaker superstructure reflection and higher axial ratio (c/a) as compared to the wellordered ε-Fe3N. Annealing the sample at 100-135oC reordered the N-superstructure (Figure 4e) and this transition follows a first order rate law with an Arrhenius-type temperature dependence of the rate constant. The reordering in the P6322 space group symmetry was best explained by the transfer of N from a disordered position (2b site) to an empty ordered position (2c site) involving an activation energy of 144 ± 5 kJ/mol. The γ′-Fe4N phase can be viewed as an interstitial solid solution FeI4(NII,VII)1(NIII,VIII)3 consisting of the FCC sublattice (I) of Fe, occupied (II) and unoccupied (III) octahedral sites in the ordered and disordered interstitial sublattices filled by N and vacancies (V), respectively.59 When N content is less than 20 atomic% the sublattice II is modified by N vacancies as constitutional defects and the relative occurrence of N atoms in the 10 ACS Paragon Plus Environment

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tetrahedral clusters increases.60 In real conditions, thermal defects also result in the transfer of N-atoms to the vacant sites, thus configurational entropy increases. Microcrystalline γ′-Fe4N powder was converted to ε-Fe3N0.95(2) single crystals at high pressures of ~15 GPa and ~1327oC, mostly due to segregation of N-poor phases at the grain boundaries or formation of elemental iron.61 Also ε-Fe3N particles annealed at 400oC, partly converted to γ′-Fe4N with enrichment of nitrogen in the remaining εphase.62 Thus depending on the treatment conditions, nitrogen diffusion results in mutual transformation of Fe-N phases. The nitrogen activity constant is a function of the number of gas molecules dissociated on the precursor metal surface per unit time and directly proportional to the product of NH3 flow rate and degree of dissociation. High flow rates lead to nonequilibrium conditions inside the furnace leading to lesser dissociation of ammonia and an equilibrium state is only possible between the solid phase and the gas. At the initial stage when N2/NH3 comes in contact with the solid surface, dissociative adsorption of NH3 takes place simultaneously, followed by gaseous diffusion inside the solid.47, 49 The dissociative adsorption of NH3 is the rate-limiting stage. Nitridation is a typical diffusion process whereby nitrogen penetrates through the surface of the metal and invades the core. It is a slow process and in case of bulk iron samples, iron nitride only forms a surface layer at high temperatures. The inward diffusion of N thus creates double/triple layers of different Fe-N phases with highest N-content at the surface, causing a depth dependence of the lattice parameters a and c.63 Since NPs have a large surface area, the diffusion rate of N2/NH3 has a comparatively lesser influence on the overall composition of the nitride phases. At the initial stage of the reaction, the dissociative adsorption of



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ammonia at the precursor NP surface is the rate limiting step.47 After nucleation, nitridation can progress uniformly throughout the NP without enough dependence on the nitrogen concentration gradient. The NPs have orders of magnitude larger defect concentration in the form of grain boundaries and dislocations than the bulk. The grain boundaries can act as diffusion pathways for nitrogen and have high density of nucleation sites which promotes nucleation of nitride phases.64,

65

Although the nitriding rate

decreases with decreasing concentration of the available sites on bulk α-Fe, the rate is influenced by the grain size distribution in NPs.66 Initially when larger α-Fe crystallites react, the reaction rate increases, and thereafter the rate decreases during the reaction of the smaller crystallites. If the inlet gas flow rate is optimized, the metastable Fe-N phases and their particle morphologies can be controlled. The gaseous diffusion depends on the optimum particle size of the iron precursors which in turn influences the Fe-N phases formed at a particular temperature and NH3 flow rate. The highest yield of 27 nm α″Fe16N2 particles was obtained from 100 nm α-Fe particles at 140oC, whereas diffusion of nitrogen into the α-Fe lattice was negligible for other particle sizes (60 and 300 nm) of the same precursor.67 4. Synthesis of Iron Nitrides Both chemical and physical procedures were used to synthesize the iron nitride binary, doped, ternary composite phases. For each of these systems, the discussion on the fabrication procedures of bulk materials and thin films preceded the synthesis protocols of the nanostructures to provide a comprehensive overview. The various synthesis methods of the Fe-N nanostructures are schematically represented in Figure 5. (a) Binary Phases: Bulk Materials and Thin Films 12 ACS Paragon Plus Environment

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A variety of iron nitrides were fabricated in the form of thin films, mostly starting from α-Fe precursors. Fe-N films were made from degreased Fe foil with 1% anhydrous hydrazine and Argon gas stream at 400oC.68 ε-Fe3N films were formed either by magnetron sputtering on Fe target in N2 and Ar flow,69 reactive ion beam sputtering,70 or by chemical vapor deposition of Fe(acac)3 and anhydrous NH3 on 50 µm thick polycrystalline Ti substrates at 600-800oC.71 In the case of thin film and bulk samples, homogeneity of the Fe-N phases is always a concern and the relative concentration of the phases changes based on the fabrication parameters employed. In the ion assisted sputter deposition technique, phase evolution (ε-Fex(≈

2)N

→ γ″-FeN → γ″′-FeN) and

nanocrystallinity of the films depend on the N2+ ion energy and flux ratio JN/JFe.72 α-Fe thin film electrodes were electrochemically converted to γ′-Fe4N + α-Fe films at 450oC 723 K when Li3N was used as the nitride ion source, Al as the counter electrode and LiAl alloy as the reference electrode.73 Plasma nitriding is another method where intense electric fields generate ionized molecules of the gas in close vicinity of the metal to be nitrided. Plasma nitriding 1 mm thick Fe disc resulted in a mixture of ε-Fe2-3N and γ′Fe4N.74 Similar phases were obtained by plasma nitriding electrolytic Fe with a plasma glow discharge in 80:20 H2:N2 at 400 Pa for 0.5-1.5 h,75 at 9 mbar pressure and 500oC up to 6 h,76 and pulsed plasma nitridation of Fe foil in NH3 at 2.5 mbar pressure for 24 h.77 A mixture of Fe + ε-Fe2+xN + ε-Fe2N layers were obtained by irradiation with 22 keV N2+ ions supplied by an electron cyclotron resonance ion source at RT.78 In fact plasma assisted implantation of N-ions on the surface of steel is a way to enhance resistance to wear and corrosion.79 Single phase Fe-N thin films could be deposited by direct current (dc)-magnetron sputtering with mixed Ar/N2 discharges on glass substrates.80 The single 13 ACS Paragon Plus Environment


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phase is sensitive to the nitrogen fraction in the gas mixture.81 Thin films of ZnS-type γ″′FeN were prepared by dc-magnetron sputtering on Si, soda-lime glass and copper substrates using N2 as sputtering gas.36, 82 FeN0.7 films consisting of ~30 nm randomly distributed grains were deposited by reactive ion beam sputtering with N2/Ar gases maintained at an overall pressure of 3 × 10-3 mbar during deposition.83 γ′-Fe4N films could be prepared by exposing sputtered Fe films to NH3-H2 (g).84 Molecular beam epitaxy (MBE) is a prominent technique to deposit single phase films. 800 nm thick Fe films deposited on Al2O3 wafers were converted to γ′-Fe4N1-x (x ≈ 0.05) films at a constant flow of a NH3/H2 mixture in a vertical nitriding furnace.85 Reactive sputtering could also deposit single phase 58 nm thick epitaxial γ′-Fe4N films with (100) and (110) orientations.86 The films and sheets of iron nitrides are popularly studied since they have the potential of direct incorporation into devices. Fe and Fe-Ni alloy sheets were prepared by vacuum induction, hot rolling and cold pressing, after which the sheets were nitrided in 1:1 NH3:H2 mixture at 500oC for 16 h to synthesize the Fe4N and (Fe, Ni)4N (Ni = 6, 12 and 20 %) sheets.87 Acicular structured γ′-Fe4N + α-ferrite sheets were synthesized from 0.75 mm thick Fe-sheet in the presence of 1.3:98.7 NH3:H2 mixture at 840oC for 2 h.88 A Ni catalyst coating of nanometer thickness on the Fe layer can control the decomposition of NH3 on the surface.89 The Fe-N phases formed by gas nitriding also depends on the temperature of reaction, composition of the gas mixture and time of exposure of the metal precursor to the flowing gas. A gas mixture of NH3 + H2 at 275325oC resulted in γ′-Fe4N within 3 min and ε-Fe3-xN at longer durations. Fe-N powders can also be prepared by firing iron hydroxides at 600-800oC in NH3 (g).90 When Fe powder was converted to γ′-Fe4N particles at 500oC with 40-60% NH3 in the NH3/H2 14 ACS Paragon Plus Environment

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flow, whereas other flow compositions result in a mixture of Fe-N crystallographic phases.84 Fe-N phases could also be obtained by spark erosion whereby α-Fe ingots were used as electrodes in NH3 at ~108 Pa and ~104 K.77 Also, high pressure methods are useful to synthesize ternary, pseudo-binary and binary Fe-N phases. For example ζ-Fe2N was transformed to single crystalline ε-Fe3N1.47 with space group P6322 under 15(2) GPa pressure and 1327oC using a resistivity-heated Walker-type two-stage multi-anvil device.38 The high pressures typically required are 10-15 GPa at temperatures of >1300oC.37 Phase pure α″-Fe16N2 is difficult to obtain both in the form of epitaxial films and powders.35 Relatively phase pure α″-Fe16N2 films was grown on GaAs (100) and In0.2Ga0.8As (100) by converting the MBE grown epitaxial Fe films under low pressure NH3-N2.91 α″-Fe16N2 was also obtained by N2+ ion implantation on Fe films grown on MgO (111).92 Mechanical tempering, straining or ageing α′-N-austenite yielded lesser pure α″-Fe16N2.93-97 The α″-Fe16N2 films with highest concentration of 36% were recently deposited by MBE in ultra high vacuum conditions.98 (b) Binary Phases: Nanostructures The stabilization of the binary Fe-N phases at the nanoscale provides opportunities for tuning the size, morphology and surface area of the particles, confinement inside carbon nanotubes (CNTs) or intercalation within graphene/graphene oxide sheets so as to generate electronic, magnetic and catalytic properties notably different from those at the bulk or thin films. Radio-frequency (rf) reactive sputtering was used to fabricate a series of 3-18 nm Fe-N granular thin films on Si (100). The NPs were embedded in the amorphous Fe-N matrix and exhibit in-plane uniaxial magnetic anisotropy.99 4-10 nm FeN particles were confined inside CNTs of inner diameters 4-8 15 ACS Paragon Plus Environment


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nm and outer diameters 10-20 nm (Figure 6a).3 The CNT channels were filled by ferric nitrate, which after controlled drying at 140oC in air were treated with NH3 at 450oC for 2 h. The above protocol provides an opportunity to stabilize the nitride phases from air oxidation and if quantitatively filled inside wider CNT channels, may lead to enhanced magnetic anisotropy of the hybrids.100,

101

Fe2N NPs dispersed on nitrogen-doped

graphene oxide was synthesized by chemical impregnation of iron acetate and a Ncontaining precursor such as 1,10-phenanthroline in ethanol to form the N-coordinated Fe complex, followed by drying, ball milling and thermal treatment in NH3 at 800oC for 2 h.102 The stability region of Fe-N phases can be shifted for nanometric materials. ~20 nm Fe particles were nitrided in 150 sccm of NH3/H2 mixtures at 400°C stepwise up to pure NH3.103 The obtained iron nitrides were reduced in the same gas mixture decreasing stepwise down to the pure H2. Stoichiometric Fe2N NPs were obtained at the final nitrogen concentration of ~11 wt% in the lattice. Additionally, the morphology of Fe-N nanostructures can be tuned by wet-chemical and solvothermal methods.104 A hydrothermal reaction was carried out on a mixture of aqueous FeCl3, isopropanol and nitilotriacetic acid to prepare the precursor for 50-60 nm thick ε-FexN (2 < x < 3) nanowires of several micrometers in length.105 Literature reports on ε-Fe3N NPs are the most abundant among all its counterparts, probably because of its wide chemical stability window and FM properties. Chemical vapour condensation of the reaction product of Fe(CO)5 and NH3/Ar (g) on a rotating chiller was used to synthesize 10-40 nm ε-Fe3N NPs passivated with 4-6 nm amorphous Fe3O4 or α-FeOOH shell (Figure 6b).106-108 ε-Fe3N NPs were formed at the gaseous composition of 10 sccm NH3 + 30 sccm Ar according to the reaction: 6 Fe(CO)5


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+ 2 NH3 → 2 Fe3N + 3 CO↑ + 3 H2↑. Core-shell morphologies of ε-Fe3N/γ′-Fe4N/α-Fe and γ′-Fe4N/α-Fe were obtained at compositions 10 sccm NH3 + 100 sccm Ar and 10 sccm NH3 + 185 sccm Ar, respectively.108 Sponge-like ε-Fe3N having surface area of 39 m2/g was synthesized by sol–gel route involving in situ nitridation, thus avoiding high temperature ammonolysis (Figure 6c, d).109 A self-expanding polypeptide foam was prepared by mixing aqueous solutions of ferric nitrate and gelatin and the dried casted films were treated in N2 at 700oC to form ε-Fe3N via nucleation of Fe3O4 NPs. This spongy nitride showed promising activity and stability in the catalytic decomposition of ammonia for hydrogen production. ε-Fe3N based magnetic fluids are usually synthesized from Fe(CO)5 (Figure 6e).110, 111 In one process, Fe(CO)5 vapour was mixed with Ar and NH3 gases and passed through a porous plate into the carrier liquid composed of αolefinic hydrocarbon synthetic oil and succinicimide as the surfactant maintained at 182oC.110 The ultrafine Fe particles formed from the decomposition of Fe(CO)5 could reduce the activation energy of the NH3 decomposition from 377 to 167 kJ/mol and εFe3N NPs coated with ~12 nm thick surfactant layer, homogeneously dispersed in the carrier liquid was obtained. In the absence of surfactant, the ε-Fe3N NPs appeared as ~23 nm clusters, due to agglomeration.112 In another process, 5-10 nm ε-Fe3N particles dispersed in kerosene were synthesized by vapour-liquid chemical reaction between Fe(CO)5 and NH3.111 The reactor containing kerosene solution of Fe(CO)5 and polybutenylsuccinpolyamine was heated stepwise up to 185oC for 1 h in the presence of NH3 introduced at 390 cm3/min. The magnetic fluids were distilled to condense the fluids or to replace kerosene with high boiling hydrocarbon oils as the carrier liquid.



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The Fe-N phases obtained from nanosized Fe-oxide precursors involve sol-gel process with subsequent nitriding the oxide phases. NH3 gas can decompose to nascent hydrogen and nitrogen at lower temperatures and hence preferred over nitriding in N2 flow, since bond energy of gaseous N2 is ~945 kJ/mol.113 The hydrogen atoms reduce the solid precursor and the latter is simultaneously nitrided due to inward diffusion of atomic nitrogen from the surface of the reduced particles. The extent of diffusion is dependent on the flow rate of NH3, temperature, duration of the reaction and the characteristics of the oxide precursor. In fact, synthesis of the desired oxide precursor is the key step for nitridation to occur at lower temperatures, which allows efficient control of the particle size of the nitride. Higher crystallinity and larger particle size lowers the reactivity of the oxide precursor since the bond dissociation energy of Fe-O is ~409 kJ/mol, and higher temperatures are required to replace oxygen by nitrogen atoms inside the lattice. Hence improper reaction conditions might lead to oxynitride phases from incomplete nitridation of the precursor. 36-60 nm ε-Fe3N1.17 particles were obtained from α-Fe2O3 NPs in NH3 at 550oC for 3 h.113 The amorphous α-Fe2O3 precursor was synthesized from ferric nitrate at pH~9.3 and ageing at 100oC for 48 h. The semi-crystalline oxide precursors could also be prepared by air-ignition of citrate or oxalate gels, the latter synthesized by refluxing stoichiometric amounts of citric/oxalic acid and metal nitrates. Oval shaped ~20 nm εFe3N particles were obtained from ultrafine α-Fe2O3 in NH3 atmosphere at 500oC for 4 h, the oxide precursor being prepared by decomposition of iron oxalate complex.41 Temperature and the nitriding time could change the nitrogen stoichiometry since 200 cm3/min NH3 treatment of α-Fe2O3 NPs synthesized from commercial ferric citrate yielded ε-Fe2.98N1.2 and ε-Fe2.99N1.1 at 550oC for 12 h, and ε-Fe2.999N1.001 at 700oC for 4 h 18 ACS Paragon Plus Environment

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(Figure 6f).31 γ'-Fe4N NPs initially formed at 500oC, can convert to ε-Fe3N at higher temperatures. A slight excess of nitrogen can introduce hexagonal ε-Fe2-3N and orthorhombic ξ-Fe2N impurities. If the diffusion of nitrogen is poor inside the NPs, metallic Fe/Fe-O/Fe-O-N can coexist with γ'-Fe4N and ε-Fe3N. Smaller NPs of the binary nitrides of high purity can be synthesized at lower temperatures, if synthesized from the precursor NPs with high surface area enabling better diffusion of nitrogen atoms. α-Fe NPs are prone to air oxidation and form a surface oxide layer. Hence these metallic NPs are difficult to stabilize in air, if the NP surface remains unprotected by surfactants. This was evident from the higher nitriding temperature (500-550oC) required to convert α-Fe NPs to ε-Fe3N0.99 by diffusion of nitrogen through the porous oxide surface layer.42 The α-Fe NPs were synthesized by sodium borohydride reduction of FeSO4. In another method, Fe2O3 could be converted to 200 nm ε-Fe3N1+x particles when autoclaved with molten NaNH2 at a significantly low temperature of 240oC.114 Reports on the synthesis of phase pure γ′-Fe4N NPs are not vast compared to those for ε-Fe3N because of difficulties in stabilizing γ′-Fe4N. ~70 nm clustered particles of γ′-Fe4N with ε-Fe3N impurities were synthesized from citrate gel precursor but with lower NH3 flow rate of 54 cm3/min at 600oC.34 Nanocrystalline iron was nitrided at 500oC to a mixture of α-Fe and γ′-Fe4N, where the bigger iron crystallites could be converted to γ′-Fe4N at lower nitriding potential and vice versa.115, 116 Solvothermal route was adapted for the synthesis of sub-micron sized γ′-Fe4N particles with Fe3C coating.117 The synthesis was carried out at 400oC for 12 h with FeCl2.4H2O as the Fe source in the presence of NH4Cl and NaN3 as nitrogen sources. Reactive gas flow sputtering at N/Fe



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ratio of ~0.2 could synthesize 50 ± 15 nm diameter γ′-Fe4N particles with α-Fe impurities (Figure 6g).118 The sputtering was carried out with a mixture of Ar and N2 gas flow maintained at 500 and 0-0.5 sccm, respectively at 150oC inside an evacuated chamber. εFe2-3N + α-Fe phases were formed when the N/Fe ratio was increased to 0.4. However, pure γ′-Fe4N powder was obtained from freshly prepared FeOOH, after the latter was reduced to α-Fe NPs under flowing H2 at 440oC and later reacted with a flowing mixture of NH3 + H2 at 350oC.119 Highly FM α″-Fe16N2 fine powders with decent yield and reproducibility were obtained by tuning the particle sizes of the oxide precursors.120 The oxide precursors consisted of Fe3O4 and γ-Fe2O3 particles which were synthesized by reacting iron acetylacetonate (Fe(acac)3) and non-aqueous benzyl alcohol in an autoclave at 200oC for 48 h. The ~40 nm oxide particles were reduced to 100-300 nm α-Fe particles in H2 flow at 400oC for 10 h and nitrided in NH3 at 160oC for 15 h to give NPs of α″-Fe16N2. Smaller particle size of the oxides resulted in lower reduction temperatures to nanosize αFe with high surface area and thereby requiring lower nitriding temperature for the formation of α″-Fe16N2. The yield of the nitride phase was 66 wt% and the rest consisted of unreacted α-Fe NPs. The particle size of the iron oxide precursor largely determines the final phase composition of the nitride product. 30, 60 and 200 nm α-Fe2O3 particles were successfully reduced in hydrogen stream at 500oC to 120 nm α-Fe particles which were nitrided in NH3 flow at 130-170oC for 100 h to yield ~27 nm α″-Fe16N2 particles with unreacted α-Fe.67 The 60 nm α-Fe2O3 particles resulted in the best yields of α″Fe16N2. Monodispersed < 30 nm α-Fe particles synthesized from Fe(CO)5 in octyl ether


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with equimolar mixture of oleic acid and oleyl amine as stabilizer also holds promise for lower nitriding temperatures in the stabilization of α″-Fe16N2 NPs. The method of reducing the Fe-oxides in H2 at 250-600oC for 3-5 h and subsequent nitriding the Fe NPs in NH3 or NH3 + H2 (g) mixture at 100-200oC had been used to make ~20 nm α″-Fe16N2 NPs with a thinner oxide coating.121,

122

The surface of the α″-Fe16N2 NPs could be

coated with Y or Al compounds to protect the NPs from oxidation and sintering.123 α″Fe16N2 nanocrystallites embedded in an amorphous Fe92B8 matrix was obtained by annealing the amorphous ribbon in 92:8 NH3:H2 at 600oC for 45 min, homogenized in Ar for 30 min, quenched in water and tempered at 150oC for 2 h.124 The nitrogen content of the amorphous matrix was found to be 7.8 atomic%. (c) Doped Phases: Bulk Materials and Thin Films Iron nitrides were doped by replacing iron atoms in the Fe-N lattice by Group IIIA elements or transition metals without altering the crystal structure of the host nitride for the purpose of modulating the magnetic properties of the materials. γ′-Fe4N and εFe3N are the two most commonly studied host phases used for doping. The doped Fe-N bulk materials were synthesized by conventional mixing and sintering processes. Phase pure GaxFe4-xN (x = 0.125-1) powders were obtained by stoichiometrically mixing and grinding Ga2O3 and Fe2O3 powders, followed by sintering at 1100oC for 1 min and nitriding the sintered powder in 1:1 NH3:H2 gas at 530oC for 3 h.125,

126

A similar

approach was followed for the synthesis of InxFe4-xN (0 < x < 0.8) wherein In2O3 and Fe2O3 powders were stoichiometrically mixed, ground and sintered at 600oC for 5 h. The sintered powders were nitrided in 1:1 NH3:H2 gas at 500oC for 3 h.127 Phase pure InxFe4xN

was obtained only till x ≤ 0.75. Higher doping led to elemental indium as a side phase, 21 ACS Paragon Plus Environment


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which could be dissolved by diluted HCl washing leading to phase pure In0.8Fe3.2N. 200 nm thick (Fe1-xNix)4N (0.2 ≤ x ≤ 0.6) films were prepared on Si (111) substrates by electron cyclotron resonance microwave plasma assisted evaporation method.128 The Si substrate was kept at 450oC during deposition in the presence of N2 inside the evacuated chamber. The samples were inserted into a quartz tube and were heated at 550oC. Mndoped Fe3N powders according to the stoichiometry (Fe1-xMnx)3N (x = 0.25, 0.5, 0.6) were synthesized from mixing Mn2N and Fe2N powders and treating the mixture twice in NH3:H2 gas at 600oC.129 (d) Doped Phases: Nanostructures The extent of achievable doping at the nanoscale is sensitive to the synthesis procedures, surface passivation effects and particle size of the doped Fe-N products. NPs tend to be pure eliminating foreign elements from the lattice and hence in most cases, side phases exist above 20 to 40 atomic% doping leading to composite materials. The literature is limited and it was observed that only Cr, Co, Ni and Ga were doped in εFe3N and γ′-Fe4N lattices at reduced dimensions. The ε-Fe3-xNixN (x ≤ 0.4) system was synthesized by citrate precursor route where aqueous solutions of Fe(NO3)3.9H2O and Ni(NO3)2.6H2O was refluxed with citric acid solution at 70oC for 12 h. The dried Fe-Nicitrate precursor gel was decomposed in air to Fe-Ni-oxide NPs at 500oC for 4 h. The decomposition was due to the elimination of adsorbed water, CO (g) and the conversion of carboxylate groups, methylene groups and hydroxyl groups to form acetone and CO2. The Fe-Ni-oxide NPs was nitrided in flowing NH3 gas at 400-550oC for 12 h (Figure 6h).130 Incorporation of Ni into γ′-Fe4N lattice could result in single phase doped nitrides over a wide doping range. 60-80 nm γ′-Fe4-xNixN (0.2 ≤x ≤ 0.8) particles were also 22 ACS Paragon Plus Environment

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synthesized by citrate precursor route starting from ferric citrate and nickel chloride and the mixed metal citrate was decomposed in air to Fe-Ni-oxide particles at 400oC for 2 h, followed by nitriding in NH3 (g) (140 cm3/min) at 200-550oC for 3-15 h.45 Cobalt substituted ε-Fe3N phases were obtained by oxalate precursor route or NaBH4 reduction of stoichiometrically mixed metal salt precursors followed by nitriding in NH3 (g) at 500550oC for 4 h.131, 132 It was observed that only 20 atomic% Cr can be incorporated into the ε-Fe3N lattice to give the ε-Fe2.8Cr0.2N phase. Similar citrate precursor route was adapted to synthesize the Fe-Cr-oxide NPs which were nitrided at 710oC for 24 h (Figure 6i).133 Single phase ε-Fe2.8Ga0.2N was similarly synthesized by nitriding Fe-Ga-oxide NPs under flowing NH3 (g) (flow rate: 200 cm3/min) at 550oC for 24 h.134 At higher doping concentrations, composites of ε-Fe3N and GaN were obtained. Recently a new doped material α′′-(Fe1−xCox)16N2 (x = 0.03) phase was suggested by XRD and X-ray absorption spectroscopy.135 ~100 nm particles of a mixture of α′′-(Fe0.97Co0.03)16N2 and BCC Fe0.97Co0.03 alloy was synthesized from (Fe0.97Co0.03)3O4, prepared by hydrolysis of iron and cobalt acetylacetonate in benzyl alcohol. (Fe0.97Co0.03)3O4 was reduced to the Fe0.97Co0.03 alloy at 400°C in H2 and finally nitrided at 150°C. (e) Ternary Phases: Bulk Materials and Thin Films Ternary Fe-N phases were mostly synthesized in bulk morphologies and as thin films. In order to develop ternary nitrides, elements on the right side of Fe in the periodic table can be used to substitute Fe based on the difference of chemical affinity between ironnitrogen and metal-nitrogen.136 A new ternary iron nitride ε-Fe2IrN0.24 was synthesized by high-pressure reaction of ζ-Fe2N and Ir mixture at 8-15 GPa and 1100-1250oC,137 similarly to the other metastable nitrides ε-Fe2CoN and ε-Fe2NiN.138 GaFe3N powders 23 ACS Paragon Plus Environment


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were synthesized from Ga2O3 and Fe2O3, stoichiometrically mixed in the ratio 1:3, ground and treated under 1:1 NH3:H2 in two steps.139 The micron sized powders were briefly sintered at 1100-1200oC for 1 min and nitrided at 500oC for 3 h. GaFe3N crystallizes in the perovskite structure with space group Pm 3 m and lattice parameter a = 3.7974(1) Å in which Ga occupies the Fe (I) sites of the γ′-Fe4N lattice. AlFe3N was similarly synthesized by replacing Ga2O3 with Al as the starting precursor. γ′-InFe3N was prepared by ball-milling commercial indium powder with 15 µm ζ-Fe2N powder, previously synthesized by ammonolysis of Fe powder at 347oC.136 Grinding was performed keeping ball-to-powder ratio at ~10. Due to the lower chance of diffusion of In inside the lattice, the milling was continued for 64 h followed by heat treatment under inert atmosphere at 557oC to homogenize the disordered perovskite structure with lattice parameter a = 3.8673 Å. γ′-ZnFe3N along with an impurity phase was synthesized by replacing indium with commercial zinc powder. The lattice parameter of the perovskite γ′-ZnFe3N was a = 3.8003 Å and the crystallite size calculated from the broadening of XRD peaks was ~20 nm. Although the electronic structure of a similar anti-perovskite CoFe3N was predicted by plane-wave pseudopotential method,140 the phase was not synthesized experimentally until the Co3FeN thin films were fabricated by MBE recently.141 A semihard itinerant ferromagnet RhFe3N as synthesized by synthesized by ball milling Rh and Fe/ε-Fe3N and subsequent nitridation.142, 143 Rh occupies the Fe (I) sites and has a MS of 8.3 µB per formula unit. Rf diode sputtering was used to fabricate Co-Fe-N films from Co-Fe (Co 49% Fe 49% V 2%) targets kept initially under vacuum of ~10-7 Torr followed by introduction of N2/Ar mixture.144 The final products were deposited on water cooled glass substrates. DC 24 ACS Paragon Plus Environment

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magnetron sputtering was used to deposit Fe-Cr-N films with α-Fe(Cr) side phases on water cooled glass and Al foil substrates from target consisting of Fe (99.9%) + Cr (99.9%) plates in Ar/N2 plasma.145 Fe-Cr-N ingots were prepared by melting Fe–Cr (22, 24%)–N (1.1%) alloys in an induction furnace at 1200oC under N2 pressure up to 3 MPa.146 The tendency of phase separation at higher temperatures affected the phase purity of the nitride. Alloys such as Y2Fe17N2.3, Y2Fe17C1.0N2.3, Sm2Fe17N2.1, Sm2Fe17C1.1N2.3 were prepared by introducing interstitial nitrogen in the mixed metal alloys under N2 at 450-650oC.147 To improve thermal stability of the Sm2Fe17N2.1 alloy, Co was doped in Sm2(Fe1-xCox)17 by induction melting to get Sm-Fe-Co-N powders.148 Similar alloys were prepared by high energy vibration ball milling, vacuum annealing at 700-800oC for 1 h and nitriding the alloys in N2 at 460oC for 3 h.149 The other type of ternary iron nitrides consist of η-carbide type Fe-W nitrides (Fe3W3N, Fe6W6N) which were synthesized by hot isostatic pressing method when cold pressed mixture of Fe and W powders were subjected to 1500oC and 200 MPa pressure for 1-2 h.150 (f) Ternary Phases: Nanostructures Ternary Fe-based nitride nanostructures exist either as nanocrystalline films or as NP powder, the precursors of the latter synthesized by sol-gel routes. Magnetron sputtering was utilized to fabricate nanocrystalline soft magnetic FeAlN films with grain sizes 9-15 nm on glass or Si (111) substrates either ex-situ or in-situ and further crystallized by vacuum annealing at 300-350oC.151 The ex-situ process consisted of the growth of FeAlN films on water cooled substrates and post-annealed at 200-450oC under N2 to give α-Fe phase structure with lattice dilation of 0.19-0.37%. In the in-situ process,



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the FeAlN films were grown on substrates maintained at 60-220oC under N2 to provide a composite phase structure of α-Fe + γ′-Fe4N. 29-35 nm particles of Fe-Ni-N ternary nitrides according to the formula γFe0.75Ni0.25N0.0026, γ-Fe0.5Ni0.5N0.002 and γ-Fe0.25Ni0.75N0.0012 were synthesized by sodium borohydride reduction of stoichiometrically mixed aqueous solutions of ferric nitrate and nickel chloride, followed by nitridation of the resulting Fe-Ni alloy precursors at 500oC for 4 h using NH3 (g) at ~120 cm3/min (Figure 6j).152 The nitrogen content in the nitrides were estimated by the Kjeldahl method using acetanilide as the source of nitrogen, and/or C, H, N analyzer. Fine particles of the interstitial ternary nitride Fe3Mo3N was synthesized by ammonolysis of FeMoO4 at 800oC for 12 h.153 The structure of Fe3Mo3N consists of NMo6 octahedra corner shared with the Fe atoms located at the sites between the octahedra. NPs of MMoN2 (M = Fe, Co, Fe0.8Mo0.2, Fe0.8W0.2) with WC-type layered structure were also synthesized by sol gel methods.154-157 FeMoN2 was synthesized by ammonolysis of Fe3Mo3N at 300oC for 6 h whereas (Fe0.8Mo0.2)MoN2 and (Fe0.8W0.2)WN2 were prepared from Fe2(MoO4)3 and Fe2(WO4)3 under NH3 (g) at 700oC for 96 h, respectively. Similarly FeWO4 was treated with NH3 (g) at 800oC for 12 h to obtain the FeWN2 NPs.158 In FeWN2, the Fe-N bonding is more ionic than the W-N bonding and hence leaching of Fe is possible by H2SO4 treatment to create Fe-deficient FexWN2 (x ~ 0.72, 0.74, 0.90) plates of 50 nm thickness.159, 160 Fe0.74WN2 was found to be FM at RT.160 (g) Composites: Bulk Materials and Thin Films The literature reported bulk composite phases were mostly due to the existence of undesired side phases along with the main nitride phase. For example, ε-Fe2-3N and γ′-


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Fe4N coexisted during plasma nitriding of Fe and ammonolysis of iron hydroxide at 600800oC.74,

84

DC magnetron sputtering with mixed Ar/N2 discharges gave composite

phases of ε-Fe3N - γ′-Fe4N with 10% N2 and the composite of FeN0.056 - α″-Fe16N2 at 5% N2.81 Composite layers of Fe + ε-Fe2+xN + ε-Fe2N were fabricated by irradiation of Fe/Si bilayers with 22 keV N2+ ions at RT.78 Plasma ion nitriding of α-Fe in the presence of N2/H2 mixture with 8 × 10-2 partial pressure of H2 under 0.4-1.3 Pa pressure resulted in the mixture of α-Fe + Fe4N + Fe16N2 + ε-Fe24N10.161 Composites of Li3FeN2 and Fe were obtained when Fe powder was microwave irradiated with slight excess of Li3N under N2 gas.162 However in none of the above techniques, the percentage composition of the phases could be controlled. (h) Composites: Nanostructures Chemical synthesis processes offer the advantage of controlling the composition and morphology of the individual nitride phases at the nanoscale, such that improved properties of the composite could be obtained over the individual components. The ~25 nm wide ε-Fe3N-GaN core-shell nanowires (Figure 6k, l) were prepared by citrate precursor technique.163, 164 The 2.5 cm long Fe-Ga-citrate fibers were grown by controlled evaporation of the citrate gel prepared from ferric nitrate and gallium nitrate according to the stoichiometric Fe/Ga atomic percentage ratio of 78/22. The citrate gel was decomposed to Fe-Ga-oxide nanowires at 550oC for 4 h which were then nitrided in flowing NH3 (g) at 750oC for 24 h to give the ε-Fe3N-GaN nanowires. ε-Fe3N and GaN with space groups P63/mmc and P63mc, respectively are immiscible hence facilitating the composite core-shell morphology. The GaN shell was ~4 nm and the ε-Fe3N/GaN mass ratio was 54/46. During the flow of NH3 (g) at 200 cm3/min, the Fe-Ga-oxide was 27 ACS Paragon Plus Environment


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reduced into the corresponding elements. However due to different melting points, Fe remained as solid whereas Ga in the liquid state. With the increase in temperature ε-Fe3N formation took place around 450oC whereas Ga in the liquid state encapsulated ε-Fe3N and later converted to GaN at ~700oC.164 With lower Fe/Ga atomic percentage ratios of 89/11 and 83/17, the ε-Fe3N/GaN mass ratios were 73/27 and 62/38, respectively without an observable core-shell morphology of the ~25 nm particles.134 For ε-Fe3N, the lattice parameters a (= 2.765-2.768 Å) and c (= 4.421-4.422 Å) did not alter systematically with Ga concentration, whereas lattice parameters of GaN were (a = 3.178 Å and c = 5.178 Å) were constant over all Fe/Ga concentrations. Composites of γ´-Fe4N, Fe4-xGaxN and GaN phases with three percentage compositions of Fe:Ga (70:30, 50:50 and 30:70) were obtained as strips comprising 30-40 nm spherical particles, synthesized from Fe- and Ganitrates and citric acid (Figure 6m, n).165 The basic difference in the synthesis protocol of γ´-Fe4N-Fe4-xGaxN-GaN to that of ε-Fe3N-GaN system was the lower NH3 (g) flow rate in the former. The oxide NPs were treated with NH3 at 700oC for 8 h with a flow rate of 14 cm3/min to obtain the composite of γ´-Fe4N, Fe4-xGaxN and GaN. ~25 nm ε-Fe3N-CrN composite system was prepared from Fe- and Cr-nitrates converted to a mixed metal citrate gel which was fired in air at 500oC for 4 h and then nitrided under 240 cm3/min NH3 flow at 710-830oC for 24 h.166 The Fe:Cr atomic percentage ratios of 86:14, 79:21 and 72:28 corresponded to ε-Fe3N:CrN, 68:32, 58:42 and 53:47 nanocomposite systems, respectively as determined by XRD-Rietveld and elemental analyses. When the Fe:Cr ratio was 72:28, ε-Fe3N-CrN nanorods of diameter ~30 nm and the length 400 nm were obtained.167 ε-Fe3N and CrN formed as separate phases since the diffusion of Cr into Fe was sluggish and hence kinetically an immiscible 28 ACS Paragon Plus Environment

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system was favored. When 50-80 atomic% Ni was incorporated into ε-Fe3N by the similar citrate precursor route, nanocomposites of hexagonal ε-Fe3-xNixN and fcc γ'-Fe4yNiyN

phases were formed.130 Ammonolysis of the oxide NPs were carried out at 400-

550oC for 12 h. The lattice parameters did not follow any particular trend with the increasing Ni concentration. 5. Magnetism of Iron Nitride Nanostructures The major interest in the magnetism of Fe-N compounds stems from finding an alternative to metallic iron especially with better stability. Fe NPs are pyrophoric and extremely reactive which limit their applications.168 Saturation magnetization (MS) of bulk iron is 222 emu/g,169 whereas the superparamagnetic relaxation and the oxide surface layers on Fe NPs reduce the moment. Fe NPs obtained from Fe(Co)x-oleylamine reacted precursor show MS of 160 and 192 emu/g for 2.3 and 10 nm particles, respectively.170 Additionally, the surfactant coated Fe NPs suffer loss in magnetization. Strongly attached surfactant molecules increase the coordination around the Fe atoms, which decreases its net moment by broadening the density of states.168 Also, the net magnetization from the surface atoms is zero since the surface atoms remain pinned at an orientation relative to the surface in the absence of directional bonding of the surface atoms. Thus the magnetization of Fe NPs suffers either way whether bare or coated by the surfactants. Although introducing nitrogen atoms in the α-Fe lattice dilutes the intrinsic magnetic moment, the stability of Fe-N nanostructures is comparatively better than Fe NPs. The Fe-N nanostructures have demonstrated diverse electronic and magnetic phenomena from RT down to low temperatures. The discussion is arranged in the decreasing order of magnetic moment i.e. α″-Fe16N2 followed by γ'-Fe4N and ε-Fe3N.



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The magnetic properties of doped, ternary and composite, ε-Fe3N and γ'-Fe4N nanostructures are also discussed. 5.1. α″-Fe16N2 α″-Fe16N2 has an average spin moment higher than that of BCC α-Fe (2.2 µB). Theoretically, stronger on-site Coulomb interactions between the localized 3d electrons in α″-Fe16N2 were observed by the Heyd-Scuseria-Ernzerhof screened hybrid functional method giving an average spin moment of 2.9 µB/Fe.171 The perturbative many-body correlations through the GW approximation and the on-site Coulomb correlations through the generalized gradient approximation method calculated 2.6-2.7 and 2.7 µB/Fe, respectively. In α″-Fe16N2 the nitrogen atoms sit at the center of a cluster of six Fe atoms and additional Fe atoms stay in between these clusters (Figure 2e). The electrons in the Fe atoms within the clusters tend to get localized and contribute to the enhanced magnetic moment. Although the exceedingly high MS of α″-Fe16N2 was first discovered in 1951 by Jack,172 followed by Kim and Takahashi in 1972 on 55 nm polycrystalline mixed phase Fe-N epitaxial film,173 and the data reproduced by Sugita and co-workers in 19911994,174, 175 the controversy surrounding the high magnetic moment of α″-Fe16N2 exists till today.176 Most of the magnetic data of α″-Fe16N2 reported till date are on mixed Fe-N films and bulk powders.35 MS of 5 and 55 nm epitaxial films were reported as 298 and 315 emu/g, respectively.91, 173 In the α″-Fe16N2 unit cell with stoichiometric nitrogen content, 16 Fe atoms remain distributed among three non-equivalent sites with relative site population ratio of 1:2:1. RT Mössbauer spectrum of single phase 100 nm diameter spherical particles of α″Fe16N2 showed three sextets due to differences in the distance of first (I), second (II) and 30 ACS Paragon Plus Environment

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third (III) nearest Fe and N atoms from the crystallographic symmetry (Figure 7a).121 The Fe I, II, and III sites had hyperfine fields of 29.8, 31.5, and 40.3 T, respectively with a relative intensity of 1:2:1. The isomer shifts and area% of I, II and III Fe sites were 0.048, 0.175, 0.154 mm/s and 25.6, 49.7, 24.7%, respectively. The surface oxide layer is detected from the presence of a paramagnetic doublet as was observed in 20 nm α″Fe16N2 particles with a 2 nm Fe-oxide shell (Figure 7b).122 At 4.2 K, the non-magnetic surface oxide layer got antiferromagnetically (AF) ordered evident from an additional sextet assigned to Fe3+ ions (Figure 7c). Because of the presence of this amorphous oxide layer, the overall MS was 108 emu/g and that of the FM core part was 196 emu/gFe. In general, the overall MS will decrease if the samples contain additional phases with lower magnetization than α″-Fe16N2, and the individual MS values can be estimated if the volume fraction of each component is known.124 In the 20 nm particles, the RT coercivity (HC) was 3350 Oe which increased rapidly below 25 K. The presence of a magnetically ordered oxide shell has the advantage of pinning the core FM spins at the core/shell interface and demonstrating the presence of exchange anisotropy through a shift of the hysteresis loop along the field axis. In this case, a hysteresis loop shift of 200 Oe was observed at 5 K. The magnetic uniaxial anisotropy was found to be 4.4 ×106 erg/cm3 from the magnetic torque measurement under 2.2 T field following the coherent rotation model.122 For the pure 100 nm α″-Fe16N2 particles,121 MS increased from 226 to 234 emu/g on decreasing the sample temperature from 300 to 5 K (Figure 7d). In this case the magnetocrystalline anisotropy energy constant Ku = HsatMS/2 was estimated to be 9.6 ×106 erg/cm3, where Hsat is the field corresponding to the maximum anisotropy field. When these NP compacts were sintered under high pressure,176 Ms increased as the



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sintering temperature (TS) increased up to 190oC, since sintering suppressed the surface oxidation (Figure 7e). Before air exposure, the RT HC was 2300 Oe and Ms was 177 emu/g. When the compacts were sintered up to 222oC, MS decreased due to a phase transformation from α″-Fe16N2 to ε-Fe2-3N and α-Fe. Thus optimized sintering conditions are required to lower the surface oxidation and prevent phase transformation. The decrease in HC with TS < 222oC was mainly due to the loss of shape anisotropy due to compaction and inter-particle magnetic interactions. The decrease of HC at 222oC was attributed to the phase transformation to the ε phase.176 An organic surfactant coating over the α″-Fe16N2 NPs can prevent surface oxidation and retain the small size of the NPs. Single phase NPs with stable magnetic order and high HC are essential requisites for applying the NPs to technologies of high density recording and magnetic hard disks,123 and there is still a lot of research scope in this direction. 5.2. γ′-Fe4N The chemical bonding of γ′-Fe4N is directly correlated to its FM properties. From Figure 2d, Fec and Fef positions can be referred to as Fe (I) and Fe (II), respectively. In the density of states curves (Figure 8a), three band regions for both majority and minority spin bands can be distinguished.14 In the lowest region, the N 2p states remain strongly hybridized through (pdσ)-type coupling with the d(3z2 – r2) orbitals of Fe (II) atoms. In the middle, all the d-bands are occupied in the majority spin bands and do not show any hybridization with the N p-orbitals. The Fermi level lies at the center of the d-band on the minority spin side. The anti-bonding p-d bands are situated at the top. Because of the strong hybridization between the N (p) and Fe (II) eg orbitals, a strong covalent Fe (II) – N bond exists leading to a decrease of the Fe (II) moment. However, the Fe (I) moment is 32 ACS Paragon Plus Environment

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almost unchanged. The Muffin-tin magnetic moments on Fe (I), Fe (II) and N are 2.73, 1.97 and -0.005 µB, respectively. Thus the magnetic moment of γ′-Fe4N is reduced to 2 µB as compared to ≈2.5 µB for FCC Fe in the high spin state.14 However, γ′-Fe4N is still a FM since the minority bands remain half-filled. Moreover, ab-initio calculations using the FP-LAPW codes showed that with increase in lattice constant, the magnetic moment of Fe (I) remains constant due to charge transfer between the d-orbitals of Fe (I) whereas the magnetic moment of Fe (II) increases with the lattice constant due to changes in the dxy, dxz, dyz and dx2-y2 orbitals.177 Conversion electron Mössbauer spectroscopy (CEMS) studies on a 36 nm thick γ′-Fe4N film capped with a 16 nm Cu3N layer and grown on (001) MgO substrate corroborated the above results and showed that the Fe (I) atoms have a local cubic structure and possess zero quadrupole splitting.44 For the two Fe (II)-A sites (Figure 2d), the internal magnetic field is oriented 90o with the principal axis of the electric field gradient tensor and for the Fe (II)-B sites, the corresponding angle is zero. Thus for the magnetically non-equivalent Fe (II) sites, the ratio of the quadrupole splitting of Fe (II)-A : Fe (II)-B is 1:-2. The spectrum in Figure 8b was deconvoluted into three magnetic components with intensity ratios of 3:4:1:1:4:3. ~3% paramagnetic component was used in the Lorentzian fitting to incorporate the contribution from the interface between γ′-Fe4N and the substrate. The cubic symmetry and easy [100] magnetization directions of the film were studied by the Kerr domain observations measured by transverse magnetooptic Kerr effect (MOKE) with the field applied in-plane and at different angles with respect to the crystalline axes.44 The remanence and coercive fields varied with the applied field directions. High-resolution vectorial Kerr measurements performed on nanometer thick γ′-Fe4N epitaxial thin films are particularly 33 ACS Paragon Plus Environment


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useful in determining the orientation of the easy and hard magnetization axes, additional anisotropy contributions and magnetization reversal mechanisms.178 In γ′-Fe4N NPs, the non-equivalent environments of Fe (II) may not be observed if the resolution of the spectra is low owing to large linewidths.179 Moreover, γ′-Fe4N NPs are likely to contain secondary phases such as ε-Fe3N and surface Fe-oxide phases, all of which can be identified by temperature dependent Mössbauer spectroscopy and by applying external magnetic field. Because of the presence of these additional phases and iron vacancies in the γ′-Fe4N lattice, MS will also reduce.179 If the impurity phases are amorphous, they could not be detected in the XRD patterns. In 70 nm γ′-Fe4N particles, 13-24% ε-Fe3N could be detected from the sextet with isomer shift 0.34 mm/s and quadrupole shift of 0.15 mm/s.34 ε-Fe3N gave rise to the fourth sextet in the Mössbauer spectrum recorded at 77 K, where the highest internal field of 34.2 T was attributed to the Fe (I) site of γ′-Fe4N (Figure 8c). A controlled flow of the nitriding gases can stabilize pure γ′-Fe4N powders and retain high MS. For example, single phase γ′-Fe4N obtained with 40-60% NH3 flow in a NH3/H2 gas mixture over Fe(CO)5 at 500oC for 3 h had MS of 188 ± 2 emu/g.84 With a lower NH3 flow, magnetization was higher due to presence of unnitrided Fe whereas with higher flow, ε-phases can coexist lowering the MS (Figure 8d). γ′-Fe4N always shows a lowering of MS by 17-18% as compared to α-Fe, due to volume expansion of 16% in γ′-Fe4N. Figure 8e shows the comparative hysteresis loops of 25 µm thick α-Fe and γ′-Fe4N sheets, in which the average magnetic moment per Fe atom were 2.19 and 2.14 µB, respectively.87 The additional factors for reduction in MS of γ′-Fe4N are the chemisorption of oxygen which can result in secondary oxynitride and/or


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oxide surface layer, canted spin structure at the surface of the NPs sometimes giving rise to a magnetically dead layer and presence of superparamagnetic fractions due to single domain particles within an ensemble of differently sized NPs.34 5.3. ε-Fe3N The RT Mössbauer spectrum of ε-Fe3N is usually fitted to a single sextet due to a single Fe site with equivalent coordination.42 The 22 nm ε-Fe3N particles had isomer shift δ = 0.35 mm/s, quadrupole splitting, ∆ = 0.03 mm/s, hyperfine field ∆Hf = 21.14 T, and outer linewidth of 0.55 mm/s (Figure 9a). The central superparamagnetic doublet had δ = 0.54 mm/s and ∆ = 0.08 mm/s. However, reducing the particle size to 16 nm presents only a single doublet without any hyperfine splitting. The above parameters can be explained taking into consideration the charge transfer processes in the formation of molecular orbitals (MO) with N 2p and Fe 3d orbitals.42 Fe has a metallic character in εFe3N due to decrease in electron density at the Fe nucleus since N 2p electrons occupy the bonding MO and shields the 4s electrons of Fe. ε-Fe3N has smaller s-electron density and/or greater p- and d-electron densities than γ′-Fe4N. The spin-down d electrons at Fe site and a small fraction of spin-up p electrons from N sites contribute to the density of states at the Fermi level (Ef) of ε-Fe3N. A major contribution from the p and d valence electrons leads to partial covalent character, larger Fe-N bond distance and improved symmetry at the Fe sites. The presence of doublet in a FM material can be corroborated to several factors. The first is the presence of smaller NPs than the average size which can give superparamagnetic fractions due to the presence of smaller particles.42 Secondly at RT, the spin relaxation time of superparamagnetic NPs is in the order of 10-11 to 10-12 s, shorter than the nuclear sensing time (10-8 s) and the sextet collapses into a singlet or 35 ACS Paragon Plus Environment


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doublet, which depends on the quadrupole splitting.180 At low enough temperatures, the relaxation time becomes comparable to the Larmor precession time and hence the sextet is observed. Third, if the Fe sites are coordinated with more number of N atoms due to higher nitrogen diffusion in the low dimensional particles, the doublet can be observed even at 77-80 K.113,

180

These doublet features disappear once the Mössbauer

spectroscopic studies are performed at temperatures such as 5 K, and also by applying external magnetic fields.113, 180, 181 When N occupies only the Wyckoff 2c sites in wellordered ε-Fe3N (Figure 2c), a single Fe-N coordination exists giving rise to a single sextet. Diffusion of extra nitrogen atoms leads to additional filling of nitrogen in Wyckoff 2b sites. In such circumstances, the distribution of hyperfine fields result from different types of Fe-N coordination and thus the magnetic ordering is reflected from two different sextet patterns (Figure 9b).113 In fact, the nitrogen stoichiometry (y) in ε-Fe3Ny is the key factor in maintaining the FM properties. For a perfectly stoichiometric ε-Fe3N, the magnetic moment per Fe had been estimated by band structure calculations to be 1.9 µB and the easy axis lies along the c-axis of the hexagonal structure (Figure 2c).182 On the contrary ε-Fe2.8N films had MS per Fe atom of 1.8 µB at 300 K.183 ~73 nm ε-Fe2.999N1.001 NPs were FM from 300 to 5 K.31 At 300 K with a negligible fraction of the superparamagnetic particles, MS was estimated to be 94.4 emu/g using the Langevin function whereas at 5 K, MS was 129 emu/g (Figure 9c), which closely matched the extrapolated MS of 133 emu/g at 0 K for bulk ε-Fe3N1.17.184 It is remarkable that the bulk magnetization could be stabilized in the NPs. Correspondingly HC increased from 72 Oe at 300 K to 159 Oe at 5 K. On the other hand, the ~20 nm ε-Fe3-yN1+y (0.1≤ y ≤ 0.2) NPs had a much lower MS of 52.8 emu/g at 5 36 ACS Paragon Plus Environment

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K (Figure 9c).31 Larger spin-pairing effects and covalent bonding between the N 2p and Fe 3d orbital electrons reduced the magnetization in ε-Fe3-yN1+y (0.1≤ y ≤ 0.2) with excess nitrogen. With exact nitrogen stoichiometry, larger number of free electron spins can ferromagnetically align along the hexagonal c-axis giving higher moment. The zero-field cooled (ZFC) and field cooled (FC) magnetization of both these samples showed strong inter-particle ‘exchange-like’ and dipolar interactions. The spins were blocked below the blocking temperature. AC susceptibility measurements however showed the coexistence of a low temperature spin glass phase due to freezing of the disordered surface spins and exchange interactions between the random spins at the interface of the nitride core and the oxide/oxynitride shell (Figure 9d, e). In ε-Fe3-yN1+y (0.1≤ y ≤ 0.2) NPs, the peak (Tf) shifts to higher temperatures due to increase in the time scale of measurement (t = 1/ω) with the increase in frequency (ω). Similar shifts in the in-phase (χ′) and out-of-phase (χ′′) components are signatures of both superparamagnetic blocking and spin glass phase. The two phenomena were distinguished by observing the relative shift (∆Tf) of the χ′ component in the ∆log10(f) frequency interval as Φ = ∆Tf / Tf∆log10(f). For ε-Fe3-yN1+y (0.1≤ y ≤ 0.2) and ε-Fe2.999N1.001 NPs, Φ is 0.004 and 0.002, respectively which fall within the range 0.005-0.05 of canonical spin glass.31 However it is always difficult to draw a distinct line between the spin glass phase and superparamagnetic blocking of spins. A similar situation of cooperative existence of blocked moments and freezing moments of interacting particles in a spin glass phase was observed in 10-40 nm ε-Fe3N NPs.106 In most of the Fe-N NPs, the surface Fe-oxide layers are mostly evidenced by Xray photoelectron spectroscopy (XPS),31, 106 and when the samples are field cooled, the AF or ferrimagnetic surface layers pin the ferromagnetically aligned spins at the core, 37 ACS Paragon Plus Environment


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leading to the exchange bias effect.106, 185 Also, the thin Fe-O/Fe-O-N surface layer over the FM Fe-N NPs consists of canted spins (Figure 9f), which freeze cooperatively below Tf. In doped FM systems, when the long-range FM ordering is broken by competing AF interactions at the core of the NP, re-entrant spin glass state results.185,

186

The low

temperature spin glass phase does not possess any long range magnetic order and below the freezing temperature, the spins freeze in a large number of possible metastable states.187 The random disorder is generally the result of competing random FM and AF interactions and the transition to the spin glass phase is denoted by a kink in the susceptibility versus temperature plot. The spin glass phase has been well established in frozen magnetic fluids of εFe3N with different particle number density, n.188-192 When n is reduced, MS and the dipolar interaction between the NPs tend to decrease. For example, MS of single domain ε-Fe3N NPs in frozen kerosene was reduced from 29.1 to 0.32 emu/cm3 by dilution.191 The spin glass phase is more prominent in the dense interacting NP systems than the diluted non-interacting NPs. The single domain non-interacting NPs are predominantly superparamagnetic although features such as a plateau in the temperature dependence of FC susceptibility and disappearance of the ZFC susceptibility at extremely low temperatures are similar to a spin glass phase (Figure 9g, h). The above two consequences in a spin glass state signify temperature independence of the equilibrium FC magnetization and existence of magnetic moments with short relaxation temperatures. Apart from AC susceptibility, experiments on magnetic relaxation can verify the spin glass phase. The relaxation rate depends on the wait time, the applied field and the temperature of measurement. Figure 9i shows the relaxation curves against inverse time


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for the dense ε-Fe3N fluid.191 The magnetization relaxations were measured in 20 Oe field at 50 K during ~105 s. The initial states SZFC, SFC18, SFC19, SFC20 and SFC21 represent the applied cooling fields of 0, 18, 19, 20 and 21 Oe. SHT20 state maintains a constant H/T ratio during cooling i.e. a cooling field of 20 × T/Tm, where Tm is the measurement temperature. For a spin glass phase, the FC magnetization remains nearly independent of temperature and is thus referred to as the equilibrium magnetization. Also, the onset of irreversibility between the ZFC and FC plots should be simultaneous with the peak of equilibrium susceptibility χ eq (T). In Figure 9i, all the relaxation curves converge at χ = M/Heff/MS = 6.6 × 10-3 Oe-1 at infinite time (t-1 → 0). This is also χ eq since the NPs will attain the equilibrium magnetization in infinite time at a certain temperature and field. Moreover, as observed from Figure 9h, χ eq is nearly same to that of the FC susceptibility which decreases slightly below 70 K. This is simultaneous with the ZFC-FC irreversibility and the χ ZFC (T) → 0, which was interpreted as the presence of magnetic moments with short relaxation times in the temperature range of the spin glass phase. Apart from cooperative freezing of the spins, random magnetic fields account for the broad peak of χ ZFC (Figure 9h).191 5.4. Doped, Ternary and Composite ε-Fe3N and γ′-Fe4N Nanostructures Doping metal ions at the Fe sites in ε-Fe3N and γ′-Fe4N leads to interesting magnetic ordering especially at low temperatures such as re-entrant spin glass and exchange bias phenomena. Lattice doping also induces magnetic dilution and randomization of spins breaking down of the long range FM order. This results in a strongly inhomogeneous magnetic ordering for example band structure calculations have



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shown that the AF and FM order in Mn doped ε-Fe3N depends on the nearest neighbor exchange interactions.129 There are only few reports available on the magnetism of doped Fe-N nanostructrues. Ni could be doped in the ε-Fe3N lattice only up to 40 atomic%.131, 185, 193, 194

yNiyN

Higher Ni doping results in the precipitation of a small fraction of FM γ'-Fe4-

along with the major superparamagnetic ε-Fe3-xNixN phase.195 In γ'-Fe4-yNiyN, the

added Ni atoms preferentially substitute the Fe (I) sites, where the isomer shift, quadrupole splitting and hyperfine field vary with the Ni content. In 30 ± 3 nm εFe2.6Ni0.4N particles, a sharp peak at 26.1 K in the ZFC plot was attributed to breaking of the translational symmetry at the 8 nm thick surface spin-glass layer.185 Re-entrant spin glass phase was only observed with 20-40 atomic% Ni from the kink or sharp peak in the FC magnetization plot. The spin glass phases were confirmed from the thermoremnant magnetization and critical slowing down analysis by dynamic scaling based on the peaks obtained from the frequency dependent AC susceptibility plots.185 Spin glass can be verified from the fitting parameters to the equation τ = τ0[Tf/(Tc-Tf)zν] where Tf is the frequency dependent freezing temperature, Tc is the critical temperature of the transition (Tc = Tf when ω → 0) and τ0 is the spin-flip time of atomic moments.185 For ε-Fe2.6Ni0.4N NPs, the fitted parameters τ0 = 10-8 s and zν = 4.5 confirmed the existence of a spin glass phase (Figure 10a). Cobalt doing in ε-Fe3N could be obtained till high atomic% and in εFe2.4Co0.6N NPs, the spin freezing process was speculated from the flattening of FC magnetization below 25 K.131 With chromium as the dopant, NPs of a single phase ε-Fe2.8Cr0.2N could be obtained whereas higher Cr concentrations give ε-Fe3N-CrN nanocomposite systems.166 Cr doping can reduce the magnetization below 15 emu/g since AF CrN stays randomly 40 ACS Paragon Plus Environment

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distributed within the FM ε-Fe3N matrix. The lowering of magnetization was additionally due to the presence of AF CrN/Cr/Cr-O-N/Fe-O-N phases on the NP surface, localized spin-disorder at the AF-FM interface, spin-pairing and finite-size effects. With 47 wt% CrN, nanorods were obtained and the effective anisotropy constant (Keff = 1/8 [Area of the hysteresis curve]) was determined to be ~6.3 × 103 ergs/cm3 at 10 K. The contribution to Keff was from the magnetocrystalline, shape and exchange anisotropy due to AF-FM interfacial coupling both at the bulk between ε-Fe3N and CrN and at the surface between the FM core and the AF surface layer.166, 167 The exchange bias after 3 T field cooling was however low

Iron Nitride Family at Reduced Dimensions: A Review of Their

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