Size-dependent magnetic properties of NiO nanoparticles synthesized via Ni–hydroxyacetate decomposition

Abstract

This study focuses on a systematic investigation of the structural and magnetic properties of NiO nanoparticles (NPs) synthesized via the thermal dehydroxylation of a polyol-derived nickel layered hydroxyacetate salt (Ni–LHS). The turbostratic brucite-like Ni–LHS precursor undergoes a phase transition into pure NiO at 260–600 °C, yielding single-phase NPs with crystallite sizes ranging from ∼4.5 to 18 nm. Structural analyses reveal a transition from a sheet-like to a rougly spherical morphology and a progressive lattice contraction with increasing annealing temperature. Magnetic measurements at 5 K reveal characteristic signatures of nanoscale antiferromagnetism, including unsaturated hysteresis loops, high coercivity values, and pronounced exchange bias (EB). Both coercivity and EB exhibit a non-monotonic size dependence, reflecting a crossover from surface-disordered to core-ordered magnetic behaviour. These findings provide new insights into the interplay between finite-size effects and interfacial anisotropy in antiferromagnetic NiO NPs.

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Introduction

The magnetic behaviour of antiferromagnetic (AFM) nanoparticles (NPs) can markedly differ from that of their bulk counterparts. Typical size-dependent phenomena include the reduction of the Néel temperature,^1–4 the emergence of ferromagnetic (FM) like behaviours, superparamagnetic (SPM) blocking,^5,6 exchange bias,^7–9 or even spin-glass like phenomena.^5,10,11 These effects are largely due to uncompensated surface moments arising from the high surface-to-volume ratio.^12,13

NiO nanoparticles are a well-suited system for studying size-dependent magnetic properties. Although their magnetic structure was initially described by a simple core–shell-like model (in which an AFM core is exchange-coupled to a shell of canted, uncompensated spins^13–15), additional complexities have been reported in the literature, pointing to richer scenarios. For instance, Kodama et al.¹⁶ proposed a size-dependent evolution of the magnetic structure from the two-sublattice type-II AFM (observed in bulk NiO^17,18) to a 4-, 6- or 8-sublattice organization, to explain the large particle moments and exchange bias of NiO NPs. Rinaldi-Montes et al.^5,6 have also reported the loss of long-range AFM order at sizes around 2 nm.

Here, we investigate the size-dependent magnetic properties of NiO NPs in the ∼4.5–18 nm range, which span the crossover from strongly finite-size-dominated to near-bulk behaviour. Samples of different sizes were obtained by annealing a nickel hydroxyacetate salt (Ni-LHS) Ni(OH)_1−x(CH₃CO₂)_x·nH₂O synthesized via the polyol method.^19–21 Zero-field-cooled (ZFC) isothermal magnetization curves reveal large coercivities and unsaturated, ferromagnetic-like loops, while field-cooled (FC) magnetization curves exhibit the characteristic signatures of exchange bias-enhanced coercivity, a horizontal loop shift, and increased remanence.²² In what follows, we show how annealing-driven size control of NiO NPs allows us to study the crossover from finite-size-dominated to near-bulk magnetic behaviour, clarifying the mechanisms underpinning exchange bias in antiferromagnetic nanomaterials.

Experimental methods

Sample preparation

The starting reagents were all purchased from Sigma-Aldrich. NiO NPs of different sizes were prepared by annealing a layered nickel hydroxyacetate salt (Ni–LHS) powder, prepared by means of the polyol method.^19–21 Firstly, Ni–LHS was synthesized by dissolving 19.9 g of nickel acetate tetrahydrate (Ni(O₂CCH₃)₂·4H₂O, 98%) in 1 L of diethylene glycol (DEG, 99%) and 14.4 mL of distilled water. The mixture was then stirred and heated at 180 °C for 3 h (heating rate of 6 °C min⁻¹) and then cooled with a water/ice bath to room temperature. The green powder produced in this process was then recovered and washed by centrifugation with abundant ethanol and dried overnight at 60 °C.

Subsequently, the powder was annealed in air for 2 h at temperatures ranging from 200 to 600 °C, using a heating rate of 20 °C min⁻¹. This post-treatment step in air was performed to promote phase evolution. The complete dehydroxylation and transformation into NiO occur near 300 °C, as shown by temperature-resolved X-ray diffraction studies.^23,24 The different annealing temperatures result in crystal sizes varying from ∼4.5 to 18 nm. Hereafter, NiO NP samples will be named after their size, e.g., NiO_X, where X stands for the crystal size (nm).

Instrumental techniques

X-ray diffraction (XRD) patterns were recorded in the 5–110° 2θ range on a Panalytical diffractometer (multichannel X’celerator detector; Co Kα). The NP crystal sizes were assessed by Rietveld refinement performed with the MAUD software.²⁵ The data of NiO samples were modelled according to the face-centered cubic (fcc) rock-salt crystal structure (Fm3̄m space group²⁶). The effect of the rhombohedral distortions caused by magnetostriction at T < T_N is in the order of 0.1%, and was therefore neglected.^6,14,27 Transmission electron microscopy (TEM) was performed on a JEOL JEM-100CX-II. Thermogravimetric analysis (TGA) was performed on the Ni–LHS samples with a Setaram TG92-12. The analysis was performed in air at a 5 °C min⁻¹ heating rate up to 800 °C.

The magnetic investigation was performed with a Quantum Design MPMS-5S superconductive quantum interference device (SQUID) magnetometer. Field-dependent isothermal magnetization curves M(H) were measured at 5 K in a ±7 T range in zero-field cooled (ZFC) mode. M(H) curves were also measured in field cooled (FC) mode by cooling the samples from 300 to 5 K under a 7 T field to study exchange bias effects. Horizontal shifts of the M(H) curves were expressed in terms of exchange bias field H_EB = (H⁺_c − H⁻_c)/2 (where H^±_c are the positive and negative coercivity values estimated from the FC loops). M(H) curves were also measured on the pristine Ni–LHS sample.

Results and discussion

Layered hydroxyacetate: structure, morphology and magnetic properties

The XRD pattern of the as-prepared Ni–LHS (Fig. 1a and b) shows the layered hydroxyacetate structure with turbostratic disorder,^21,28 and a Ni(OH)_2−x(CH₃CO₂)_x·nH₂O stoichiometry. Turbostratic disorder is a particular type of stacking fault typically observed in layered structures, in which the periodic stacking along the crystallographic c axis is preserved, while the single layers are shifted and/or twisted with respect to one another. This structural feature results in low angle repetitive symmetric (00l) reflections with decreasing intensity, due to the stacking of single brucite-like layers, and sawtooth-shaped peaks at ∼40° and ∼70°, which can be attributed to the intralayer (100) and (110) planes, respectively^29–31 (Fig. 1a and b). This is in agreement with literature values for polyol-made Ni–LHS²¹ and β-Ni(OH)₂ (hk0) reflections.³² The Ni–LHS structure was observed until 240 °C, above which the material is transformed into the NiO phase. Before the transformation occurs, a shift at higher angles of the (00l) peak was observed, corresponding to a shrinkage of the interlayer basal spacing from ∼10.8 to 8.3 Å (Fig. 1b and Table S1). This shrinkage is presumably ascribed to the loss of the interlayer water and acetate anions, typically present in metal–hydroxyacetates produced by the polyol method.²¹

Fig. 1 (a) XRD patterns of the thermal evolution of Ni–LHS during annealing; the patterns are measured every 20 °C, with the transformation of Ni–LHS into NiO happening between 240 and 260 °C. (b) XRD patterns of the pristine, 220 and 240 °C treated Ni–LHS samples. (c) TGA (left axis, black curve) and DTA (right axis, red curve) of the pristine Ni–LHS powder. (d) Evolution of the NiO lattice parameter (full dots) and crystal size (empty dots) with respect to the annealing temperature. The dashed lines are a guide to the eye.

The TGA curve for pristine LHS (Fig. 1c) shows a first 10% loss completed after 100–150 °C, which is attributed to the loss of intercalation water, and another 25.2% loss at ∼310 °C, which is attributed to the departure/combustion of the remaining anions, leading to the formation of pure NiO. No significant losses are observed at higher temperatures, indicating that the transformation of the hydroxyacetate phase is completed in the vicinity of 300 °C. The mass loss accounts for ∼37%, which is compatible with the losses measured by Poul and co-workers.²¹ These results ensure that the LHS phase is totally decomposed, meaning that the magnetic data of NiO can be interpreted without considering the presence of other magnetic impurities. From the measured weight losses, the stoichiometry of the compound was determined to be Ni(CH₃CO₂)_0.68(OH)_1.32·0.64H₂O (M_w = 132.82 g mol⁻¹), in good agreement with previously reported values.^33,34 The layered structure is confirmed by the observed TEM morphology, which appears as aggregated, crumpled thin sheets (Fig. 2a), commonly observed in turbostratic Ni–LHS.^21,34

Fig. 2 TEM images of (a) Ni–LHS and of its products annealed for 2 hours at (b) 240 °C (still Ni–LHS), (c) 260 °C, (d) 300 °C, (e) 340 °C and (f) 360 °C, (g) 380 °C, (h) 400 °C, (i) 420 °C, (j) 440 °C, (k) 460 °C, (l) 480 °C, (m) 500 °C, (n) 520 °C, (o) 540 °C and (p) 600 °C (NiO).

M(H) curves of the Ni–LHS were measured at 5 K, observing a hysteretic loop with a saturation magnetization (M_S) of ∼87 Am² kg⁻¹ (Fig. 3) and a coercivity (µ₀H_C) of 77 mT, indicating a ferromagnetic response. Such intense low-temperature ferromagnetism (below ∼10–20 K) is a well-documented feature of layered double hydroxides, which are often characterized by saturation magnetizations as high as several tens of Am² kg⁻¹.^13,34–38 From the stoichiometry derived by TGA, we can calculate a 2.07 µ_B/Ni²⁺ moment, which is in agreement with the expected value for Ni²⁺ ions in a brucite-like structure.^21,34 Below the transformation temperature, while the brucite structure is still preserved and the interlayer distance is reduced, the low-temperature ferromagnetic behaviour is maintained, with almost unchanged saturation magnetization value (see also Table S2). This intense low-temperature ferromagnetism implies that residual LHS can affect the interpretation of the low-temperature magnetic properties of antiferromagnetic materials.³ However, all the samples were produced above the decomposition temperature of Ni–LHS observed in TGA measurements – or at least above the decomposition onset – and none of the NiO M(H) curves show features that can be associated with the presence of residual Ni–LHS (including those of Ni_4.5), meaning that NiO magnetic properties can be interpreted without considering the possible presence of impurities.

Fig. 3 M(H) curves (5 K) of the pristine (black) and the 240 °C-treated (red) LHS samples.

NiO NPs: structure, morphology and magnetic properties

The conversion of the Ni–LHS occurs between 240 and 260 °C, after which only the peaks of the fcc NaCl-like NiO phase (ICDD No. 00-047-1049) are visible, with no evidence of secondary phases. The phase transformation was also observed comparing TEM pictures of the samples annealed at different T, which show the growth of more or less aggregated roughly spherical NPs (Fig. 3).

The fast Fourier transform (FFT) patterns inferred from the high-resolution micrographs of the observed aggregates (Fig. S1) were unambiguously indexed as the fcc structure with a lattice parameter of about 4.2 Å (close to that of bulk NiO³⁹). At temperatures higher than 400 °C, the NiO nanocrystals start coalescing, leading to larger grains which start percolating at temperatures higher than 460 °C. From this temperature up to 600 °C, the TEM particle size becomes quite difficult to determine, thus the analysis of the diffraction line broadening of all the collected XRD patterns was preferably used to determine the NiO crystal size. The results are displayed in Fig. 1d.

Upon thermal annealing from 260 °C to 600 °C, the crystallite size of NiO exhibits two distinct linear growth regimes. Initially, up to approximately 450 °C, the growth is relatively slow; beyond this point, the growth rate increases significantly (approximately 5.5 times higher). This change suggests a transition in the dominant mechanism of crystallite evolution, likely related to enhanced atomic mobility and to the onset of coalescence or recrystallization processes once a critical thermal energy threshold is exceeded. Concurrently, the lattice parameter shows a progressive decrease with increasing annealing temperature, approaching 4.176 Å (Table S3). This contraction is likely due to the gradual elimination of structural defects such as oxygen vacancies or internal strain, which are commonly present in low-temperature-made or nanostructured NiO.⁴⁰

The ZFC M(H) curves of NiO samples at 5 K are shown in Fig. 4a, with an example of ZFC and FC loops for NiO_5.5 in Fig. 4b (complete datasets are available in Fig. S3 and S4). All the samples exhibit typical non-saturating hysteretic behaviour associated with AFM systems (key magnetic parameters reported in Table S4).

Fig. 4 (a) M(H) curves of NiO samples with different sizes (zoom in the inset), and (b) ZFC and FC loops for NiO_5.5, displayed as an example for better clarity.

As the NPs’ size increases, the ZFC coercivity increases to a maximum for the ∼6 nm sample (0.422 T), before decreasing for larger sizes. The FC coercivity displays a similar evolution, with enhanced values across the entire size range due to exchange coupling effects.²²

The coercivity increment ΔH_c = H^FC_c − H^ZFC_c also follows a maximum-like trend (Fig. 5a), consistent with previous findings.^7,16 This behaviour likely stems from a competition between size-dependent contributions: the growing AFM core anisotropy and the diminishing influence of disordered/uncompensated surface spins.⁵ A similar interplay is observed in the size dependence of the exchange bias field µ₀H_EB, which peaks at about 1.26 T for NiO_10.5 and then declines for larger particles. Such non-monotonic behaviour is characteristic of core@shell AFM systems (Fig. S5) where surface effects dominate at small sizes but become progressively less influential as the core volume increases.

Fig. 5 (a) ZFC (black squares) and FC (red circles) coercivities, plotted alongside the FC-ZFC difference (blue triangles), for NiO samples of different sizes; (b) size trend for the uncompensated moment (M_sFM, full squares), with the high field susceptibility χ_HF (open circles). The connecting lines are a guide to the eye.

Overall, the evolution of coercivity and exchange bias with particle size reflects the interplay between structural ordering in the core and magnetic disorder on the surface, which are distinct contributions modulated differently by the annealing temperature. In particular, while annealing improves crystallographic ordering within the core, surface spin disorder remains partly decoupled from the structural refinement and is largely governed by finite-size and surface-coordination effects. This distinction highlights the key role of nanoscale morphology in determining the balance between core and surface magnetic contributions. While AFM materials cannot saturate at experimentally achievable magnetic fields,¹⁴ it is common to express their magnetization as M(H) = M_FM(H) + χ_HFH, where χ_HF is the high-field susceptibility extracted by linear interpolation at fields higher than 4 T, and M_FM(H) is the field-dependent FM-like contribution,^8,14 for uncompensated surface spins. A graph with the loop and its contribution is shown in Fig. S2. From the resulting loop, it is possible to extract a ferromagnetic-like saturation magnetization M_sFM, which shows a steep decreasing trend upon size increase, as shown in Fig. 5b. This decrease is related to the decrease of the surface-to-volume ratio of the particles, meaning that bigger particles will have a smaller amount of uncompensated surface moments contributing to the total magnetization. Interestingly, χ_HF follows the same size-decay of M_sFM in the 4.5–10.5 nm range only to stabilize to a finite value of ∼0.12 Am² kg⁻¹ T⁻¹ for larger sizes. This strong correlation suggests that the two values share the same magnetic origin (i.e., the uncompensated shell moments), or at least that at sizes lower than 10 nm, the shell provides the biggest contribution to χ_HF. The remaining susceptibility contribution is likely caused by the compensated AFM structure and the core uncompensated moments.¹⁴

All the measured samples show exchange bias, which is here quantified with its bias field µ₀H_EB. The trend of µ₀H_EB shows a maximum at 10.5 nm, at a remarkable value of 1.26 T, which is in no qualitative contradiction with previously reported size-dependent values for NiO.^5,7,16 In the case of bilayer AFM/FM films, the interfacial unidirectional energy density Δσ is commonly defined by H_EB = Δσ/M_sFMt_FM, where M_sFM is the FM saturation magnetization and t_FM is the FM layer thickness of the FM film.⁴¹ While this model is formally derived for layered structures with well-defined FM and AFM interfaces, it is frequently adopted to study AFM nanoparticles, where the uncompensated surface acts as the FM component.^4,7,8 The model predicts an inverse relationship between H_EB and M_sFM. In our case, however, this inverse correlation is only observed within the 4.5–10.5 nm size range (Fig. 6). Beyond this threshold, H_EB decreases, breaking the expected trend. This deviation suggests that additional factors – such as partial compensation of shell magnetic moments, shell roughness, or changes in magnetic anisotropy – come into play at larger sizes.^22,41 In particular, the observed decay of the bias field at larger sizes might be influenced by the fact that the particles are polycrystalline (i.e., crystal diameter is smaller than the particle size observable by TEM, see Fig. 2), meaning that beyond a certain threshold, a considerable portion of the individual crystal surfaces will be in contact with other NiO particles by means of grain boundaries, which is rather different from the case of isolated single crystalline particles. This behaviour sets a crossover point between surface- and core-dominated regimes. The observed trends highlight the critical role of crystal size, morphology, and interface quality in governing exchange bias and related phenomena in antiferromagnetic NiO NPs (Fig. 7).

Fig. 6 Comparison of M_sFM and bias fields (μ₀H_EB, open circles) for different NiO sizes (d). The lines are a guide to the eye.

Fig. 7 An illustration of the investigation presented in this study is shown. NiO samples with different particle sizes were obtained by annealing the Ni–LHS precursor at increasing temperatures. The magnetic structure of each particle consists of an antiferromagnetically ordered core (orange) and a surface shell with uncompensated magnetic moments (blue). The interaction between these two regions – and the relative weight of each – varies with particle size, giving rise to the observed magnetic behaviour.

Conclusions

The controlled transformation of the turbostratic nickel hydroxyacetate precursor into NiO nanoparticles results in the transition from a layered to a spherical morphology, and the associated lattice contraction reflects the progressive ordering of the NiO crystal lattice with increasing annealing temperature. However, while annealing improves the core crystallographic order, the surface magnetic disorder remains only partially mitigated, as it is largely governed by finite-size and coordination effects rather than by the degree of structural refinement. The observed large coercivities and pronounced exchange bias fields, which reach their maximum for nanoparticles with ∼10 nm crystallites, therefore arise from a delicate balance between persistent surface spin disorder and the strengthening antiferromagnetic anisotropy of the ordered core. This crossover from surface- to core-dominated behaviour underscores the pivotal role of interfacial spin coupling in defining the macroscopic magnetic response of NiO nanoparticles. These findings advance the understanding of finite-size effects in antiferromagnetic systems and establish a framework for tailoring interfacial coupling independently of structural ordering, thereby enabling precise control of magnetic properties at the nanoscale.

Author contributions

All authors have contributed substantially to the experiments, writing and revisions of the manuscript in the current form.

Conflicts of interest

There are no conflicts to declare.

Data availability

The data is available from the corresponding author upon reasonable request.

The data supporting this article have been included as part of the supplementary information (SI). Supplementary information: additional data including TEM pictures and full magnetic hysteresis data for all samples. See DOI: https://doi.org/10.1039/d5cp03812j.

Acknowledgements

The authors acknowledge Labex SEAM (Science and Engineering for Advanced Materials and devices), ANR-10-LABX-096 and ANR-18-IDEX-0001 grants, for its financial support. Dr Patricia Beaunier (Sorbonne University) is acknowledged for her assistance in TEM measurements. P. M. and D. P. acknowledge the support of the “Network 4 Energy Sustainable Transition-NEST” project (code PE0000021), adopted by the “Ministero dell’Università e della Ricerca (MUR),” according to attachment E of Decree No. 1561/2022.

Notes and references

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