Synthesis and Redox Properties of Iron and Iron Oxide Nanoparticles Obtained by Exsolution from Perovskite Ferrites Promoted by Auxiliary Reactions

Abstract

Nanoparticles of iron and iron oxides, as well as their composites, are of great scientific and technological interest. However, their properties and sustainability strongly depend on the preparation methods. Here, we present an original approach to synthesizing Fe and FeNi_x metal nanoparticles by exsolution, in a reducing environment at elevated temperatures from perovskite ferrites (La_1−xCa_xFeO_3−γ, CaFeO_2.5, etc.). This approach is made possible by the auxiliary reactions of non-reducible A-site cations (in ABO₃ notation) with the constituents of reducing compounds (h-BN etc.). The nanoparticles exsolved by our process are embedded in oxide matrices in individual voids formed in situ. They readily undergo redox cycling at moderate temperatures, while maintaining their localization. Fe nanoparticles can be obtained initially and after redox cycling in the high-temperature γ-form at temperatures below equilibrium. Using their redox properties, a new route to producing hollow and layered oxide magnetic nanoparticles (Fe₃O₄, Fe₃O₄/La_1−xCa_xFeO_3−γ), by separating the oxidized exsolved particles, was developed. Our approach provides greater flexibility in controlling exsolution reactions and matrix compositions, with a variety of possible starting compounds and exsolution degrees, from minimal up to ~100% (in some cases). The described strategy is highly important for the development of a wide range of new functional materials.

1. Introduction

Nanomaterials based on nanoparticles of transition metals and/or their oxides are of great interest from both scientific and practical viewpoints. Their applications are very diverse and include catalytic, magnetic, electronic, sorption, and pharmaceutical functional materials, utilizing various types of nanoparticles (supported, embedded, individual nanoparticles of various shapes) [1,2,3]. Many transition metal nanoparticles are chemically active and, therefore, suffer from a lack of stability, which is further exacerbated by their tendency for recrystallization, agglomeration, and growth, etc., especially at elevated temperatures. Supported/embedded nanocomposite materials have the potential to combine useful properties of both nanoparticles and matrices and maintain long-term stability and durability [1,3,4]. Moreover, those materials can develop new features due to the synergy between nanoparticles and the support, which is attributed to the intrinsic properties of the components and the strength of their interaction [4,5,6,7].

There are a wide variety of ways to produce these diverse nanoparticle-based nanomaterials [5,8,9]. One of the actively developing strategies for the synthesis of nanocomposite materials with stronger metal nanoparticle–substrate interactions (and, therefore, enhanced functionality and stability) is based on the thermal decomposition of complex oxides under reducing conditions. This kind of decomposition can be complete or partial and, especially the latter, is termed metal exsolution. In this approach, the complex oxides must contain both reducible and non-reducible metals under conditions of synthesis [10,11,12,13]. The main progress in this field is associated with the use of perovskite-related oxides of transition metals. It is a vast family of oxides with various stoichiometry (ABO₃, A_n+1B_nO_3n+1 Ruddlesden–Popper (RP) phases, etc.), whose crystal structures are based on [AO₃] close-packed layers, where the main constituents of A-sites are non-reducible rare-earth or alkaline-earth cations, while B-sites, in the interstices, can contain reducible cations of transition metals (Fe, Co, Ni, Cu, etc.) [14,15]. A series of Ln₂O₃/M nanocomposites, obtained by complete redox decomposition of the perovskite precursors of easily reducible transition metals (M = Ni, Co, etc.), were reported as metal nanoparticle-supported catalytic materials with enhanced productivity and stability [16,17]. However, the synthesis of nanocomposites by partial redox decomposition of perovskites requires methods to control the processes of metal nanoparticle exsolution and matrix transformations. Several approaches are proposed in the literature to facilitate metal exsolution in this manner [10,11,12,18]. One approach includes modification of the initial oxides, such as introducing an A-site deficiency [19,20], doping of B-sites with easily reduced cations [19,21,22], and promoting the formation of RP phases by adjusting the composition of A-sites [13], as well as others. Another approach is to use special reduction conditions, such as lattice strains, voltage biasing, and plasma assistance, as well as varying the temperature and pO₂ [13,18].

Metal reduction and exsolution is accompanied by changes in the parent perovskite matrices, which can occur in different ways. One such change is that the starting A-site-deficient perovskites are converted into stoichiometric ones, with minimal topotactical changes [21,23]. Another possible change is the formation of the corresponding oxides of non-reducible metals (La₂O₃, etc.) [17]. Yet another change is that the initial ABO₃-based perovskites can be transformed into RP phases, with different A to B ratios for some A-site compositions (containing Sr²⁺, etc.) [13].

Although the metal exsolution process is widely employed to enhance surface properties, it can also be used to modify the bulk properties of functional nanocomposites [24,25,26]. The exsolution at the interior grain and phase boundaries, together with the simultaneous transformation of host oxide matrices, provide a new pathway for the modification of a wide range of electrical, optical, and magnetic properties [25,26].

Overall, metal exsolution is a smart and effective strategy, with great potential for the preparation of new functional materials, with enhanced and unique properties. Its implementation, however, largely depends on the specific parent perovskite oxides and the appropriate reaction conditions [10,13,25]. Consequently, further development of methods for tuning the properties of the resulting nanocomposites, by controlling the exsolution process, is needed.

Herein, we present a novel approach to metal exsolution using perovskite-related ferrites, where B-metal exsolution is promoted by the additional auxiliary reactions involving non-reducible A-site cations. We have explored the behavior of a number of La and Ca-based ferrites in this new type of reaction, resulting in the formation of nanocomposites with nanoparticles embedded in individual voids, which are significantly different from the exsolved materials reported previously. The unique properties of the nanoparticles obtained in this manner, as well as the transformation of oxide matrices, will be discussed. It was shown that exsolved metal nanoparticles can be reversibly transformed into oxides in redox cycling, while maintaining their location inside the matrices. Moreover, the use of the chemical properties of nanocomposite matrices derived from alkaline earth-based ferrites makes it possible to separate and investigate the oxide nanoparticles obtained by the oxidation of exsolved metal nanoparticles.

2. Results and Discussion

2.1. The Main Concept of Metal Exsolution Promoted by Auxiliary Reactions

In this study, the synthesis of Fe-containing nanoparticles and nanocomposites is based on a specially developed process for metal exsolution from complex perovskite-related oxides of reducible transition metals (Fe, Ni, etc.). Our new strategy is to facilitate B-site reducible metal exsolution through additional auxiliary reactions (ARs) of non-reducible A-site (in the ABO₃ notation) cations with the constituents of reducing compounds, resulting in new complex oxides. This strategy, in a general simplified form, is described by schematic reaction (1):

ABO_3−γ + [Rⁿ] + {H₂} → B + AR¹_ηO_y + {H₂O} + {R²}

(1)

Here, [Rⁿ] is a reducing compound that reduces B cations, and at the same time, some (or all) of its constituents (R¹) form new complex oxides with non-reducible A-site cations that do not contain reducible B cations. In general, [Rⁿ] (R¹R², etc.) compounds can have a different chemical nature and aggregate states. For clarity of the explanation, the overall reaction can be nominally split into two consecutive pseudo-reactions, (2) and (3):

ABO_3−γ + [Rⁿ] + {H₂} → B + AO_y1 + ηR¹O_y2 + {H₂O} + {R²}

(2)

AO_y1 + ηR¹O_y2 → AR¹_ηO_y

(3)

Chemical reactions for the formation of complex oxides containing R¹ and non-reducible A-site cations (1b type) are designated here as auxiliary reactions. Curly brackets {} denote possible compounds. The reaction proceeds at elevated temperatures and under low pO₂ conditions, which are required for the formation and preservation of the B components in metal form. An important feature of reaction (1) is that the main source of oxygen for the formation of the resulting oxides is the initial oxides. This imposes limitations on the possible ratios of R¹ and A cations in the resulting complex oxides. If the initial oxides include two or more different B cations, then reaction (1) will be more complex depending on the B cation properties, as shown schematically below:

AB’_χB″_ψO_3−γ + [Rⁿ] + {H₂} → B’ + {B″} + AR¹_ηO_y + AB’_χ−ξB″_ψ+ϕO_3−β + {H₂O} + {R²}

(4)

Reaction (4) illustrates that all the different B cations can be exsolved to various degrees, or that the irreducible or less reducible ones can remain in perovskite oxides. The exsolved metals can also form alloys.

2.2. LaFeO₃ and Ca₂Fe₂O₅-Based Nanocomposites

For LaFeO₃ as a starting material and hexagonal BN as a reducing compound, the A-site cation AR (ACAR)-promoted exsolution can be written as reaction (5) (for clarity purposes, along with the nominal pseudo-reactions (6) and (7)):

LaFeO₃ + BN → Fe + LaBO₃ + 1/2N₂

(5)

{2LaFeO₃+ 2BN → 2Fe + La₂O₃+ B₂O₃ + N₂}

(6)

{La₂O₃ + B₂O₃ → 2LaBO₃}

(7)

Herein, the processes involving reaction (5) were carried out at elevated temperatures of 700–750 °C, in a reducing atmosphere, provided by a 10% H₂/Ar flow. Although hydrogen is not formally involved in the reaction, its role here is to maintain low pO₂ conditions, essential for the existence of metal Fe nanoparticles. The formation of lanthanum borate LaBO₃ provides an additional driving force for the reduction of LaFeO₃ and decreases the reaction temperature compared to a reduction with hydrogen. The specified temperatures of reaction (5) are below the decomposition temperature of LaFeO₃ in a hydrogen-containing atmosphere, which is reported to be >850 °C [27,28]. Thus, its decomposition, described by reaction (3), does not occur under these conditions.

LaFeO₃ + 3/2H₂ → Fe + 1/2La₂O₃ + 3/2H₂O

(8)

It is important to maintain reaction (5) as a main process. It should be noted that the higher temperature stability under reducing conditions makes ferrites more preferable starting materials compared to other perovskite-like compounds of reducible transition metals (cobaltites, nickelates, cuprates, etc.) [27].

Moreover, h-BN has a number of properties making it suitable for reactions of type (1, 5). First of all, it has decent reactivity in reactions of type (1, 4) with rare-earth (La, Y, etc.) and alkali-earth (Ca, Sr, etc.) perovskite-related ferrites. At the same time, it is exceptionally stable in inert and reducing atmospheres, and is also quite resistant to oxidation at moderate temperatures in air.

Figure 1 shows the powder XRD patterns of the LaFeO₃ ferrite-based samples after interaction with h-BN at ~750 °C with different amounts of exsolved Fe, depending on the LaFeO₃/h-BN ratio. As the reaction proceeds, the peaks corresponding to LaBO₃ and metal Fe appear in patterns, which are consistent with reaction (5) (Figure 1(a1)). The metal Fe peaks are distinguishable, but they strongly overlap with the LaBO₃ ones. The corresponding Mössbauer spectra are shown in Figure 1(b1,b2). Each of the spectra consist of two magnetically split sextet components and a paramagnetic singlet component. All the components are well resolved and do not broaden. According to their hyperfine parameters (Table 1), the first of two sextets with an isomer shift (δ) of 0.37 mm s⁻¹ and a hyperfine magnetic field (H_hf) of 52.4 T correspond to LaFeO₃ and the second, with δ ~0 mm s⁻¹ and H_hf = 33 T, correspond to αFe. The paramagnetic singlet, according to its hyperfine parameters (δ~−0.1 mm s⁻¹ at 298 K), corresponds to metal γFe. Moreover, γFe is a high-temperature paramagnetic form, with a close-packed fcc crystal structure. It is important to note that γFe is metastable at temperatures below 910 °C [29], while the synthesis temperature was ~750 °C. The metal γFe nanoparticles, however, may be undetectable in the XRD patterns (Figure 1(a2,b3)). The SEM images of LaFeO₃ after the ACAR exsolution display smooth surfaces, without distinguishable metal Fe nanoparticles distributed on them (Figure 2a). Figure 2b,c shows typical TEM images of such samples. For the samples at the low or moderate extent of reaction (Figure 2b,c), the images reveal agglomerated grains of different contrast, probably due to variations in thickness, but exsolved metal Fe particles are not clearly distinguishable. For the samples at the high extent of reaction >40%, obtained after prolonged heating for more than 15 h, the metal Fe particles are also largely undistinguishable, but some whisker-like metal Fe agglomerates are visible (Figure 2d) in small amounts with respect to the reaction extent. Additionally, the EDX analysis revealed numerous La-rich oxide grains, which, according to the XRD analysis, are actually LaBO₃, since boron is undetectable. Therefore, the metal Fe particles are embedded in the oxide matrix after the ACAR-promoted exsolution, which makes them hard to distinguish, except for the formation of whiskers, if any.

Figure 1. (a) Powder XRD patterns of the LaFeO₃/h-BN-derived nanocomposites: (a1) 29% of the total Fe amount exsolved, where 11% is in γ form; (a2) 54% of Fe exsolved (8% γFe); (a3) sample (a1) oxidized at 300 °C and subsequently reduced at 700 °C, all 29% of Fe exsolved is in α form (all the Fe contributions were evaluated by Mössbauer spectroscopy). (b) Corresponding Mössbauer spectra: (b1,b2) of sample (a1) at RT and 78 K, respectively; (b3) of sample (a2); (b4) of sample (a3) at RT.

Figure 1. (a) Powder XRD patterns of the LaFeO₃/h-BN-derived nanocomposites: (a1) 29% of the total Fe amount exsolved, where 11% is in γ form; (a2) 54% of Fe exsolved (8% γFe); (a3) sample (a1) oxidized at 300 °C and subsequently reduced at 700 °C, all 29% of Fe exsolved is in α form (all the Fe contributions were evaluated by Mössbauer spectroscopy). (b) Corresponding Mössbauer spectra: (b1,b2) of sample (a1) at RT and 78 K, respectively; (b3) of sample (a2); (b4) of sample (a3) at RT.

Table 1. Hyperfine parameters of Mössbauer spectra of the LaFeO₃/h-BN-derived nanocomposites measured at RT (except Figure 1(b2) at 78 K): δ—isomer shift, ΔEQ—quadrupole splitting, H—hyperfine magnetic field, A—relative area, Γ—linewidth.

Sample	Component	δ (mm s−1) ±0.01	ΔEQ (mm s−1) ±0.01	H (T) ±0.1	A (%) ±0.5	Γ (mm s−1) ±0.01	Comments
In Figure 1(b1)	s11	0.37	−0.07	52.4	71	0.26	Fe3+ oct. in LaFeO3
	s21	−0.01	−0.01	33.0	18	0.24	αFe
	d11	−0.10	0.00	-	11	0.25	γFe
In Figure 1(b2) 78 K	s12	0.48	−0.07	56.2	71	0.25	Fe3+ oct. in LaFeO3
	s22	0.11	0.01	33.8	18	0.25	αFe
	d12	0.01	0.00	-	11	0.25	γFe
In Figure 1(b3)	s13	0.37	−0.07	52.4	46	0.27	Fe3+ oct. in LaFeO3
	s23	0.00	0.00	33.0	45	0.24	αFe
	d13	−0.10	0.00	-	8	0.25	γFe
In Figure 1(b4)	s15	0.37	−0.07	52.3	72	0.25	Fe3+ oct. in LaFeO3
	s25	0.00	0.00	33.0	28	0.24	αFe
In Figure 3(1)	s123	0.37	−0.06	52.4	68	0.27	Fe3+ oct. in LaFeO3
	s223	0.00	0.01	33.1	23	0.23	αFe
	d123	−0.10	0.00	-	9	0.26	γFe
In Figure 3(2)	s124	0.37	−0.07	52.2	70	0.26	Fe3+ oct. in LaFeO3
	s224	−0.01	0.01	33.0	17	0.21	αFe
	d124	−0.10	0.00	-	13	0.23	γFe
In Figure 3(3)	s125	0.37	−0.07	52.2	71	0.25	Fe3+ oct. in LaFeO3
	s225	0.35	−0.17	51.0	5	0.33	Fe3+ oct. in αFe2O3
	s325	0.00	0.01	33.0	23	0.23	αFe
	d125	−0.10	0.00	-	1	0.26	γFe
In Figure 3(4)	s126	0.37	−0.07	52.2	69	0.26	Fe3+ oct. in LaFeO3
	s226	0.00	0.00	33.0	30	0.22	αFe
	d126	−0.12	0.00	-	1	0.25	γFe
In Figure 3(5)	s127	0.37	−0.08	52.3	70	0.25	Fe3+ oct. in LaFeO3
	s227	0.38	−0.18	51.4	30	0.25	Fe3+ oct. in αFe2O3
In Figure 3(6)	s128	0.37	−0.06	52.2	70	0.26	Fe3+ oct. in LaFeO3
	s228	0.00	0.00	33.0	14	0.21	αFe
	d128	−0.10	0.00	-	16	0.24	γFe

Table 1. Hyperfine parameters of Mössbauer spectra of the LaFeO₃/h-BN-derived nanocomposites measured at RT (except Figure 1(b2) at 78 K): δ—isomer shift, ΔEQ—quadrupole splitting, H—hyperfine magnetic field, A—relative area, Γ—linewidth.

Sample	Component	δ (mm s−1) ±0.01	ΔEQ (mm s−1) ±0.01	H (T) ±0.1	A (%) ±0.5	Γ (mm s−1) ±0.01	Comments
In Figure 1(b1)	s11	0.37	−0.07	52.4	71	0.26	Fe3+ oct. in LaFeO3
	s21	−0.01	−0.01	33.0	18	0.24	αFe
	d11	−0.10	0.00	-	11	0.25	γFe
In Figure 1(b2) 78 K	s12	0.48	−0.07	56.2	71	0.25	Fe3+ oct. in LaFeO3
	s22	0.11	0.01	33.8	18	0.25	αFe
	d12	0.01	0.00	-	11	0.25	γFe
In Figure 1(b3)	s13	0.37	−0.07	52.4	46	0.27	Fe3+ oct. in LaFeO3
	s23	0.00	0.00	33.0	45	0.24	αFe
	d13	−0.10	0.00	-	8	0.25	γFe
In Figure 1(b4)	s15	0.37	−0.07	52.3	72	0.25	Fe3+ oct. in LaFeO3
	s25	0.00	0.00	33.0	28	0.24	αFe
In Figure 3(1)	s123	0.37	−0.06	52.4	68	0.27	Fe3+ oct. in LaFeO3
	s223	0.00	0.01	33.1	23	0.23	αFe
	d123	−0.10	0.00	-	9	0.26	γFe
In Figure 3(2)	s124	0.37	−0.07	52.2	70	0.26	Fe3+ oct. in LaFeO3
	s224	−0.01	0.01	33.0	17	0.21	αFe
	d124	−0.10	0.00	-	13	0.23	γFe
In Figure 3(3)	s125	0.37	−0.07	52.2	71	0.25	Fe3+ oct. in LaFeO3
	s225	0.35	−0.17	51.0	5	0.33	Fe3+ oct. in αFe2O3
	s325	0.00	0.01	33.0	23	0.23	αFe
	d125	−0.10	0.00	-	1	0.26	γFe
In Figure 3(4)	s126	0.37	−0.07	52.2	69	0.26	Fe3+ oct. in LaFeO3
	s226	0.00	0.00	33.0	30	0.22	αFe
	d126	−0.12	0.00	-	1	0.25	γFe
In Figure 3(5)	s127	0.37	−0.08	52.3	70	0.25	Fe3+ oct. in LaFeO3
	s227	0.38	−0.18	51.4	30	0.25	Fe3+ oct. in αFe2O3
In Figure 3(6)	s128	0.37	−0.06	52.2	70	0.26	Fe3+ oct. in LaFeO3
	s228	0.00	0.00	33.0	14	0.21	αFe
	d128	−0.10	0.00	-	16	0.24	γFe

Figure 2. SEM and TEM images of the LaFeO₃/h-BN-derived nanocomposites: (a,b) SEM and bright field (BF) TEM images of the nanocomposites with ~45% of Fe total exsolved; (c,d) high-angle annular dark-field scanning TEM (HAADF-STEM) image and energy-dispersive X-ray (EDX) elemental analysis in selected locations of the nanocomposites with ~12% of Fe total (~10% γFe); (d) formation of the Fe whisker (w) in sample (a,b).

Figure 2. SEM and TEM images of the LaFeO₃/h-BN-derived nanocomposites: (a,b) SEM and bright field (BF) TEM images of the nanocomposites with ~45% of Fe total exsolved; (c,d) high-angle annular dark-field scanning TEM (HAADF-STEM) image and energy-dispersive X-ray (EDX) elemental analysis in selected locations of the nanocomposites with ~12% of Fe total (~10% γFe); (d) formation of the Fe whisker (w) in sample (a,b).

Figure 3. RT Mössbauer spectra of the LaFeO₃/h-BN-derived nanocomposites during redox cycling: (1) as prepared with 32% of Fe exsolved (~9% γFe); (2) sample (1) after consecutive oxidation (air, 500 °C) and reduction (10% H₂/Ar, 700 °C) (~30% of Fe total, ~13% γFe); (3) sample (2) after oxidation at 300 °C in air (~23% αFe, ~1% γFe, ~5% αFe₂O₃); (4) after reduction of sample (3) (10% H₂/Ar, 700 °C) (~30% αFe); (5) after oxidation of sample (4) (air, 500 °C) (~30% αFe₂O₃); (6) after reduction (10% H₂/Ar, 700 °C) of sample (5) (~30% of Fe total, ~16% γFe).

Figure 3. RT Mössbauer spectra of the LaFeO₃/h-BN-derived nanocomposites during redox cycling: (1) as prepared with 32% of Fe exsolved (~9% γFe); (2) sample (1) after consecutive oxidation (air, 500 °C) and reduction (10% H₂/Ar, 700 °C) (~30% of Fe total, ~13% γFe); (3) sample (2) after oxidation at 300 °C in air (~23% αFe, ~1% γFe, ~5% αFe₂O₃); (4) after reduction of sample (3) (10% H₂/Ar, 700 °C) (~30% αFe); (5) after oxidation of sample (4) (air, 500 °C) (~30% αFe₂O₃); (6) after reduction (10% H₂/Ar, 700 °C) of sample (5) (~30% of Fe total, ~16% γFe).

This is rather different from the metal exsolution from the A-site-deficient perovskites or the high-temperature hydrogen reduction of stoichiometric perovskite oxides reported previously, where the exsolved metal nanoparticles are typically clearly visible on the surfaces of oxide matrixes [10,11,12,13]. In these processes, the reducible metal cations migrate toward the outer and inner grain surfaces, which are exposed to the reducing environment, and reduce to metal, with subsequent agglomeration, grain growth, etc. The oxide matrices, in turn, undergo shrinkage because of a loss of oxygen and metal constituents, promoting socketing of the exsolved nanoparticles [5,10,12,23]. In our ACAR-promoted exsolution process, metal reduction/exsolution is accompanied by the formation of other complex oxides of two or more elements instead of the initial perovskite oxide, even those of lower density, i.e., LaBO₃ vs. LaFeO₃. Their formation begins in the contact areas between the initial ferrite grains and the reducing reagents. These newly formed Fe-free oxides (viz. LaBO₃) on the top of the initial perovskite grains presumably create a diffusion barrier for the Fe species. Consequently, the exsolved metal Fe nanoparticles do not appear on top of the oxide grain surfaces, but instead localize underneath their surfaces in generated in situ voids.

Perovskite-like ferrites of alkaline-earth elements, such as Ca₂Fe₂O₅ and Sr₂Fe₂O₅, can also be used in reaction (1, 4) to produce metal nanoparticle-bearing composites. Because Ca is lighter than La and Sr, it may be more favorable for microscopic investigations in terms of the nanoparticle–matrix contrast. As stated above, the starting ferrite is the main source of oxygen for the formation of the resulting oxides in ACAR-promoted exsolution processes involving reaction (1, 4). Consequently, in the case of La ferrite, other La borates, such as LaB₃O₆, cannot be formed from the LaFeO₃ precursor due to a lack of oxygen. For Ca₂Fe₂O₅ and h-BN as starting materials, the formation of several Ca borates is possible, viz. Ca₂B₂O₅ and Ca₃B₂O₆ (without/with CaB₂O₄). The powder XRD patterns of Ca₂Fe₂O₅ after the ACAR-promoted exsolution reactions with h-BN and the corresponding Mössbauer spectra are shown in Figure 4. According to Figure 4(a2), the main complex oxide product is Ca₃B₂O₆ and, therefore, the primary reaction can be written as follow:

3Ca₂Fe₂O₅ + 4BN + 3H₂ → 6Fe + 2Ca₃B₂O₆ + 2N₂ + 3H₂O

(9)

Hydrogen, in addition to maintaining the low pO₂, reacts with excessive oxygen in this case. It is worth noting that no diffraction peaks corresponding to Ca₃B₂O₆ and/or other possible resulting oxides were observed at the low extent of reaction, indicating that the oxide products were X-ray amorphous (Figure 4(a1)). Ca₂Fe₂O₅ is more active in this type of reaction compared to LaFeO₃ and is less stable in reducing conditions as well, so minor additional peaks of CaO can be observed in the patterns of some samples depending on the preparation and reduction conditions (Figure 4(a3)). The Mössbauer spectra of the Ca₂Fe₂O₅-derived samples are shown in Figure 4b. The spectra are comprised of three magnetically split sextet components and a paramagnetic singlet component. The Ca₂Fe₂O₅ (=CaFeO_2.5) has a brownmillerite structure, which is oxygen deficient compared to ABO₃ perovskites with fully ordered oxygen vacancies at room temperature, where Fe³⁺ cations equally occupy distorted octahedral and tetrahedral oxygen polyhedra. Consequently, two of the sextets of nearly equal spectral contributions, according to their hyperfine parameters, correspond to the brownmillerite subspectrum, i.e., Fe³⁺tet (δ~0.19 mm s⁻¹ and H_hf = 43.4 T) and Fe³⁺oct (δ~0.37 mm s⁻¹ and H_hf = 51.2 T) in Ca₂Fe₂O₅. The third sextet (δ~0 mm s⁻¹ and H_hf = 33 T) and a singlet (δ~−0.1 mm s⁻¹) correspond to metal Fe in α and γ forms, respectively, i.e., to the metal Fe subspectrum (Table S2). Like in the LaFeO₃ case, at ~650–730 °C, the temperatures during the ACAR-promoted exsolution process were much lower than 910 °C. Both the Fe forms can be distinguished in the powder XRD patterns (Figure 4a). Similar to the previous case, the SEM images of the Ca₂Fe₂O₅-derived samples display smooth surfaces, without distinguishable metal Fe particles (Figure 5a). The TEM images (Figure 5b,c) show agglomerates of different contrasts, without metal Fe particles being clearly visible on the surfaces. At the same time, a small number of Fe whiskers can be observed in some samples (Figure 5d). However, as it follows from Figure 5c, the EDX mapping in several locations shows areas of Fe segregation, which, taking into account the distributions of other elements, can be identified as Fe-embedded nanoparticles. These exsolved, embedded Fe nanoparticles, identified by the EDX analysis, were about 15–25 nm in size. The EDX analysis also displays Ca-rich O-containing agglomerates (i.e., Ca₃B₂O₆, etc.).

The data show that the Ca₂Fe₂O₅/h-BN system demonstrates similar behavior to the LaFeO₃/h-BN system in the exsolution reactions facilitated by ACARs. In both the systems, the exsolved metal Fe nanoparticles, which can be determined by EDX analysis, are mostly located within oxide matrices, consisting of initial ferrites and newly formed iron-free borates, in the generated in situ voids. Presumably, the reaction zones and, hence, the adjacent voids are somehow connected to the low pO₂ environment, otherwise, in most cases, the metal reduction reaction will not proceed to any significant extent. Metal Fe in the exsolved nanoparticles can exist in two forms, namely metastable at temperatures <910 °C γFe and stable αFe, the latter can also form whiskers in some cases. The interior particle growth in a confined space in voids appears to be a key factor, along with a nanosized dimension, for the γFe formation at temperatures well below 910 °C. The fcc close-packed crystal structure of γFe is denser than the bcc structure of α(δ)Fe [30,31], so that the compressive strain developed when the particles of nano-scale dimensions grow under confined conditions is conducive to the formation of γFe nanoparticles and their subsequent stabilization upon cooling. The strain-induced formation of γFe nanoparticles, smaller than ~20 Å, under confined conditions in oxide matrices (Al₂O₃, MgO), which are stable at an ambient temperature, has been reported previously [32,33]. The stabilization of γFe was also observed in iron coatings produced by arc plasma deposition on porous alumina substrates, when Fe was localized inside pores with a diameter <160 nm [34]. Note that in this case, the deposited Fe layers only covered the inner walls of the pores and did not completely fill their interior space, leaving central gaps.

The γFe/αFe ratio depends on the reaction conditions and extent: the contributions of γFe are larger in the initial stages and at the low extent of reactions. However, the nanocomposites containing only γFe have not been obtained using individual Ca₂Fe₂O₅ or LaFeO₃ ferrites. Ca₂Fe₂O₅ is more active during the described process than LaFeO₃ and reacts at lower temperatures, and the reaction can proceed almost completely, whereas for LaFeO₃, its extent is limited to ~50–60% (based on the Fe content in coexisting phases obtained from the Mössbauer spectra).

2.3. ACAR-Promoted Metal Exsolution Using the Substituted Ferrites

The properties of Ca₂Fe₂O₅ or LaFeO₃ ferrites can be significantly modified by substitutions. Accordingly, the effect of cation substitution on ACAR-promoted metal exsolution and γFe formation was investigated. Rare-earth and alkaline-earth ABO_3−γ perovskite ferrites allow substitution in both A- and B-sites and a wide variation in oxygen content of 0 ≤ γ ≤ 0.5 [14,15]. Figure 6b shows the powder XRD patterns of A and B double-substituted compounds, Ca_1.4Y_0.6Fe_1.8Zn_0.2O_5.2 (=Ca_0.7Y_0.3Fe_0.9Zn_0.1O_2.6), after ACAR-promoted exsolution using h-BN. Y was chosen as the lightest rare-earth 3+ cation. The Ca_1-xY_xFeO_3−γ solid solutions have not been studied in detail in the literature, but for our process it is important that they belong to a pseudobinary system, i.e., no additional phases other than the perovskite solid solutions coexist. The solubility ranges from both the Ca side (the brownmillerite type solid solutions with oxygen excess) and the Y side (the LaFeO_3−γ type ones with mainly disordered oxygen vacancies) are not well defined and, presumably, depend on the temperature [35]. Similar to the case of unsubstituted Ca ferrite, the resulting borates can only be correctly identified at relatively large reaction extents, such as in the sample in Figure 6(b2), with ~29% of Fe exsolved, according to the Mössbauer spectrum shown in Figure 6(a2). The XRD phase analysis revealed that substitution with Y increases the number of oxide products formed, and the ACAR-promoted exsolution process in this case can be written in a simplified form as follows:

Ca_0.7Y_0.3Fe_0.9Zn_0.1O_2.6 + BN + H₂ → Fe + Ca₃B₂O₆ + Y₂O₃ + Ca_0.7−zY_0.3+zFe_0.8−βZn_0.1+βO_2.6+γ + {Zn₅B₄O₁₁} + N₂ + H₂O

(10)

The main oxide products are Ca₃B₂O₆ and Y₂O₃ (Zn₅B₄O₁₁ could be a minor phase). The Mössbauer spectra (Figure 6(a2,a3)) show that 29% of the Fe total content was exsolved in this sample, where γFe accounted for ~9% and αFe for 20%. The substitution manifests itself in a significant broadening of lines in the Mössbauer spectra, due to the introduction of local distortions and the disruption of the magnetic superexchange interactions. Consequently, the components of the brownmillerite subspectrum that correspond to the tetrahedral and octahedral plus pentagonal positions in the ferrite structures were fitted with combinations of several Zeeman sextets. Sextets with δ~0.34–0.37 mm s⁻¹ and H_hf~49–51 T correspond to Fe³⁺ cations in octahedral and pentagonal positions (poorly resolved at RT), while those with δ~0.18–0.20 mm s⁻¹ and H_hf ~40–43 T correspond to the tetrahedral ones. The two other components with narrow lines correspond to the metal Fe subspectrum, comprising of the αFe sextet with δ~0 mm s⁻¹ and H_hf ~33 T and the γFe singlet with δ~−0.1 mm s⁻¹. The Mössbauer measurements at 78 K result in the narrower lines of the components, but the spectral contributions of the subspectra remain about the same. At a lower exsolution level of ~3–6%, all metal Fe was in the γ form (Figure 6(a1), Table S3). The TEM images of the latter samples are shown in Figure 7. Similar to the unsubstituted Ca₂Fe₂O₅-derived samples, it was difficult to differentiate the Fe particles among the agglomerates with different contrasts. However, the Fe nanoparticles <50 nm in size embedded in the oxide matrix were clearly identified by the EDX analysis and mapping (Figure 7a,c). The HAADF-STEM images, together with the Fourier transform imaging, confirmed that the exsolved nanoparticles are metallic γFe (Figure 7b). The EDX mapping also shows some degree of Y and O segregation on the scale of tens of nanometers, which is consistent with the XRD phase analysis.

Since Ca is more active in the exsolution reactions than Y, the remaining Ca–Y ferrite can also be enriched in Y despite the formation of Y₂O₃. This can be seen from the spectra of the samples with high degrees of exsolution, as shown in Figure S1 and Table S4. They reveal that the contributions of spectral components corresponding to Fe³⁺ in tetrahedral coordination decrease significantly because of the transformation of Ca ferrite-based Ca_1−xY_xFeO_3−γ solid solutions with x = 0.3 to Y ferrite-based solid solutions (presumably with x > 0.6 [35]).

Preferable formation of Ca-rich borates was similarly observed for the La-substituted Ca_1−xLa_xFeO_3−γ starting ferrites. This system is also pseudobinary. Within this system, the Ca₂LaFe₃O₈ Grenier phase, with a crystal structure intermediate between brownmillerite and perovskite types, was reported. In addition, the formation of microdomains of close compositions is possible at different values of x and γ [36,37,38]. The main oxide product in ACAR-promoted exsolution reactions of Ca–La ferrites with h-BN is Ca₃B₂O₆, like for Ca_0.7La_0.3FeO_2.65 in Figure 8(a2,a3)), while the formation of La₂O₃ was not observed. At higher exsolution degrees, double borate Ca₃La₃(BO₃)₅ was additionally formed (Figure 8(a3)). Since La is retained more in the Ca_0.7-zLa_0.3+zFeO_2.65+δ solid solution, its content increases, as does the oxygen index value. At some level of La content, the resulting ferrite solid solution loses oxygen vacancy ordering, i.e., it becomes the La_0.3+zCa_0.7−zFeO_3−γ type perovskite-like solid solution. This is reflected in the Mössbauer spectra as the disappearance of Fe³⁺ tetrahedral spectral components (Figure 8(b2), Table S5). A simplified reaction is shown in (11):

Ca_0.7La_0.3FeO_2.65 + BN + H₂ → Fe + Ca₃B₂O₆ + La_0.3+zCa_0.7–zFeO_3−γ + {Ca₃La₃(BO₃)₅} + N₂ + H₂O

(11)

In the case of LaFeO₃-based compounds, the substitution of Fe³⁺ cations by non-reducible cations can also lead to the formation of γFe exclusively. Figure 9 shows that the Zn-substituted solid solution, LaFe_0.8Zn_0.2O_2.9, reacts with h-BN, according to the simplified reaction:

LaFe_0.8Zn_0.2O_2.9 + BN → Fe + LaBO₃ + LaFe_0.8−βZn_0.2+βO_3−γ + N₂ + {ZnO}

(12)

At ~8% of the total metal Fe content, all of the exsolved Fe is in the γ form (Figure 9(b1), Table S5). The TEM EDX mapping of this LaFe_0.8Zn_0.2O_2.9-derived sample with ~8% of γFe exsolved, allowed for the identification of the embedded Fe metal nanoparticles (Figure 10a). The HAADF-STEM images, together with the Fourier transform imaging, confirmed that the exsolved nanoparticles are highly twinned metal γFe nanocrystals (Figure 10b,c).

In the case of double substitution by Ca and Zn, the ACAR exsolution with h-BN will proceed as follows (Figure S2(a2,a3,b2,b3)):

La_0.8Ca_0.2Fe_0.8Zn_0.2O_2.8 + BN + H₂ → Fe + LaBO₃ + Ca₃La₃(BO₃)₅ + La_0.8+zCa_0.2−zFe_0.8−βZn_0.2+βO_3−γ + {ZnO} + N₂ + H₂O

(13)

At low reaction extents only γFe will be exsolved (Figure S2(b2)). In both cases, Zn-containing phases may precipitate at high reaction extents, presumably in oxide (ZnO) or borate forms, but in most of our samples they did not appear in the XRD patterns.

The lattice parameters of the cubic fcc cell of the exsolved γFe nanocrystals can be estimated from the room-temperature powder XRD patterns, in the range of 0.357–0.358 nm. These values correspond to those calculated for austenite solid solutions at room temperature when extrapolated to pure iron [39,40]. Since the γFe nanoparticles synthesized here at 650–750 °C are crystalline and are confined in oxide/borate matrices, they are not pyrophoric in air at room temperature. Moreover, they persist during heating/cooling cycles in a reducing atmosphere. However, a mild mechanical impact on the nanocomposites with high γFe content, by gently grinding in a mortar and pestle for 1–2 min, leads to the transformation of most of the γFe into αFe, and irreversibly so (Figure 9(a2,a3,b2,b3) and Figure S2(a3,a4,b3,b4), Table S6). This transition is, to some extent, analogous to transformations of retained austenite caused by mechanical deformation at low temperatures [41,42].

Note that for all the substituted ferrites investigated, the formation of Fe whiskers, which is considered undesirable during the described processes, was not observed.

2.4. Redox Behavior of the Exsolved Nanoparticles

The exsolved metal nanoparticles can be oxidized in air to iron oxides at elevated temperatures. The oxidation of metal nanoparticles is generally associated with a decrease in density and an increase in volume. As it was shown in [23], for CoNi exsolved, socketed nanoparticles, these volume changes lead to their migration from the initial sites under redox cycling conditions. In addition, redox cycling can cause the growth and coarsening of oxidized metal nanoparticles or their reintegration into perovskite matrices during oxidation [10,43,44]. For all the nanocomposites obtained herein, the embedded nanoparticles are also completely oxidized to αFe₂O₃ at 500 °C and above, in air. The γFe nanoparticles are oxidized first, and at lower temperatures of 200–300 °C, and their oxidation is accompanied by the transition of residual γFe into the α form (Figure 3(3,5), Table 1). However, the nanocomposites produced by ACAR-promoted exsolution exhibited peculiar behavior during redox cycling, or at least at the temperatures investigated. In these nanocomposites, the initial metal nanoparticles, after complete oxidation into αFe₂O₃ at 500 °C and above, can revert back to the metallic state by subsequent reduction in 10% H₂/Ar at 650–750 °C and, remarkably, the metal γFe nanoparticles can be reinstated in the γ form through such a reduction of αFe₂O₃. Moreover, the γFe fractions may increase compared to the initial content, especially when Ca₂Fe₂O₅ and its solid solutions have been used (Figure 3(1,2) and Figure 4(a3,a4,b3,b4)). To our knowledge, the formation of γFe nanoparticles from Fe oxides by hydrogen reduction at temperatures below 910 °C has not been reported previously. Like the formation of the initial γFe nanoparticles, this is presumably a compressive strain-driven behavior, which can be attributed to the preservation of the localization of all the nanoparticles, i.e., initially exsolved metallic-derived oxide and restored metallic oxide, within the voids, along with their strong binding to the void walls during the redox cycles. The γFe formation can, thus, be considered as a kind of indicator of the localization of nanoparticles inside the voids during these processes. Note that at the same time, there should also be free space inside the voids close to nanoparticles, sufficient to compensate for the metal/oxide volume difference, otherwise the matrix may be destroyed.

Using the above assumption on the nanoparticle redox behavior, the reversible transformation of γFe into αFe, and vice versa, through redox cycling at different oxidation temperatures can be realized (Figure 3, Table 1). The sample of the LaFeO₃/h-BN based nanocomposite with ~30% of Fe exsolved was subjected to a preliminary oxidation (at 500 °C)/reduction cycle to maximize the γFe contribution (Figure 3(1,2), Table 1 (d123, d124)). To convert most of the Fe into the α form, the first stage involved oxidation at a low temperature of ~300 °C, where more chemically active γFe nanoparticles were partially oxidized/partially converted into αFe (Figure 3(3)), but the resulting oxides did not sufficiently sinter with the matrix. In the second stage, subsequent reduction at 700 °C yields αFe (Figure 3(4), see also Figure 1(a3,b4)). To regenerate γFe, the obtained αFe nanoparticles were first completely oxidized at 500 °C (Figure 3(5)), which provides sufficient sintering of the formed αFe₂O₃ particles and the matrix. Subsequent reduction, at the same temperature of 700 °C, regenerated the γFe nanoparticles (Figure 3(6)). According to the Mössbauer spectra, the Fe⁰ (metal)/Fe³⁺ (in matrix) ratio remained approximately the same during redox cycling (Table 1).

The TEM images of the reduced sample with ~11% of γFe and ~6% of αFe, obtained by the ACAR-promoted exsolution reaction of Ca_0.7Y₀₃(Fe_0.9Zn_0.1)O_2.6 with h-BN, after an oxidation–reduction cycle demonstrate features similar to uncycled (only exsolved) nanocomposites: using EDX analysis, the Fe nanoparticles can be identified in some locations, embedded in the oxide matrix, along with some Y segregation (Figure S3).

2.5. ACAR-Promoted Exsolution of FeNi_x Alloys

It is well-established that the existence of a region of the γ form can be significantly extended to lower temperatures by creating an Fe alloy with certain elements, such as Ni [29]. The exsolution of NiFe alloy nanoparticles from Ni and Fe-containing perovskite-like oxides has also been reported [10,16,21], so the formation of FeNi_x alloys was expected during our process as well. Ni is less chemically active and more reducible compared to Fe, so Ni additions to La, and especially Ca, ferrites significantly decrease their stability in a reducing H₂-containing atmosphere [27]. For this reason, the starting ferrites with low Ni content were used herein, to avoid decomposition. According to the XRD patterns shown in Figure 11a, the ACAR-promoted exsolution with h-BN from LaFe_0.8Ni_0.2O_2.9, La_0.8Ca_0.2Fe_0.8Ni_0.2O_2.8, and La_0.5Ca_0.5Fe_0.9Ni_0.1O_2.7 can be written as follows:

LaFe_0.8Ni_0.2O_2.9 + BN → FeNi_x + LaBO₃ + LaFe_0.8+βNi_0.2−βO_2.9+δ + N₂

(14)

La_0.8Ca_0.2Fe_0.8Ni_0.2O_2.8 + BN + H₂ → FeNi_x + LaBO₃ + Ca₃La₃(BO₃)₅ + La_0.8+zCa_0.2−zFe_0.8+βNi_0.2−βO_2.8+δ + N₂ + H₂O

(15)

La_0.5Ca_0.5Fe_0.9Ni_0.1O_2.7+ BN + H₂ → FeNi_x + Ca₃La₃(BO₃)₅ + Ca₃B₂O₆ + La_0.5+zCa_0.5−zFe_0.9+βNi_0.9−βO_2.7+δ + N₂ + H₂O

(16)

In these reactions, the resulting borates are the same as for the Ni-free ferrites described before (Figure 11(a1,a2,a4)). Note that while it is difficult to achieve exsolution levels of more than ~50% using metal Fe for the Ni-free ferrites due to cation mobility limitations, the reaction extent can be significantly higher for Ni-containing ferrites. As follows from the Mössbauer spectra shown in Figure 11(b1,b2,b4) and Table S7, Ni additions effectively stabilize the fcc structure of the exsolved nanoparticles. Their spectra are mainly comprised of two subspectra. The first, magnetically split with broad lines, correspond to Fe³⁺ cations in ferrites. It was fitted as a set of sextets, with δ~0.37 mm s⁻¹ and H_hf~49–53 T. The second subspectrum, which was fitted as a paramagnetic singlet with δ~−0.07 mm s⁻¹, corresponds to Fe in γFeNi_x alloys. According to the Fe–Ni phase diagram, at temperatures close to ambient, the αFe-based bcc phase αFeNi_x coexists with intermetallic compounds of Fe₃Ni and FeNi fcc types in the Fe-rich region; although, the phase boundaries at these temperatures are difficult to determine [29]. The thermodynamically stable phases of FeNi_x alloys, viz. the Fe-rich bcc and the Ni-rich fcc alloys, are magnetically ordered at room temperature [45]. At elevated temperatures, there is a continuous solid solution of γFeNi_x with the eutectoid temperature of ~345–400 °C at ~50 at% of Ni. At ~10–20 at% of Ni, which matches the initial ferrite stoichiometry, the transition temperature is about ~650–700 °C. Since Ni is more reducible than Fe, the alloys at low degrees of exsolution will be enriched in Ni and their transition temperatures will be even lower. These temperatures are below the temperatures of 700–750 °C at which reactions (14, 15, and 16) were carried out. When cooled to room temperature, the XRD patterns show that all the exsolved FeNi_x nanoparticles retained their γ structure at room temperature (Figure 11(a1,a2,a4)). The Mössbauer subspectra corresponding to the FeNi_x nanoparticles consist of singlets with δ~−0.1 mm s⁻¹ (Figure 11(b1,b2,b4)), evidencing that they are paramagnetic. It is consistent with their γ form, since paramagnetic behavior at room temperature has been reported for metastable γFeNix alloys with high Fe content [46,47]. The stabilization of the most exsolved γFeNi_x nanoparticles in γ form upon cooling, which can be explained by the strain developed due to their localization in voids, suggests that this is the main type of localization. Similar to γFe exsolved nanoparticles, a large part of γFeNi_x nanoparticles can be transformed into ferromagnetic αFeNi_x (Figure 11(b3)) by gently grinding in a mortar and pestle for 1–2 min. This treatment significantly reduced the FeNi_x reflections visually, in the XRD pattern of those samples (Figure 11(a3)). The mechanical stress-induced martensite γ to α transformation in Fe-rich Fe-Ni bulk alloys at room temperature has been previously reported in the literature [48,49].

The exsolved γFeNi_x nanoparticles can be completely oxidized at 500 °C and above, to the spinel solid solution Ni_1±xFe_2±xO₄ (Figure 11(a5,b5)). Subsequent reduction in 10% H₂/Ar atmosphere at 650–700 °C will restore the γFeNi_x metal nanoparticles (Figure 11(a6,b6)). It is noteworthy that the exsolution level for this sample was >80% (Table S7).

The TEM images of Ni-containing samples after the ACAR exsolution from La_0.5Ca_0.5Fe_0.9Ni_0.1O_2.7 ferrite, with ~41% of Fe exsolved (Figure 11(a1,b1)), are shown in Figure 12a–d. Similar to other samples, the images show agglomerates of different contrast (Figure 12a–c). The embedded FeNi_x nanoparticles <30 nm in size, however, can be identified by EDX analysis, e.g., like in locations 2, 3, and 4 in Figure 12d. The alloy nanoparticles contained approximately about ~30 at% of Ni at this exsolution degree. Figure 12e shows the TEM images of this sample after gently grinding in a mortar with a pestle. There is not much difference compared to the unground ones, except for some rounded agglomerates of <25–30 nm located on the grain surfaces enriched in Fe and Ni, which can be identified as alloy nanoparticles (Figure 12e, loc. 2). Figure 12f shows the TEM images of the same unground sample after the oxidation (at 500 °C, air)/reduction (700 °C, 10% H₂/Ar) cycle. Here again, it looks similar to the original sample with embedded FeNi_x nanoparticles, which were identified by EDX analysis (Figure 12f, loc. 2).

2.6. Separation of the Individual Nanoparticles

The reversible oxidation/reduction of metallic nanoparticles, while maintaining their localization within the in situ-created individual voids, is a remarkable feature of the nanocomposites produced by ACAR-promoted exsolution. It makes it possible to transform the initial metallic nanoparticles into various oxides, viz. Fe_1−xO, Fe₃O₄, α/γFe₂O₃ etc., while preventing their agglomeration. The components of matrices, viz. the starting and resulting ferrites, resulting borates, etc., in turn, have very diverse chemical properties depending on their composition. In particular, unsubstituted, and certain substituted, Ca ferrites are susceptible to hydrolysis and dissolve in dilute mineral acids like HCl. Ca₃B₂O₆ is also soluble in dilute acids. Since Fe₃O₄ (and γFe₂O₃) particles can dissolve reasonably slowly in dilute acidic aqueous solutions at RT, the crystallized and sintered Fe oxide nanoparticles, produced by the oxidation of the exsolved metal nanoparticles, can be separated from such oxide matrices by acid hydrolysis.

The TEM and SEM images of the nanoparticles separated from the Ca₂Fe₂O₅/h-BN and Ca₂FeAlO₅/h-BN-derived nanocomposites (~30–35% of Fe exsolved) are shown in Figure 13 and Figure S5. Ca₂Fe_2−xAl_xO₅ ferrites of the brownmillerite type, including Ca₂FeAlO₅, obtained by the isovalent substitution of Fe³⁺ cations with Al³⁺, are also prone to hydrolysis as unsubstituted Ca₂Fe₂O₅. They react with h-BN, similar to Ca₂Fe₂O₅, according to the simplified reaction (17), where Al is mainly retained in the ferrite phase (Figure S4):

Ca₂Fe_2−xAl_xO₅ + BN + H₂ → yFe + Ca₃B₂O₆ + Ca₂Fe_2−x−yAl_x+yO₅ + N₂ + H₂O

(17)

The RT Mössbauer spectrum of the Ca₂FeAlO₅/h-BN-derived nanocomposite is shown in Figure S4(1). The subspectrum corresponding to Ca₂Fe_2−xAl_xO₅ is significantly broadened and unresolved due to the magnetic dilution by diamagnetic Al³⁺ cations. Therefore, the spectral contribution of exsolved Fe, which totaled ~35% (~5% γFe), was evaluated from the 78K Mössbauer spectra (Figure S4(2)).

Figure 13(a1,b1) shows the powder XRD patterns and Mössbauer spectra (Table S8) of the Fe₃O₄ nanoparticles separated from the Ca₂FeAlO₅/h-BN-derived nanocomposite. The transformation of metal Fe nanoparticles into Fe₃O₄ nanoparticles was carried out by their oxidation to αFe₂O₃ at 500 °C, followed by their reduction in 10% H₂/Ar flow at 460–480 °C. The matrices were dissolved in 0.5–1.5% HCl aqueous solutions at RT. The Fe₃O₄ nanoparticles were magnetically separated and washed thoroughly with distilled water. It is worth noting that the leaching process of Fe₃O₄ can be carried out with high yields. This is facilitated by the fact that the particles were produced at relatively high temperatures at both stages of synthesis, viz. at oxidation of the metallic particles to Fe₂O₃ and at their subsequent reduction to Fe₃O₄ and were, therefore, crystallized and sintered well, which should reduce their solubility in diluted acids. Similar acid leaching in diluted aqueous HCl solutions has previously been successfully utilized to extract SrFe₁₂O₁₉-based nanoparticles, during high-temperature glass–ceramic synthesis [50].

The HAADF-STEM and SEM images of the separated nanoparticles obtained from the Ca₂Fe₂O₅/h-BN and Ca₂FeAlO₅/h-BN-derived nanocomposites (~30–35% of Fe exsolved) reveal that most of them are cap-shaped (hemispherical) hollow nanostructures with sizes below ~200 nm and ~100 nm, respectively (Figure 13c,d and Figure S5). According to Figure 13 and Figure S5, the wall thickness of the oxide nanoparticles is of the order of a few nm, thus the wall thickness of the parent metallic nanoparticles should be of the same scale.

The separated nanoparticles provide insight into the shape and size of the parent metal nanoparticles within the matrices. These hollow cap-shaped oxide nanoparticles indicate that the exsolved metal nanoparticles are formed as an inner layer in individual voids generated in situ within the matrices. This particle shape can potentially provide free space sufficient to accommodate the volume increase during oxidation, if the hollow particles are not completely filled with other reaction products. At the same time, these shapes and locations appear to play an important role in creating the conditions necessary for the development of compressive strain in metal nanoparticles sufficient to stabilize γFe, which is somewhat analogous to the γFe formation in cylindrical pores [34].

The extracted nanoparticles were different when La-substituted Ca_1−xLa_xFeO_2.5+δ ferrite was used as the starting material in the ACAR-promoted exsolution process. Figure 13e,f and Figure S6 show the TEM images of the oxide nanoparticles extracted from the nanocomposite obtained by the reaction of Ca_0.8La_0.2FeO_2.6 with h-BN (~35% of Fe exsolved). Metallic Fe nanoparticles were converted into Fe₃O₄ nanoparticles in a similar manner as described above. According to reaction (14), the La content in the remaining perovskite oxide Ca_0.8−zLa_0.2+zFeO_2.6+δ increases and, thus, its solubility in the dilute HCl aqueous solution decreases.

Using Mössbauer spectroscopy and XRD analyses, it was found that the extracted nanoparticles consisted mainly of spinel Fe₃O₄ and perovskite La_0.2+zCa_0.8−zFeO_3−γ phases (Figure 13(a2,b2), Table S8). The TEM with EDX analysis revealed that the La_0.2+zCa_0.8−zFeO_3-γ crystallites were segregated around Fe₃O₄ nanoparticles (Figure 13e,f and Figure S6). Most of them were almost completely covered by the resulting La–Ca ferrite, forming open shells. The size of the Fe₃O₄ nanoparticles was estimated to be smaller (<40–50 nm) than in the previous case. The composition and microstructure of the extracted nanoparticles reflect the relative arrangement of the metal/resulting ferrite phases in the parent nanocomposites. They show that the growth of exsolved metal nanoparticles in matrices is accompanied by the formation of shells, consisting of the resulting La-rich ferrites.

2.7. Variety of ACAR-Promoted Metal Exsolution Reactions

The new approach we have developed is not limited to the aforementioned ferrites and h-BN as starting compounds. The range of suitable compounds is wider and includes, but is not limited to, for example, Sr-containing ferrites and some other nitrides like Si₃N₄, etc. In particular, Figure S7 (Table S9) illustrates the ACAR-promoted exsolution of metal Fe nanoparticles from Sr₂Fe₂O₅, with Si₃N₄ as a reducing agent, according to the simplified reaction (18):

Sr₂Fe₂O₅ + 1/3Si₃N₄ + H₂ → 2Fe + Sr₂SiO₄ + H₂O + 2/3N₂

(18)

Even CO gas can be utilized in reactions of this type with Ca and Sr ferrites, when it is added to 10%H₂/Ar flow at ~650 °C, according to reaction (19):

(Sr,Ca)₂Fe₂O₅ + 2CO + H₂ → 2Fe + 2(Sr,Ca)CO₃ + H₂O

(19)

This reaction also leads to embedded metal α and γ Fe nanoparticles, as shown in Figure S8 [51].

The chemical activity of reducing reagents in ACAR processes is different, so that Si₃N₄ and CO are active with Ca and Sr-based ferrites, forming the corresponding silicates or carbonates, but are inactive towards La-based ferrites.

2.8. In Situ Reactions of Exsolved Nanoparticles

The exsolved metallic nanoparticles and their oxides can also be converted into other functional compounds, directly inside the matrices. For instance, embedded αFe₂O₃ nanoparticles obtained by the oxidation of exsolved metallic nanoparticles inside (La-based ferrite La_0.8Ca_0.2Fe_0.8Zn_0.2O_2.8)/h-BN-derived nanocomposites (Figure 14(1)) can be converted into embedded χ-Fe₅C₂ (Hagg carbide) nanoparticles by the reaction with gaseous CO at 300–350 °C (Figure 14(2), Table S10), according to the scheme in Figure S9. This carbide is considered to be an active phase in Fischer–Tropsch synthesis [52,53,54]. Its Mössbauer subspectrum at RT consists of three broad Zeeman sextets, with δ of 0.18–0.25 mm s⁻¹ and H_hf of 10–22 T, providing reliable identification [52]. The matrix, consisting of La ferrites and La borates, does not react with CO under the above conditions and remains unchanged. The α and, more significantly, γ Fe metallic nanoparticles can then be regenerated by the oxidation of carbide to αFe₂O₃ at 500 °C, followed by reduction to metal (Figure 14(3)), indicating that the metal/oxide/carbide nanoparticles retain their localization in the voids.

3. Materials and Methods

3.1. Materials Preparation

The initial perovskite-like compounds used for the nanocomposite syntheses were prepared using a sol–gel routine from the respective nitrate precursors and citric acid as a complexing agent. All the chemicals were of reagent grade, from commercial sources (AO Reachem, Moscow, Russia). The aqueous solution of metal nitrates in the calculated amounts were mixed with concentrated citric acid solutions in the molar ratio of citric acid to metals at ~1.1. The pH of the solution was adjusted by a dilute ammonia solution to values of ~5–7. Several samples, mostly Sr-containing ones, were synthesized using ethylenediaminetetraacetic acid (EDTA) as a complexing agent. In this case, the respective nitrate solution was dropwise added to an EDTA solution in diluted aqueous ammonia and then the pH was adjusted to ~8. The water was evaporated from the mixed solution and the resulting viscous gels were heated to ~200 °C to yield porous brownish materials. The materials were ground and then calcined in air at 500 °C (heating rate was 5 °C/min) for 4 h. The samples were then calcined in air for 8–10 h at 700–900 °C. The substituted perovskites, usually containing Fe⁴⁺ and/or Fe^(3+δ)+ cations in mixed oxidation states, were additionally treated in 10% H₂/Ar gas flow at 600 °C for 1 h to reduce them to Fe³⁺.

The nanocomposites with exsolved particles were synthesized through the solid-state method. The required amounts of perovskite powders and h-BN (Plasmotherm, Moscow, Russia) or other reducing reagents were dry grounded in an agate mortar until homogeneous for 20–30 min. The mixture powders were placed into a silica reactor in alumina crucibles and heated in 10% H₂/Ar gas flow of ~1–4 mL/min. The synthesis temperatures were ~650–700 °C and ~700–780 °C for the Ca- and La-based ferrites, respectively. For most of the samples used in this work, the exsolution reactions were carried out in a single step over ~10–16 h to reach completeness or a steady state. The experimental conditions for the synthesis of the nanocomposites are summarized in Table S1. The different extents of the exsolution reactions in this work were mainly predetermined by the ferrite/(h-BN, etc.) ratio, and the final amounts of exsolved Fe were quantified experimentally using Mössbauer spectroscopy.

3.2. Materials Characterization

The phase composition of the nanocomposites was characterized by X-ray powder diffraction (XRD), with a Huber G670 Image Plate Guinier diffractometer (CuK_α1 radiation, curved Ge monochromator, image plate detector) (Huber Diffraktionstechnik GmbH & Co., Ltd., Rimsting, Germany).

The ⁵⁷Fe Mössbauer spectra were recorded in terms of the transmission geometry on a commercial MS-1104EM spectrometer, using a ⁵⁷Co(Rh) source (13.5 mCi), equipped with a custom liquid nitrogen bath cryostat (77–320 K) (ZAO, Kordon, Russia). The fitting procedure was performed with a special original least-squares fitting software developed at Lomonosov MSU. All the isomer shift values refer to αFe at room temperature (RT).

Scanning electron microscopy (SEM) analysis was performed using a Quattro S scanning electron microscope (LaB6 field emission cathode, Thermo Fisher Scientific, Bleiswijk, The Netherlands). Transmission electron microscopy (TEM) images and energy-dispersive X-ray (EDX) spectra were obtained with a JEOL 2100 F/Cs, operated at 200 kV and equipped with a JEOL system for EDX analysis (JEOL Ltd., Tokyo, Japan), and using an aberration-corrected Titan Themis Z transmission electron microscope at 200 kV (Thermo Fisher Scientific, The Netherlands), equipped with a Super-X system for EDX analysis. For the TEM studies, the powder samples of the nanocomposites were prepared by dispersing in ethanol or heptane without additional grounding to avoid sample destruction, and then depositing a few drops of the suspension on special grids.

4. Conclusions

The novel approach to metal exsolution using ABO_3−γ (B = Fe, Ni, etc.) complex perovskite oxides of reducible transition metals has been developed to produce Fe-containing metal and metal oxide nanoparticles and nanocomposites. Our strategy is based on the additional auxiliary reactions of irreducible alkaline-earth and/or rare-earth A-site cations of the parent perovskite oxides with the constituents of reducing compounds resulting in the formation of complex oxides, which facilitate the exsolution of reducible transition metals from the B-sites of perovskites. The approach was applied mainly to Ca- and La-based unsubstituted and substituted ferrites as starting compounds and h-BN as a reducing compound. The metal nanoparticles exsolved by the described approach are largely localized within oxide matrices, in individual voids generated in situ. Moreover, the metal nanoparticles of Fe and FeNi_x alloys can be formed and stabilized in γ form, presumably due to induced compressive strain that develops in such locations.

The redox behavior of the exsolved nanoparticles was investigated. It was shown that they readily undergo redox cycling at moderate temperatures, where they maintain localization in the voids. The latter allows, under certain conditions, the regeneration of γFe and γFeNi_x from oxides during redox cycling. The nanoparticles can be chemically modified by reagents inside the matrices, while maintaining their localization.

The oxide matrices consist of unreacted initial perovskites and reaction products (borates, silicates, resulting perovskites, etc.) in various combinations and exhibit a wide variety of chemical and physical properties. They undergo significant transformation during the exsolution process, and their phase and chemical compositions depend on various factors, such as the composition of the initial ferrites and reducing reagents, the extent of reactions, and the reaction conditions, etc.

For certain Ca and Sr ferrite-based nanocomposites, a method has been developed for the separation oxide nanoparticles produced by the oxidation of the initially exsolved metal nanoparticles. The separated Fe₃O₄ nanoparticles (which can also be γFe₂O₃, etc.), obtained using Ca ferrites as starting materials, had a cup-shaped hollow shape. The separated nanoparticles obtained using La-substituted Ca ferrites were composite Fe₃O₄/(La-based resulting ferrite) layered particles. The separated nanoparticles provide insight into the shape, size, and relative arrangement of the phases of the parent metal nanoparticles within the matrices. This approach provides a new high-temperature templateless route for the synthesis of hollow and layered nanoparticles of Fe oxides and related compounds.

Our proposed strategy is applicable to a wide range of both perovskite-like oxides and reducing compounds. The developed approach provides greater flexibility in controlling the reaction extent, as well as the composition and properties of the resulting nanocomposites and nanoparticles. The approach can also be applied to the surface modification of perovskite oxide ceramics, since the starting perovskite compounds are considered stable under the reaction conditions. Our novel strategy is expected to be viable for the development of various new functional materials.