Review Article pubs.acs.org/JPCC
Iron Nitride Family at Reduced Dimensions: A Review of Their Synthesis Protocols and Structural and Magnetic Properties Sayan Bhattacharyya* Department of Chemical Sciences, Indian Institute of Science Education and Research (IISER) Kolkata, Mohanpur 741246, India 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, and 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. 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), facecentered 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 in catalysis depending on the availability of Nactive sites,3−5 in biomedical fields since the iron nitrides are relatively less cytotoxic than the iron oxides, and as high wearresistant and corrosion-resistant coatings.6,7 As compared to other iron compounds, the nitrides possess a greater degree of magnetization than the iron oxides and are more cost-effective than ferromagnetic (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, corrosionresistant, 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 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, © 2014 American Chemical Society
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 350−550 °C. 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 pseudobinary 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 and γ′-Fe4N. The radius ratio, rx/rm, of the nonmetal (x) and metal (m) atoms plays an important role in the formation of these interstitial alloys.13 Below a critical Received: October 22, 2014 Revised: November 24, 2014 Published: December 19, 2014 1601
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The Journal of Physical Chemistry C magnitude of rx/rm = 0.59, simple structures such as cubic close-packed (CCP), BCC, or HCP are formed and 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− nonmetal 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 depend on the type of 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 nonmetal 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, ε-Fe3N1+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 combustion analysis which involves flash combustion 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.15−19 With the help of X-ray diffraction (XRD) studies, the first Fe−N phase diagram for temperatures 300−800 °C 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 aging 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 shows the largest homogeneity range.27 Several intermediate phases can also exist. For example, ε-
Figure 1. Fe−N phase diagram based on: (a) results of X-ray diffraction studies. Reprinted with permission.23 Copyright 1952, Wiley-VCH Verlag. (b) Data on magnetic transformations. Adapted with kind permission.24 Copyright 1987, Springer Science and Business Media. (c) Nitridation at atmospheric pressure. Adapted with kind permission.25 Copyright 2002, Springer Science and Business Media.
Fe24N10 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 Jack20,21 and are illustrated in Figure 2. The γ″-FeNy (y = 0.9− 1.0) phase has a cubic zinc-blende type structure (space group F43̅ m) with lattice parameter a = 4.33 Å.28 First-principles TBLMTO-ASA (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 1602
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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 two octahedral interstices occupied by the N atoms. There are three Fe positions: (0, 1/2, 1/4), (0, 0, 0.31), and (1/4, 1/4, 0) with site symmetries 4d(4m ̅ 2), 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 (ΔHf0 in kJ/mol) of the nitrides Table 1. Lattice Parameters of Iron Nitrides with Different Morphologiesa Figure 2. Crystal structures: (a) γ″-FeN (black spheres, Fe; gray spheres, N). Adapted with permission.29 Copyright 2000, Institute of Physics. (b) ς-Fe2N. Reprinted with permission.37 Copyright 2005, Elsevier. (c) ε-Fe3N. Adapted with permission.57 Copyright 2004, Elsevier. (d) γ′-Fe4N. Adapted with permission.44 Copyright 2004, American Physical Society. (e) α″-Fe16N2. Adapted with permission.1 Copyright 1999, Elsevier.
sample
structures,29 although only the γ″-FeN phase was reported experimentally.30 In the ZnS-type structure the nitrogen atoms occupy the tetrahedral interstices, whereas the NaCl-type structure has the lowest energy at the theoretical equilibrium volume and the N atoms occupy the middle of the edges of the 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, and the nitrogen atoms are arranged in zigzag chains parallel to the orthohexagonal b axis.10 The εFe3N 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 the spacer along the c-axis in the Fe lattice. The vacant sites are orderly and alternately filled with N atoms. An ideal hexagonal structure of ε-Fe3N has a 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 transform 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 the FCC Fe lattice with N atoms at the center of the unit cell. There are two equivalent Fe sites, one occupying the corners (Fe I) and the other at FCC (Fe II) positions of the cube. Due to the introduction of N atoms in the FCC γ-Fe
a
1603
morphology
space group
FeN36 γ″-FeN0.9110 Fe2N37
thin film thin film powder
----oP12
ς-Fe2N10
powder
---
ε-Fe3N1.47(1)38
P312
ε-Fe3N1.3310 ε-Fe3N1.3010 ε-Fe3N1.2210 ε-Fe3N1.1010 ε-Fe3N39
microcrystalline powder powder powder powder powder single crystal
--------P312
ε-Fe3N10 ε-Fe3N40
powder powder (336 °C)
--P6322
ε-Fe3N40
powder (25 °C)
P6322
ε-Fe3N40
powder (22 °C)
P6322
ε-Fe3N40
P6322
ε-Fe3N31
powder (−264 °C) NPs
ε-Fe3N41
NPs
ε-Fe3N42
NPs
ε-Fe3N43 ε-Fe3N43 γ′-Fe4N44 γ′-Fe4N40 γ′-Fe4N10 γ′-Fe4N45 α″-Fe16N210 α″-Fe16N221
NPs (450 °C) NPs (900 °C) thin film single crystal powder NPs bulk bulk
P63/ mmc P63/ mmc P63/ mmc P312 P312 --Pm3̅m ---------
unit cell parameters (Å) a = 4.53 a = 4.33 a = 4.423(4), b = 5.531(3), c = 4.821(3) a = 4.426, b = 5.529, c = 4.831 a = 4.8016(2), c = 4.4269(2) a = 4.774, c = 4.416 a = 4.760, c = 4.413 a = 4.743, c = 4.402 a = 4.718, c = 4.388 a = 4.7241(2), c = 4.3862(2) a = 4.696, c = 4.370 a = 4.7209(6), c = 4.4188(9) a = 4.7080(6), c = 4.3885(9) a = 4.6982(3), c = 4.3789(4) a = 4.6919(4), c = 4.3670(4) a = 2.695, c = 4.362 a = 2.75, c = 4.40 a = 2.70, c = 4.39 a a a a a a a a
= = = = = = = =
4.691, c = 4.367 4.716, c = 4.394 3.79 3.7900(6) 3.799 3.787 5.710, c = 6.283 5.72, c = 6.29
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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 the 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 were measured.47 Figure 3b shows the gasphase composition changes at 450 °C. As expected in the catalytic ammonia decomposition, the amount of N2 is 3-fold less 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 407 °C after eight cycles of 36 h, and ε-Fe3N1.39 was formed at 367 °C 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 εFe3N1.39: 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 357 °C lower than the ε-Fe3N nitriding temperature of 550 °C, a 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 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 longrange 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 well-ordered ε-Fe3N.
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 generate less negative ΔHf0: 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 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 ≥400 °C. 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 + 6H → N2 + 3H2). 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 the 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
Figure 3. (a) TG curves for nitriding process of nanocrystalline Fe at different indicated temperatures. (b) Gas-phase composition changes occurring during the nitridation process at 450 °C. Concentrations of (i) H2 in nitridation process (nitriding + NH3 decomposition), (ii) H2 in only nitriding reaction, (iii) N2 in catalytic NH3 decomposition, (iv) NH3 in the nitridation process. Adapted with permission.47 Copyright 2009, American Chemical Society. 1604
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Figure 4. (a) ε-Fe3N crystal structure with aε = 31/2aHCP and cε = cHCP superstructure cell in the P6322 space group symmetry showing the Fe atoms with HCP-type arrangement with AB stacking along [001] and the octahedral sites (o.s.). Enlarged view of the (b) edge-sharing and (c) face-sharing pairs of Fe6 octahedra, where the nitrogen motion pathways between o.s. perpendicular and parallel to [001] are indicated by arrows, respectively. (d) The o.s. of (a) are labeled as per the Wyckoff 2c, 2b, and 2d sites. Adapted with permission.58 Copyright 2007, Elsevier. (e) Reversible temperature-dependent N occupancy of the 2b Wyckoff sites obtained from neutron diffraction data. Reprinted with permission.56 Copyright 1999, Elsevier.
Annealing the sample at 100−135 °C reordered the N superstructure (Figure 4e), and this transition follows a firstorder 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 and 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 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 ∼1327 °C, mostly due to segregation of N-poor phases at the grain boundaries or formation of elemental iron.61 Also ε-Fe3N particles annealed at 400 °C partly converted to γ′-Fe4N with enrichment of nitrogen in the remaining ε-phase.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 the 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 the 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 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 140 °C, 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, and composite phases. For each of these systems, the discussion on the fabrication 1605
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Figure 5. Schematics showing the synthesis procedures of binary, doped, ternary, and composite iron nitride nanostructures.
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. 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 400 °C.68 ε-Fe3N films were formed either by magnetron sputtering on an Fe target in N2 and Ar flow,69 by 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 600−800 °C.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 ionassisted sputter deposition technique, phase evolution (εFex(≈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 450 °C (723 K) when Li3N was used as the nitride ion source, Al as the counter electrode, and Li−Al 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 discs 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 500 °C 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 + εFe2+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 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 500 °C 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 a 0.75 mm thick Fe sheet in the presence of the 1.3:98.7 NH3:H2 mixture at 840 °C 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 depend on the temperature of the reaction, the composition of the gas mixture, and time of exposure of the metal precursor to the flowing gas. A gas mixture of NH3 + H2 at 275−325 °C 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 600−800 °C in NH3(g).90 When Fe powder was converted to γ′-Fe4N particles at 500 °C with 40−60% NH3 in the NH3/H2 flow, 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 1606
DOI: 10.1021/jp510606z J. Phys. Chem. C 2015, 119, 1601−1622
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The Journal of Physical Chemistry C synthesize ternary, pseudobinary, and binary Fe−N phases. For example, ζ-Fe2N was transformed to single crystalline εFe3N1.47 with space group P6322 under 15(2) GPa pressure and 1327 °C using a resistivity-heated Walker-type two-stage multianvil device.38 The high pressures typically required are 10−15 GPa at temperatures of >1300 °C.37 Phase pure α″Fe16N2 is difficult to obtain in the form of both epitaxial films and powders.35 Relatively phase pure α″-Fe16N2 films were 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 aging α′-N-austenite yielded lesser pure α″-Fe16N2.93−97 The α″-Fe16N2 films with the 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 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 nm and outer diameters 10−20 nm (Figure 6a).3 The CNT channels were filled by ferric nitrate, which after controlled drying at 140 °C in air were treated with NH3 at 450 °C 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 Fe2 N NPs dispersed on nitrogen-doped graphene oxide were synthesized by chemical impregnation of iron acetate and a N-containing precursor such as 1,10phenanthroline in ethanol to form the N-coordinated Fe complex, followed by drying, ball milling, and thermal treatment in NH3 at 800 °C 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 vapor condensation of the reaction product of Fe(CO)5 and NH3/ Ar(g) on a rotating chiller was used to synthesize 10−40 nm εFe3N NPs passivated with 4−6 nm amorphous Fe3O4 or αFeOOH 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: 6Fe(CO)5 + 2NH3 → 2Fe3N + 3CO↑ + 3H2↑. 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,
Figure 6. Transmission (TEM) and scanning electron microscopy (SEM) images of Fe−N nanostructures. TEM images of (a) FeN/ CNT. Adapted with permission.3 Copyright 2012, Elsevier. (b) Intricate chains of γ′-Fe4N/α-Fe composite nanoparticles. Inset shows the oxide coating on the nitride core. Adapted with permission.108 Copyright 2004, Elsevier. (c) SEM and (d) TEM image of sponge-like ε-Fe3N. Adapted with permission.109 Copyright 2011, Royal Society of Chemistry. (e) ε-Fe3N magnetic fluid synthesized from a solution with 200 g of Fe(CO)5 in 50.1 g of kerosene. Adapted with permission.111 Copyright 1993, Elsevier. (f) TEM image of ε-Fe3−yN1+y (0.1 ≤ y ≤ 0.2) nanoparticles. Adapted with permission.31 Copyright 2010, Elsevier. (g) SEM image of deposited Fe−N particles with a N/Fe ratio of 0.23. Adapted with permission.118 Copyright 2008, The Japan Society of Applied Physics. (h) TEM image of ε-Fe2.6Ni0.4N. Adapted with permission.130 Copyright 2006, Elsevier. (i) SEM image of εFe2.8Cr0.2N nanoparticles. Adapted with kind permission.133 Copyright 2008, Springer Science and Business Media. (j) γ′-Fe0.75Ni0.25-N nanoparticles. Adapted with permission.152 Copyright 1999, American Institute of Physics. (k) TEM and (l) SEM images of ε-Fe3N-GaN core−shell nanowires. Adapted with permission.163 Copyright 2005, Institute of Physics. FESEM image of the 70:30 (Fe:Ga) nanocomposite γ′-Fe4N/Fe4−xGaxN sample shows (m1) strip-like morphology at low magnification and (m2) individual spherical nanoparticles within the strips. (n) TEM image of the γ′-Fe4N/Fe4−xGaxN particles. Adapted with permission.165 Copyright 2010, American Chemical Society.
respectively.108 Sponge-like ε-Fe3N having surface area of 39 m2/g was synthesized by a 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 700 °C 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-Fe3Nbased magnetic fluids are usually synthesized from Fe(CO)5 (Figure 6e).110,111 In one process, Fe(CO)5 vapor was mixed with Ar and NH3 gases and passed through a porous plate into 1607
DOI: 10.1021/jp510606z J. Phys. Chem. C 2015, 119, 1601−1622
Review Article
The Journal of Physical Chemistry C
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 240 °C.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 600 °C.34 Nanocrystalline iron was nitrided at 500 °C to a mixture of αFe and γ′-Fe4N, where the bigger iron crystallites could be converted to γ′-Fe4N at lower nitriding potential and vice versa.115,116 A solvothermal route was adapted for the synthesis of submicron-sized γ′-Fe4N particles with Fe3C coating.117 The synthesis was carried out at 400 °C 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 a N/Fe 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 150 °C inside an evacuated chamber. ε-Fe2−3N + 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 440 °C and later reacted with a flowing mixture of NH3 + H2 at 350 °C.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 nonaqueous benzyl alcohol in an autoclave at 200 °C for 48 h. The ∼40 nm oxide particles were reduced to 100−300 nm α-Fe particles in H2 flow at 400 °C for 10 h and nitrided in NH3 at 160 °C for 15 h to give NPs of α″-Fe16N2. Smaller particle size of the oxides resulted in lower reduction temperatures to nanosize α-Fe 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 500 °C to 120 nm α-Fe particles which were nitrided in NH3 flow at 130−170 °C 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
