Coupled defect chemistry and redox dynamics in a bismuth gahnite system for self-adaptive tribological interfaces

Abstract

The development of self-adaptive, wear-resistant surfaces is critical for advancing next-generation tribological materials. In this study, a robust spinel zinc aluminate matrix is engineered with bismuth to incorporate lubricious phases that activate dynamically under sliding conditions. The strain and local structural distortion are brought about by introducing larger bismuth ions into the gahnite matrix. This doping also promotes oxygen vacancy formation, enabling a redox-active environment during wear. Under atmospheric oxygen and frictional heating, bismuth undergoes oxidative transformation at the wear track to form a lubricious bismuth oxide layer. At the same time, defect-rich, high-stress regions promote the reduction of bismuth through vacancy-mediated electron transfer, establishing a reversible redox loop. This repeated cycle enhances surface reactivity, enabling bismuth migration and the formation of surface oxides. Tribological testing using electroless NiP matrices on steel surfaces confirms bismuth redox cycling and tribofilm behavior, which critically influences the wear response. Raman spectroscopy and OSP show dynamic surface transformations and phase evolution. These findings demonstrate a synergistic interplay between redox-active dopants, structural modulation, and oxygen vacancy dynamics, offering a pathway to design chemically adaptive, high-performance tribological coatings.

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1 Introduction

Friction is a fundamental factor influencing the efficiency and performance of machinery in industry. The elimination of unnecessary friction is highly essential for maintaining fuel efficiency and reducing maintenance costs in the automotive, aerospace, electronics, power generation, and manufacturing industries.¹ Frequent exposure to mechanical and thermal driving forces alters the surface, particularly for metals or alloys. The mechanical and physicochemical characteristics of the contact components, as well as dynamic surface changes, including the tribo-induced oxidation and chemical transformations at the sliding interface, are intrinsically related to these tribological processes.² Lubricity and tribological performance are typically linked. The development of a lubricant layer on the contacting sliding surface is known as lubricity. This tribofilm should always have an ideal thickness. Because there is less direct contact between surface asperities, high lubricity reduces friction and energy loss.³

The structural origin of lubricating materials is associated with their Magnéli structure or two-dimensional layered structure. Due to the deformed metal oxide and octahedral geometry, the elastic shear constant for these phases is low, leading to considerable interlayer distances. The lubricating properties of materials are rooted in their bonding nature and capability to form low shear interfacial phases, emphasizing the importance of structural design in the development of advanced materials for tribology.⁴

Zinc aluminate, belonging to the spinel family, has gained attention due to its exceptional chemical and thermal stability and good mechanical strength. In contrast to other surface oxides, spinel oxide layers on alloys have lower shear strength and higher negative formation and cohesive energy, which encourages long self-lubrication.⁵ Several mechanisms have been proposed in the literature to explain the superior mechanical integrity and tribological performance of binary and ternary oxides. These include (i) facile shearing enabled by the screening of cations by neighboring anions, as described in the crystal-chemical model;⁶ (ii) thermal softening corresponding to the brittle-to-ductile transition of many oxides, analogous to the lubricious response of soft metals;⁷ (iii) lubrication by oxides with layered crystal structures, where weak interplanar bonding promotes easy shear, akin to graphite and MoS₂;⁸ and (iv) grain-boundary-mediated shear in nanocrystalline oxides, where dislocation glide induces intracrystalline slip.⁹

Bismuth has been used as an active component in self-lubricating alloys and is a greener alternative lubricant. It has been found that bismuth particles can form non-bonded films on steel that are easily removed from lubricating surfaces, and this is due to the poor solubility of bismuth. Bismuth metal has a close-packed quasi-metallic layered structure at room temperature.¹⁰ Because the bare 6s and 6p electrons are easily vulnerable to various crystal fields, bismuth has numerous oxidation valences (Bi⁰, Bi⁺, Bi²⁺, Bi³⁺, and Bi⁵⁺), making it a multivalence emitter ion.¹¹ Under shear stress conditions, bismuth oxide has a high diffusion rate, which is beneficial for the formation of a glaze layer due to friction sintering, resulting in reduced wear and friction.¹² Incorporating bismuth into the gahnite phase leads to enhanced tribological characteristics.

Among several methods used for developing protective coatings over steel, electroless nickel coatings are being used more frequently in a variety of industrial sectors, such as the automotive and petroleum industries, where strong corrosion and wear resistance are required for components like tools, equipment, and valves. In the coming years, it is anticipated that the use of electroless nickel-based coatings in various applications will significantly increase due to the expected rise in performance requirements and operational endurance.^13,14 A wide range of micro and nanoparticles has been employed as reinforcements in Ni–P coatings. Solid lubricants such as WS₂,¹⁵ MoS₂,¹⁶ hexagonal boron nitride (h-BN),¹⁷ Ni–P–TiC,¹⁸ and graphene oxide (GO),^19,20 owing to their distinctive structural features, are frequently incorporated to enhance the wear resistance of the coatings. It was found that nanolubricants improved lubrication by easy conversion of sliding friction into rolling friction, due to the spherical shape of the nanoadditives.²¹ Nanoparticles can accumulate within microdefects on friction surfaces and, to some extent, contribute to a self-repairing effect.²²

In this context, the present work has explored tuning the lattice via cationic substitution and the creation of oxygen vacancy defects through lattice distortion to achieve substantially enhanced tribological performance. The study primarily focuses on crystal lattice tuning of the gahnite structure via bismuth doping for creating oxygen vacancy defects that will serve as sites for a redox-driven lubrication mechanism. Tailored morphological control can enable a grain boundary pinning effect that enhances microstructural stability and resistance to stress relaxation. The majority of systems that have been reported rely on externally supplied lubricants or passive hard-phase strengthening, which limits their ability to adapt in real time to changing contact stresses. It is still mostly unknown whether lattice defect engineering and redox-active dopants can be combined to create a self-sustaining lubrication cycle. Specifically, structural distortion, microstructural evolution, and macroscale friction-wear performance have not been consistently associated with the tribo-responsive redox chemistry of bismuth in oxygen-vacancy-rich spinel matrices. By adding Bi³⁺ to spinel ZnAl₂O₄ (gahnite) lattices in NiP electroless composite coatings, we fill this gap. Strain fields and oxygen vacancies that function as redox-active sites are produced by the lattice distortion caused by the significant ionic radius mismatch. This enables in situ oxidation combined with vacancy-mediated Bi^m+ redox cycling under tribological loading. In order to develop next-generation triboceramics that can adaptively mitigate wear, we demonstrate how to correlate tribological performance, crystallinity control, and defect chemistry. This work represents a significant advancement in material design for harsh mechanical conditions by showing that triboceramics can be fitted with autonomous, self-healing lubrication through lattice-level chemical tailoring. This investigation, offering tailored defect chemistry, will promote surface passivation and dynamic redox transitions under mechanical contact, thereby reducing wear & preserving surface integrity.

2 Experimental section

2.1 Synthesis of bismuth-doped zinc aluminate

ZnAl₂O₄ nanoparticles were synthesized using a modified citrate-assisted sol–gel technique. Zinc nitrate hexahydrate [Zn(NO₃)₂·6H₂O] (99.0%, Merck) and aluminium nitrate nonahydrate [Al(NO₃)₃·9H₂O] (99.0%, Merck) were employed as metal precursors, while citric acid (99.5%, Merck) served a dual function: as a chelating agent to stabilize metal ions and as a combustion fuel to promote crystallization at reduced temperatures. The chelation effect of citric acid facilitates early-stage nucleation and suppresses particle growth by effectively capping the metal ions, especially zinc, thus minimizing agglomeration. An increased concentration of citric acid around the metal centres favors the formation of smaller nanoparticles, particularly when calcined at lower temperatures.

In a typical synthesis, stoichiometric amounts of zinc and aluminium nitrates were dissolved in deionized water under ultrasonic agitation to ensure complete dissolution. Separately, a citric acid solution was prepared in dilute nitric acid, maintaining a molar ratio of 1 : 2 (metal ions to citric acid). The two solutions were combined and heated in a water bath at 90 °C for 4 h, followed by magnetic stirring for an additional 3 h to promote homogeneity. The resulting sol was aged at room temperature until gelation occurred. The gel was then dried at 100 °C on a hot plate for 5 h, manually ground into a fine powder using an agate mortar and pestle, and finally calcined at 700 °C for 3 h in air to yield phase-pure ZnAl₂O₄ nanoparticles. To obtain the necessary compositions of ZnAl_2−xBi_xO₄ (x = 0.05, 0.10, 0.20, 0.50), a stochiometric amount of bismuth nitrate pentahydrate [Bi(NO₃)₃·5H₂O] (98%, Sigma-Aldrich) was added in addition to zinc and aluminium nitrates. The rest of the synthesis process remained unchanged.

2.2. Fabrication of ZA@NiP coatings over mild steel

Mild steel substrates (with dimensions of 3.5 × 2.5 × 0.1 cm³, provided by TATA Steel Limited, CR2 SPD, India) were utilized for electroless deposition following a multi-step surface preparation process to ensure optimal coating adhesion. The cleaning protocol involved the following stages: (a) mechanical preparation: the steel strips were polished to a mirror-like finish using a sequence of emery paper, progressing from coarse to fine grit, followed by rinsing with distilled water to remove debris. (b) Chemical cleaning: the mechanically polished substrates were immersed in a 5% NaOH solution for 5 min to eliminate organic contaminants, rinsed thoroughly with distilled water, and then treated with a 3% HCl solution for another 5 min to remove surface oxides, followed by additional rinsing with distilled water. (c) Sensitization: the chemically cleaned samples were immersed in a 10 g L⁻¹ SnCl₂ solution for 10 min to adsorb Sn²⁺ ions onto the surface, enabling nucleation sites for activation. (d) Activation: the sensitized strips were rinsed with distilled water and then immersed in a 1 g L⁻¹ PdCl₂ solution, where Sn²⁺ ions were replaced by catalytically active Pd²⁺ ions, preparing the surface for subsequent electroless deposition. The electroless bath was prepared by dissolving 3 g of nickel sulfate (NiSO₄, Molychem, 98.90%) and 2.5 g of succinic acid (HiMedia, 98.00%) as the nickel source and complexing agent, respectively, in 100 mL of deionized water. The pH of the solution was adjusted to 4.5 using aqueous ammonia. The solution was then stirred and gradually heated. Upon reaching 75 °C, sodium hypophosphite (NaH₂PO₂, Sigma-Aldrich, 99.99%) was added as both the phosphorus source and reducing agent. The MS coupon was then immersed in the bath solution, and then 1 g L⁻¹ synthesized ZA, ZAB_0.05, ZAB_0.10, ZAB_0.20 & ZAB_0.50 was added, respectively; a bare NiP coating was also synthesized without the presence of these compounds. The deposition process was carried out at a controlled temperature of 80 ± 2 °C for 2 h. After deposition, the coated samples were rinsed with deionized water and dried at 60 °C (Scheme 1).

Scheme 1 Schematic illustration of ZAB nanopowder and its incorporation into the NiP matrix.

2.3 Characterization techniques

The crystalline nature and phase purity of the synthesized samples were characterized by X-ray powder diffraction (XRD) using a powder X-ray diffractometer (Bruker D8 ADVANCE) with Ni-filtered Cu Kα (λ = 1.5405 Å) radiation. Data were collected by step scanning over a 2θ range from 10° to 90° with a step size of 0.02. A Carl Zeiss EVO 18 scanning electron microscope, which operates at 15 kV acceleration voltage, was used for recording the SEM images. A Carl Zeiss EVO 18 secondary electron microscope with an EDAX attachment of the AMETEK EDAX octane series connected to the SEM was utilized to do the energy dispersive X-ray analysis (EDX). The transmission electron microscopy (TEM) analysis was carried out using a Talos200 (Thermo Fisher Scientific) operated at 200 kV. Bright-field TEM was used to probe the sample morphology, while high-resolution TEM (HR-TEM) provided detailed structure visualization. Selected area electron diffraction (SAED) was used for crystal structure determination, and dark-field scanning transmission electron microscopy (DF-STEM) was utilized to image the morphology in dark-field mode. Energy-dispersive X-ray spectroscopy (EDS) enabled composition determination, and STEM-EDS was performed for elemental mapping to understand the spatial distribution of elements within the samples. The entire TEM data acquisition process was carried out using Velox V 3.15 software. The surface composition and elemental oxidation states of the materials are evaluated using a Thermoscientific X-ray photoelectron spectrometery (ESCALAB 250i) with Al Kα radiation as the excitation source. The surface roughness parameters of the developed coatings were evaluated using a non-contact optical surface profilometer integrated with an M370 scanning electrochemical workstation (Uniscan Instruments Ltd, UK). The EPR study was performed on a JEOL EPR spectrometer (Model JES-X320) at room temperature (RT), using an X-band frequency of 9.6 GHz. Raman spectroscopy analysis was done with a LabRam HR Evolution Raman spectrometer (Horiba Scientific), having a resolution of 1 cm⁻¹ and an excitation wavelength of 785 nm.

The microhardness of the fabricated coatings was assessed using a Vickers hardness tester (Model: SSS-VM-100 PC) under a load of 10 kgf and a dwell time of 10 seconds to monitor hardness variation. Tribological performance, including wear resistance and friction behavior, was investigated using a pin-on-disc tribometer (DUCOM TR-208-M1, WINDUCOM software). A standard cylindrical EN-8 grade mild steel pin (6 mm diameter) was employed as the counterbody, sliding against the coated surfaces under an applied load of 20 N. The tests were conducted at a rotational speed of 100 rpm for a duration of 3600 seconds along a wear track diameter of 10 mm. All measurements were performed under ambient conditions, maintaining a relative humidity of approximately 40%.

3 Results and discussion

3.1 Spinel metal oxides as nanolubricants

Oxide-based materials offer robust structural and chemical stability.²³ Because of their high surface energy and tendency to agglomerate, aggregate, and coagulate in lubricating media, metal oxides as lubricant additives are limited.²⁴ Thus, variable surface modification is the primary goal of this research. Compared to other surface oxides, spinel-based oxide coatings have lower shear strength and higher negative formation and cohesive energy, which encourages long-term self-lubrication.⁴ It entails creating hard, wear-resistant coatings that also serve as a source of elements for tribo-chemical oxidation reactions, producing desirable lamellar oxides and oxide-derivative phases at their surface that act as high-temperature lubricants. The softening or perhaps micro-melting of the proper phases produced inside tribo-contact should be the cause of the decrease in friction at higher temperatures. Naturally, we deal with chemical wear of basic, hard coatings in such a scenario due to variations in the current defect state density. Nanoparticles are prominent due to their strong van der Waals force and large surface area, thus posing a critical challenge to distinguish between aggregated and single particles. ZnAl₂O₄ nanoparticle synthesis has drawn increasing interest because of its excellent mechanical resistance, hydrophobicity, low surface acidity, and high thermal stability—all of which are critical for anti-friction applications. Nanoparticles are capable of self-healing, serving as a third body and depositing on the wear surface to diminish abrasion. To enhance the anti-wear capability and stop microcracks from spreading further, nanoparticles could be embedded into the frictional surface.²⁵

3.2 Strain modulation and defect engineering in gahnite via trivalent bismuth doping

3.2.1 Deciphering the nano zinc aluminate structure through Rietveld refinement and core level spectroscopy. Comprehensive structural and chemical investigation of nano zinc aluminate serves as the foundation matrix for the present system. The crystalline phase, structure, and purity of the synthesized ZnAl₂O₄ nanoparticles were investigated using X-ray diffraction (XRD). The XRD pattern of the sample is shown in Fig. 1(A). All the diffraction peaks were indexed and found to correspond to the characteristic reflections of the cubic spinel ZnAl₂O₄ phase, matching well with the standard diffraction data from the ICDD PDF card no. 00-005-0669. The observed peaks at 2θ values of approximately 31.07°, 36.6°, 44.6°, 48.9°, 55.6°, 59.3°, 65.05°, 74.08°, and 77.20° correspond to the (220), (311), (400), (331), (422), (511), (440), (620), and (533) crystallographic planes, respectively. The presence of sharp, intense peaks and the absence of any impurity reflections confirm the formation of a single-phase ZnAl₂O₄ spinel structure with high crystallinity.²⁶ The average crystallite size of the synthesized ZnAl₂O₄ nanoparticles was estimated using the Debye–Scherrer equation based on the full width at half maximum (FWHM) of the (311) reflection. The calculated crystallite size was found to be approximately 20 nm, indicating the nanocrystalline nature of the prepared spinel sample. Rietveld refinement of the XRD data was conducted using the TOPAS 4.2 software to determine the structural parameters. The simulation of the powder diffraction patterns was performed using the Fd3̄m space group, assigning Zn to the 8a, Al to the 16d, and O to the 32e Wyckoff positions, respectively. The obtained lattice constant was found to be approximately 8.1446 Å, which is consistent with literature values for stoichiometric ZnAl₂O₄.^27,28 The refined parameters are listed in Table S1. The crystal structure generated is shown in Fig. 1(B).

Fig. 1 (A) XRD pattern of nano ZA, (B) crystal structure of nano ZA, (C) graphical output of the Rietveld refinement of nano ZA, and (D) high resolution O 1s XPS spectra of nano ZA.

Fig. 2 (A) XRD pattern of ZAB_0.20, (B) crystal structure of ZAB_0.20, (C) graphical output of the Rietveld refinement of ZAB_0.20, (D) high resolution O 1s XPS spectra of ZAB_0.20, (E) W–H plot of ZA, (F) W–H plot of ZAB_0.20, (G) EPR spectra of ZA, and (H) EPR spectra of ZAB_0.20.

The graphical output of the Rietveld refinement of the ZnAl₂O₄ sample is shown in Fig. 1(C). The experimental data are shown as blue dots, and the calculated pattern is shown as a red line. The difference plot shows the residual between observed and calculated patterns, as a green line. It is seen that the best possible fit is obtained. ZnAl₂O₄ crystallizes in a normal spinel structure, where Zn²⁺ ions (ionic radius ∼0.60 Å) typically occupy the tetrahedral (A) sites, and Al³⁺ ions (ionic radius ∼0.535 Å) reside in the octahedral (B) sites. The cation distribution in the spinel lattice is primarily governed by factors such as ionic radii, charge balancing, and crystal field stabilization energy (CFSE). At the synthesis temperature of 700 °C, thermal energy facilitates the diffusion and redistribution of cations into their thermodynamically preferred sites, improving the long-range order and stabilizing the spinel structure. Because Al³⁺ ions have a lower crystal field stabilization energy in octahedral coordination compared to tetrahedral coordination, aluminium in spinel structures strongly favors octahedral sites, resulting in very limited cation site inversion in aluminate spinels. Therefore, the transition metal cations can reasonably be expected to be located in the tetrahedral A sites of the AB₂O_4, while Al³⁺ resides in the octahedral B sites. To further investigate the chemical environment and oxygen vacancies on the surface of gahnite, the XPS technique was employed. Fig. 1(D) shows the O 1s spectra which exhibit peaks at 529.9, 530.9, and 531.7 eV attributed to different oxygen species. The O 1s peak at 531.7 eV is usually attributed to the presence of loosely bound oxygen on the surface of the spinel framework; the low binding energy component located at 529.9 eV is attributed to the O²⁻ bound within the spinel framework. The medium binding energy component, centred at 530.9 eV, is associated with O²⁻ in the oxygen-deficient regions within the matrix of ZnAl₂O₄.²⁹ The survey spectrum of ZA and high-resolution spectra of Zn and Al are shown in Fig. S1. In the XRD analysis, it is generally believed that the peak intensity is related to the crystallinity of the sample. The spinel lattice thus remains chemically stable and well-ordered, making the material promising for applications in wear related properties, where structural robustness and phase purity are essential. Overall, the XRD results confirm that the optimized annealing process at 700 °C is sufficient to induce the formation of a pure, crystalline ZnAl₂O₄ phase, suitable for further electrochemical or surface characterization.

3.2.2 Size mismatch and lattice contraction induced by bismuth-structure and chemical environment modulation. The substitution of trivalent bismuth into spinel-type zinc aluminate introduces a unique interplay between ionic radii mismatch and lattice strain, which impacts crystal behavior. Bismuth has a significantly larger ionic radius (1.03 Å) compared to aluminium (0.535 Å), and replaces the octahedral sites of the spinel lattice. The XRD pattern of ZAB_0.20 is shown in Fig. 2(A). The Rietveld refinement indicated a slight decrease in the lattice parameter for ZAB_0.20, and the corresponding crystal structure is shown in Fig. 2(B). The graphical output is shown in Fig. 2(C). Rietveld refinement of the other samples is shown in Fig. S3. Rietveld refinement parameters obtained for all the samples are given in Table S1, and XRD spectra of all the other prepared compositions are depicted in Fig. S2. The intensity decrease with bismuth incorporation is manifested by a decrease in crystallinity. The slight shift towards higher angles up to ZAB_0.20 depicts the defects created in the lattice. At higher doping concentration of ZAB_0.50, the bismuth ions effectively start incorporating in the lattice sites of Al³⁺, indicated by the slight increase in interatomic distances. This is primarily due to the lattice attempting to accommodate the larger bismuth ions, resulting in tensile strain.³⁰ Although the crystal lattice is altered in both lattice distortion and strain, the spatial extent and type of structural perturbation are fundamentally different. The lattice strain estimated for ZA and ZAB_0.20 using the Williamson–Hall plot is shown in Fig. 2(E) and (F), and W–H plots of the rest of the compositions are shown in Fig. S4. Localized disruptions like point defects, doping, or atomic rearrangements are commonly the cause of lattice distortion, which causes sudden and incoherent changes in the lattice configuration. On the other hand, strain typically results from internal stress accumulation, epitaxial restrictions, or a coherent lattice mismatch. The scale and continuity of the ensuing lattice alterations are what set them apart. The symmetry-breaking effect and crystal defect theory can be used to understand the nature of lattice distortion.³¹ The crystal lattice's oxygen vacancy is the source of the lattice strain. One potential method for altering the electrical structure by changing the atoms' distances from one another is strain engineering.³²

The high-resolution O 1s XPS spectra of ZAB_0.20 are shown in Fig. 2(D). The oxygen-coordinated environment of an element may be slightly distorted by these vacancies, which also cause a local charge imbalance and are associated with the positive lattice strain found in W–H plot analysis. Lattice strain may be induced when an oxygen vacancy is introduced, as it can alter or upset the lattice d-spacing. The binding energy value for the O 1s peak shifted from 530.9 to 531.1 eV. More oxygen vacancies could push the O 1s peak to move in the direction of higher binding energy.³³Fig. 2(G) and (H) show the EPR spectra of ZA & ZAB_0.20. The EPR technique is a potential tool for the quantitative characterisation of oxygen vacancies.³⁴ X-band EPR at 9.6 GHz is employed to investigate the electronic environment and defect structures. The g factor of 2.003 (Fig. 2(G)), represents a strong paramagnetic signal that is typically attributed to an unpaired electron trapped on a surface oxygen vacancy site³⁵ (a single oxygen vacancy with one trapped electron F⁺ centre). The existence of a perturbing defect near a F⁺ centre, which is advantageous because of the excess positive charge (+1) of the typical F⁺ centre, can be used to explain the relatively low thermal stability of the F⁺-type centres in the examined spinels.³⁶ Because electrons move between nearby oxygen vacancies, the EPR spectrum becomes isotropic, indicating motional averaging of the anisotropy.³⁷

The (Bi³⁺–O_V) defect complex (O_V, oxygen vacancy) may be the cause of the g value being marginally greater than 2 (2.215).³⁸ In Fig. 2(H), the broad resonance observed arises not from Bi³⁺ itself (6s², closed shell, EPR silent) but from oxygen-vacancy-related centres (F⁺) strongly perturbed by the heavy Bi ion. The stereochemically active 6s² lone pair on Bi³⁺ leads to local off-centering and lattice distortions in Bi-based oxides.³⁹ The observed g-value exhibits a slight anisotropy and deviates from the free electron value (g = 2.0023). The admixture of excited quartet states with the ⁶S ground state via spin–orbit coupling typically causes a negative g-shift. In contrast, a positive g-shift is attributed to electron transfer from the ligand to the S-state ion mediated by spin–orbit interaction.⁴⁰ High-resolution XPS images of Zn, Al, and Bi, along with a survey spectrum of ZAB_0.20 are presented in Fig. S5. As seen in Fig. S1 and S5, the Zn 2p spectra of ZnAl₂O₄ and ZAB_0.20 do not appreciably vary in binding energy, suggesting that Bi inclusion does not affect the Zn chemical state.

3.2.3 Crystallinity driven morphological assessment. The enlarged (311) diffraction peak shown in Fig. 3(A) reveals the shift to a higher 2θvalue, where (a) ZA & (b) ZAB_0.20 indicate slight changes in the lattice caused by the incorporation of Bi³⁺ with differing ionic radii. Also, a lower crystallinity observed in ZAB_0.20 is evident in Fig. 3(A). Crystallinity and oxygen vacancies are closely connected. Higher Ov concentration originates from lower crystallinity. The Ovs can serve as nucleation centers.⁴¹ Bi³⁺ dopant induces lattice strain and segregates to grain boundaries, which drags boundary migration and suppresses grain coarsening, resulting in a non-faceted mulberry-like morphology as evident in the SEM micrograph of Fig. 3(C) compared to SEM micrograph of ZA shown in Fig. 3(B). The EDX spectrum shown in Fig. S6 confirms the presence of Zn, Al, O, and Bi elements, indicating their successful incorporation into the spinel matrix. Fig. 3(D–F) and (G–J) validate the homogeneous distribution of Zn (magenta), Al (cyan), O (green), and Bi (blue) elements across the matrix. The selected area for mapping and elemental overlay of ZA & ZAB_0.20 are shown in Fig. S7. SEM micrographs and EDX spectra, and mapping of other compositions are shown in Fig. S8. Metal oxide nanoparticles grow via Ostwald ripening. Tiny particles evaporate during this process, then settle or shift to become larger particles. This causes the granules to enlarge, which lowers the surface energy. Elastic strain caused by ionic size mismatch in the lattice is relieved by the presence of a free surface.⁴² Grain dislocation slip is the primary mechanism for stress reduction. Grain boundaries, however, may obstruct these slips and make mobility more difficult. Therefore, it is true that smaller grains can increase resistance to stress relaxation. The pinning effect of concentrated bismuth atoms at grain boundaries begins to influence the migration of dislocations. The addition of bismuth stops the dislocation density from decreasing.⁴³

Fig. 3 (A) Shift of the enlarged (311) peak in ZA and ZAB_0.20, (B) FESEM image of ZA, (C) FESEM image of ZAB_0.20, (D–F) X-ray dot mapping of Zn, Al, O in ZA, (G–J) X-ray dot mapping of Zn, Al, O and Bi in ZAB_0.20.

TEM micrographs of ZA and ZAB_0.20 shown in Fig. 4(A) and (E) revealed agglomerated nanoparticles with an average diameter ranging from 10 and 8 nm (Fig. S9). High-magnification views (Fig. 4(B) and (F)) highlight their clustered arrangement, with more pronounced agglomeration in ZAB_0.20, likely arising from enhanced van der Waals interactions at reduced particle sizes. The HRTEM images in Fig. 4(D) and (H) exhibit interference fringes with d-spacings of ∼0.4 nm and ∼0.29 nm, which correspond to the interplanar distances of the (111) and (311) planes in the cubic spinel structure of zinc aluminate. From XRD analysis, the d-spacing of the (311) plane was found to be 0.30 nm for ZA. Similarly, the HRTEM images depicted in Fig. 4(H) exhibit interference fringes with a d-spacing of ∼0.31 nm representing the interplanar spacing of the (311) plane of ZAB_0.20. This uniformity is ascribed to the volume-averaged nature of X-ray diffraction analysis. The localised variations in lattice parameters, which could stem from crystallite size effects or nanoscale surface strain, tend to occur in nanomaterials, as reflected in d-spacings obtained from HR-TEM. The STEM-EDS method was employed to further analyse each element's spatial distribution within the particles. Elemental mapping was carried out on the High-Angle Annular Dark Field (HAADF) images of each sample. Elemental mapping revealed a homogeneous spatial distribution of each constituent within the samples. HAADF images are displayed in grayscale, with elemental distributions of Al, O, Bi, and Zn overlaid in red, blue, yellow, and green, respectively. Thus, elemental analysis using STEM-EDS verifies the nanoscale purity of the synthesized materials. Elemental mapping further demonstrates a uniform distribution of Zn, Al, and O (Fig. 4(A2)–(A4)) and Zn, Al, Bi, and O (Fig. 4(E2)–(E5)) within the materials, where Fig. 4(A1) and (E1) correspond to HAADF-STEM images of ZA and ZAB_0.20, respectively. In addition, Fig. 4(C) demonstrates the SAED pattern of ZA, which confirms the polycrystalline characteristics associated with the (220), (311), (400), (422), (511), and (440) planes of ZA. SAED patterns (Fig. 4(G)) exhibit diffraction rings matching the (220), (311), (222), (400), (331), (511), (440), and (531) planes of ZAB_0.20, confirming the polycrystalline spinel structure. These findings corroborate the XRD results. Elemental analysis using STEM-EDS was conducted to verify the nanoscale purity of the samples presented in Fig. S10 for ZA and ZAB_0.20, confirming their high purity. Peaks of Zn, Al, Bi, and O were identified, with the sole additional feature being the Cu–K peak from the copper grid. Table S2 illustrates the changes in the elemental composition as the particle size decreases. Images of ZA and ZAB_0.20 reveal the assembling of small crystallites with very fine sizes. The overall results obtained from TEM-EDS reveal a combination of fine crystallite size, purity, and uniform elemental dispersion, implying structural changes brought about by lattice distortions at the nanoscale.

Fig. 4 (A) TEM micrographs of ZA, (B) magnified images of ZA, (C) SAED pattern of ZA, (D) HRTEM image of ZA, (E) TEM micrographs of ZAB_0.20, (F) magnified images of ZAB_0.20, (G) SAED pattern of ZAB_0.20, (H) HRTEM image of ZAB_0.20, (A1) HAADF-STEM image of ZA and STEM-EDS elemental mapping of (A2) Al, (A3) O, and (A4) Zn in ZA, (E1) HAADF-STEM image of ZAB_0.20 and STEM-EDS elemental mapping of (E2) Al, (E3) Bi, (E4) O, and (E5) Zn in ZAB_0.20.

3.3 Tribological performance modulation in the NiP matrix

3.3.1 Suitability of electroless NiP as a tribofunctional matrix. The wear behavior of electroless Ni–P coatings has been extensively examined in previous studies. Researchers have employed various tribological testing methodologies to evaluate their performance, with notable differences in contact configurations—such as pin-on-disc and ball-on-disc—as well as test environments, including lubricated, dry, and vacuum conditions, selected based on the intended application.^44,45 Electroless nickel coatings have emerged as suitable coatings that can serve as viable substitutes for traditional electroplating. Their properties, such as hardness, low friction, wear resistance, and corrosion resistance, have led to their usage in tribological applications.⁴⁶ Besides, the uniform deposition and the ability to coat any material have served as an added advantage to their application in various areas. The coatings' exceptional resistance to wear is one of the special features of electroless nickel deposition. However, a variety of additional factors, including the type of applied stress and the surface morphology, influence a surface's wear characteristics. The amount of phosphorus present and the kind of post-heat treatment used define how resistant electroless nickel deposits are to wear. The wear mechanism of electroless Ni–P deposits is governed by the attractive interactions between nickel atoms in the coating and iron atoms in the counter disc. A beneficial tribochemical polishing effect was observed between the frictional pairs, leading to the development of a relatively smooth surface morphology.⁴⁷ This process facilitated the formation of a reaction film composed primarily of P–O species along the wear track. Furthermore, the initial surface morphology of the coatings played a critical role in determining the distribution and concentration of P–O and Ni–O compounds on the worn surface. These variations in chemical composition and surface topography had a direct impact on the ability to achieve superlubricity.⁴⁸

3.3.2 Bismuth redox cycling as an electron trap strategy for oxygen-mediated tribological behaviour. In tribological systems, metal nanoparticles play a dual role by not only promoting carrier separation but also facilitating material transfer during frictional contact. This architecture increases the density of exposed active sites, thereby enlarging the effective frictional contact area. This expanded interface enables greater electron exchange, which in turn amplifies redox activity at the surface.⁴⁹ The mechanical force that generates friction between the catalyst and its surrounding environment promotes electron transfer across atomic interfaces.⁵⁰ These transferred electrons actively participate in surface chemical reactions, defining the functional mechanism of tribocatalysts.⁵¹ Materials capable of gaining or losing electrons are thus employed to generate reactive species that drive subsequent catalytic processes.⁵² Doping with bismuth can significantly enhance the tribological performance, primarily by introducing lattice defects. These defects act as nucleation sites for slip systems, thereby modifying the electronic structure and reducing the material's shear strength.⁵³ Additionally, in oxidation reactions catalyzed by metal oxides, lattice oxygen, particularly those atoms located at or near the surface serves as the primary active species, consistent with the widely accepted Mars–van Krevelen (MvK) mechanism.⁵⁴

3.3.3 Tailored tribological behavior in electroless NiP coating via bismuth doped gahnite integration. The coefficient of friction (COF) and specific wear rate of NiP, ZA, and ZnAl_2−xB_xO₄, x = 0.05, 0.10, 0.20, and 0.50 are presented in Fig. 5(A) and (B), and the data are provided in Table S3. Among the studied compositions, the ZAB_0.20 coating exhibited the lowest average COF (0.19) as calculated based on the data from Fig. 5(A) and the minimum wear rate (0.889 × 10⁻⁵ mm³/Nm), revealing its superior tribological performance. In contrast, the pure NiP coating demonstrated the highest wear rate (12.736 × 10⁻⁵ mm³/Nm) and COF (∼0.93), showing significant abrasive wear, most likely because of poor surface conformity under sliding stress and the lack of reinforcing ceramic phases. The undoped ZnAl₂O₄ phase's tribological performance is limited by its comparatively lower hardness and poor mechanical stability, which leaves it exposed to abrasive degradation under high contact stresses and constant sliding. The coatings were compared with those of the optimally synthesized Bi-doped ZnAl₂O₄ composite of the same composition to evaluate the role of compositional adjustment. The performance of the (ZAB_0.20) composite was noticeably better, even though both ZnAl₂O₄ and Bi₂O₃ show some positive impacts on wear reduction.

Fig. 5 (A) Coefficient of friction (COF) and (B) specific wear rate of NiP, ZA, and ZnAl_2−xBi_xO₄, x = 0.05, 0.10, 0.20, and 0.50.

In contrast to the synergistic response attained in the coupled system, this augmentation results from the individual phases' poor flexibility and weak interfacial stability when present alone. By introducing lattice distortion, oxygen vacancies, and polyhedral incommensurability (AO₄–BO₆), the preferential insertion of Bi³⁺ ions into the ZnAl₂O₄ spinel lattice stabilizes the structure under sliding. The development of a defect-rich, adaptive tribolayer that offers strong protection during prolonged tribological interaction is made possible by these effects, which also inhibits particle agglomeration and lower localized stress concentration. Table S4 compares our work's outcomes with those of other recently published studies.

3.4 Reactive surface chemistry and phase reconfiguration in worn coatings

3.4.1 Tribologically activated redox mechanisms and grain-scale structured reordering. The wear mechanism and the degree of the coating's resistance to wear and friction are evident from the assessment of the morphological changes that took place on the coating surface following the wear test. Fig. 6(A–D) displays the SEM images of the outside and inside wear tracks of the ZA, ZAB_0.20 coatings for comparison, and the corresponding elemental dot mapping and elemental distributions recorded for ZA and ZAB_0.20 are presented below Fig. 6(A–D), subsequently. Fig. 6(A1) displays the EDS mapping image of the outside wear track of ZA and the corresponding elemental overlay (A2) with Ni, P, Zn, Al, O & Fe. Fig. 6(B1) displays the EDS X-ray mapping image of the outside wear track of ZAB_0.20 and the corresponding elemental overlay (B2) with Ni, P, Zn, Al, Bi &O. Scratches and surface flaws are where nucleation preferentially occurs in electroless coatings.² In addition to having large grain sizes and poorly defined grain boundaries, the pure NiP coating contains pinholes on its surface as a result of many bubbles that form during electrodeposition, increasing surface crystallinity and significantly increasing oxygen element content.⁵⁵ The Orowan hardening mechanism explains that the dislocations wrapped around the precipitate act as dislocation pinning sites, producing a strong resistive force to the dislocation motion, even though the hardness of NiP diminishes as the phosphorus content increases.⁵⁶ According to the Hall–Petch effect, materials' strength and hardness improve as their grain size decreases.⁴⁶ A smaller grain size is good for coatings as it is less likely to get deformed when subjected to friction. An oxide film also has an antifriction effect to a certain extent;⁴⁷ high-density dislocations form and spread along sub-grain boundaries to lower distortion energy during repeated friction.² Grain growth, sliding, or rotation can lower wear resistance. However, nanomaterials can increase wear resistance. Their structure changes evenly during wear, which helps prevent softening and breaking.⁵⁶ Adding zinc aluminate nano particles to Ni–P coating can change the surface's shape and morphology. Degree of crystallinity, defect sites, and morphology are the distinct parameters for the interplay between the microstructure and oxidation resistance in lamellar solids.⁵⁷ At high temperatures, molecular oxygen interacts with edge sites and defect sites of ZA, and these interactions result in the development of oxides that can obstruct lamellar shear and other mechanical abrasion. In nanodomains over a globular Ni–P matrix, wear causes cracks on the track due to abrasion and oxidation. The surface mostly has long grooves and some pits along the sliding direction. Cracks and voids grow and may create a loose wear track.⁵⁸

Fig. 6 SEM images of ZA coatings: (A) outside and (B) inside wear tracks; SEM images of ZAB_0.20 coatings: (C) outside and (D) and inside wear tracks; (A1) EDS mapping image of the outside wear track of ZA, corresponding elemental overlay (A2) with Ni, P, Zn, Al, O & Fe; (B1) EDS X-ray mapping image of the outside wear track of ZAB_0.20, corresponding elemental overlay (B2) with Ni,P, Zn, Al, Bi & O.

Shear stress during sliding is greater than the strength of loose wear particles, causing them to detach from underlying layers. If shear stress exceeds the Cauchy strength of these particles, they easily peel off from the specimen's surface layer.⁵⁹ Grain dislocation slip is the primary mechanism for stress reduction. Grain boundaries, however, may obstruct these slips and make mobility more difficult. Therefore, it is true that smaller grains can increase resistance to stress relaxation. The pinning effect of bismuth atoms concentrated at grain boundaries begins to influence the migration of dislocations. The addition of bismuth stops the dislocation density from decreasing.⁴³ The thickness of the coating is estimated to be 226.1 µm and 343.4 µm for bare NiP and ZAB_0.20@NiP, respectively, shown in Fig. S11 (A) and (B). When the bismuth concentration increases to x = 0.50 in ZnAl_2−xBi_xO₄ in the ZAB_0.50@NiP matrix, the excess bismuth starts segregating to grain boundaries and interfaces, where it forms weak Bi-rich intergranular layers (Fig. S14).

Starting from the Al 2p region, we observe that the spectrum consists of two features. The low energy one at ∼68.0 eV corresponds to metallic aluminium as seen in the outside wear track Fig. 7(A), due to spin orbit splitting into Al 2p_3/2 & Al 2p_1/2.⁶⁰ An oxide peak at a higher binding energy is observed around ∼72 eV, which has been attributed to phonon broadening.⁶¹ The shape of the Al 2p spectrum does not change significantly, but the peak shifted to a lower binding energy, indicating that this is an oxidized aluminium state with an oxygen-deficient matrix.⁶² A shift in the binding energy is observed inside the wear track Fig. 7(B) indicates modification in the electronic environment brought about by tribo-induced reactions. Fig. 7(C) shows the O 1s spectrum exhibiting four oxygen contributions, with the lowest binding energies corresponding to Zn–O &Al–O (corresponds O²⁻ surrounded by Zn and Al in the fully oxidation state (M−O lattice oxygen).⁶¹ The other two binding energy peaks at ∼531 eV correspond to the oxygen vacancy (Ov), and the peak around ∼532 eV corresponds to chemisorbed species. A shift to higher binding energy observed inside the wear track (Fig. 7(D)) suggests alleviation of the M–O bond strength.⁶³Fig. 7(E) and (F)) depicts C 1s high resolution spectra showing four distinct peaks: C sp² hybridised carbon(C=C), C sp³ (C–C), the oxide functional groups observed at ∼286 eV (C–O) and (C=O) observed at ∼288.8 eV, attributed to the presence of obvious surface oxides.^64,65 In contrast, spectra from the worn region (inside the wear track) retain the sp², sp³, and C=O features, but the C–O contribution is absent. The Bi 4f fine scan spectra of the outside wear track deconvoluted into four peaks are shown in Fig. 7(G). The Bi 4f_5/2 and Bi 4f_7/2 peaks of positively charged Bi species (Bi³⁺) are responsible for the two peaks at ∼164.4 eV and ∼159.1 eV, whereas the Bi 4f_5/2 and Bi 4f_7/2 peaks of metallic Bi (Bi⁰) are responsible for the other two peaks at ∼162.7 eV and ∼157.3 eV.⁶⁶ The fitted curves of Bi 4f spectra in Fig. 7(H) displayed three well-defined doublets, which were associated with the multiple oxidation state contributions during sliding, and the peak observed at the binding energies of ∼159.1 eV & 164.1 eV with the separation of 5.1 eV corresponds to the Bi 4f_7/2 & Bi 4f _5/2 peaks of Bi³⁺ ionic state in Bi₂O₃. The peaks with binding energies of 160.5 eV and 165.7 eV are attributable to the Bi in the Bi⁵⁺ state.⁶⁷ In Fig. S13 the pairs of peaks at 852.7/869.9 eV and 855.6/873.5 eV were attributed to the Ni 2p_3/2 and Ni 2p_1/2 signals of metallic Ni and Ni²⁺. A new pair of peaks at 853.4/870.1 eV could be attributed to the Ni 2p_3/2 and Ni 2p_1/2 signals of Ni^δ+ (Ni–P bonds).⁴⁸

Fig. 7 XPS spectra of coatings: (A) Al 2p outside, (B) Al 2p inside, (C) O 1s outside, (D) O 1s inside, (E) C 1s outside, (F) C 1s inside, (G) Bi 4f outside and (H) Bi 4f inside the wear track of developed coatings.

3.4.2 Micro topographical deformation evaluation and metal oxide tribolayer confirmation in the worn zone. Surfaces with smoother textures typically exhibit more negative skewness (Ssk) values and higher kurtosis (Sku) values. However, the parameters Ssk and Sku do not show a direct relationship with friction or the time required to reach steady-state sliding. This lack of correlation is primarily due to the highly irregular and aggressive nature of abrasive wear mechanisms. Interestingly, surfaces with reduced friction often demonstrate lower Sku values. Additionally, surfaces characterized by high Sku and significantly negative Ssk values also tend to exhibit lower friction coefficients.

In terms of surface roughness, lower values of Sa and Sq generally lead to reduced friction under lubricating conditions.⁶⁸ A negative Ssk implies a generally flat surface interspersed with deeper grooves or pores, while a positive Ssk indicates a surface made up mostly of peaks and grooves rather than being predominantly flat. Surface kurtosis (Sku) serves as an indicator of the sharpness of surface height distribution. It reflects whether the surface has sharp protrusions or indentations. If Sku is greater than 3, it indicates the presence of sharp features, suggesting a rough texture that can concentrate stress. On the other hand, Sku values less than 3 suggest a more even surface with features that are not as sharply defined.⁶⁹ The Sa parameter reflects the average vertical deviation and provides insight into variations in surface topography. Fig. 8(A) shows the variation of roughness parameters outside and Fig. 8(B) shows the variation of roughness parameters inside the wear track. Fig. 8(C) and (D) show the 3D micro surface profile parameters outside and inside of the developed coatings. A positive skewness (Ssk > 0) suggests a surface dominated by peaks, while a negative skewness (Ssk < 0) implies that the surface is composed mostly of valleys. A detailed three-dimensional (3D) surface profile of the developed coatings NiP, ZA, ZAB_0.05, ZAB_0.10, ZAB_0.20, ZAB_0.50 are shown in Fig. S12.

Fig. 8 Variation of roughness parameters: (A) outside and (B) inside wear track; (C) & (D) are the 3D micro surface profile parameters outside and inside of the developed coatings.

Raman spectroscopy was used to examine the ZAB_0.20 phase stability and structural integrity under tribological stress in both the worn (inside wear track Fig. 9(A)) and unworn (outside wear track Fig. 9(B)) areas. Significant differences in the peak intensity and broadening are revealed by comparative spectral analysis, indicating tribo-induced changes in surface chemistry or local bonding conditions, most likely connected to stress-induced lattice distortion. This allowed for a clear distinction between spectroscopic chemical information originating within and outside of the wear scars. First-order active Raman modes are non-degenerate A₁g, B₁g, and B₂g modes and doubly degenerate Eg modes. The silent modes are A₂g and B₁u. The longitudinal optical (LO), transverse optical (TO), and acoustic modes are linked to two Eu and one A 2u modes. Bands formed in the 200–550 cm⁻¹ range are the result of the Bi–O and Bi–O–Bi vibrations. The oxygen vibration in the Bi₂O₃ structure is the cause of the peaks around 223 cm⁻¹ and 278 cm⁻¹.^70,71 The peak at 405 cm⁻¹ is attributed to the Bi–O–Bi bonds in [BiO₆].⁷² The weak Raman peak near 531 cm⁻¹ in Bi₂O₃ crystals is commonly associated with structural imperfections, particularly oxygen vacancy-related defects.⁷³ The α-Fe₂O₃ (hematite) phase, which crystallizes in the D⁶_3d space group, exhibits seven Raman-active vibrational modes—specifically, two A₁g and five Eg modes. Among them, the Raman band near 497 cm⁻¹ corresponds to the A₁g symmetric stretching vibration, while the peaks located at approximately 242.3 cm⁻¹ and 609.1 cm⁻¹ are attributed to Eg modes associated with Fe–O lattice vibrations.⁷⁴ The characteristic vibrational modes of ZnAl₂O₄, specifically the Eg mode around 420 cm⁻¹, the T₂g mode near 659 cm⁻¹, and the A₁g mode at 780 cm⁻¹, are discernible despite being superimposed on broad background features.⁷⁵

Fig. 9 Raman spectra of the (A) outside and (B) inside wear track of ZAB_0.20 coatings.

The Raman D band observed around 1308 cm⁻¹ is attributed to defect-induced vibrational modes of sp³-hybridized carbon within a predominantly sp²-bonded carbon matrix. This typically originates from disordered structures, edge defects, or amorphous carbon regions formed under tribological stress. The absence of the G band suggests a high degree of disorder or a dominance of non-graphitic sp³ carbon phases in the tribofilm.⁷⁶ Bi atom participation in the lattice is indicated by the significant peak at 96.4 cm⁻¹ in the area outside the wear track, which is due to Ag symmetry.⁷³ Oxygen vacancies are probably responsible for the weak Raman signal seen at 534.5 cm⁻¹,⁶⁹ whereas the A₁g vibrational mode of the ZnAl₂O₄ spinel structure is represented by the band at 783 cm⁻¹.⁷⁵

3.5 Tribological stabilization pathway induced by spinel distortion and the redox driven layering mechanism

Fig. 10 illustrates the tribological pathway through wear resistance on the gahnite surface. The development of a complex glazing layer made up of several oxides, such as Bi₂O₃, NiO, Al₂O₃, and Fe₂O₃, is responsible for the system's improved tribological performance.² By serving as a tribochemically active protective layer, this glazing layer helps to minimize material loss during sliding. Despite being mechanically robust and structurally dense, the Bi₂O₃-enriched outer layer's intrinsic stiffness and limited deformability might cause increased wear and friction when in dry contact. The reason for this is that elemental bismuth's quasi-metallic, densely packed layered structure at room temperature changes into a stiffer oxide phase with tribo-oxidation, which lessens its capacity to conform plastically under shear.¹⁰ Through its special framework, the ZnAl₂O₄-type AB₂O₄ spinel structure adds structural lubricity. The incommensuration between AO₄ tetrahedra and BO₆ octahedra resists deformation, enhancing load distribution and reducing abrasive interactions. Without sacrificing friction stability, this microstructural stiffness improves wear resistance.⁴ The production and renewal of tribofilms are anticipated to be greatly aided by the redox-active cationic Bi³⁺ species. Reversible Bi³⁺/Bi⁵⁺ redox cycling is supported by their electronic structure, which is typified by weakly coordinated monoanionic ligands, a vacant p-orbital, and an occupied s-orbital. Intermolecular σ- and π-type interactions are facilitated by the unoccupied p-orbital, promoting oxygen exchange dynamics that are crucial for tribolayer recovery. Even if it has less of an impact, the occupied s-orbital could help through weak dative interactions. This redox flexibility is particularly relevant under sliding conditions.⁷⁷ The Mars–van Krevelen (MvK)-type redox mechanism can effectively describe the dynamic oxygen exchange behavior observed in the tribofilm, particularly involving the bismuth redox cycle.⁷⁸ During sliding, Bi³⁺ species at the surface may be reduced to metallic Bi⁰ due to oxygen loss driven by frictional interactions or chemical reactions with counterface species. This corresponds to the formation of oxygen vacancies, locally disrupting the oxide matrix. Over time, reoxidation occurs via migration of lattice oxygen from the bulk toward the surface vacancy sites, restoring the oxide layer. In this process, Bi⁰ is reoxidized to Bi³⁺ or even Bi⁵⁺, enabling a reversible redox shuttle among Bi⁰/Bi³⁺/Bi⁵⁺ states.

Fig. 10 Mechanistic pathway for wear reduction.

4 Conclusions

This work explores an innovative approach by employing the defect-engineered spinel framework of Bi-doped zinc aluminate (gahnite) to substantially enhance the wear resistance of steel substrates via coupled defect chemistry and redox dynamics of bismuth. The study systematically unveils the impact of oxygen-vacancy-rich, self-adaptive tribolayers—comprising a mixture of spinel oxide, and tribo-oxide phases on the sustained low-friction and wear-resistant behavior of the Bi–ZnAl₂O₄ composite coating. By preferentially incorporating bismuth ions into the ZnAl₂O₄ spinel lattice, a coupled lattice distortion is generated, and oxygen vacancies that potentially prevent structural relaxation under tribological loading are created. The suitable morphological tuning of the ZnAl₂O₄ spinel framework by the unique lattice incorporation of bismuth facilitates the effective interaction between oxygen-vacancy sites and the surrounding oxide matrix, resulting in a stabilized and defect-rich surface. Additionally, the structural integrity of the spinel matrix with inherent structural incommensurability between AO₄ tetrahedra and BO₆ octahedra imparts localized lattice flexibility. The physicochemical nature of the tribolayer formed during sliding plays a decisive role in determining the coating's tribological performance. Under mechanical loading, the ZnAl₂O₄ framework in synergy with the tribochemically generated phases including Bi₂O₃, ZnO, Al₂O₃, and Fe₂O₃, develops a conformal, stress-dissipating interfacial layer. This adaptive tribolayer accommodates shear, mitigates localized contact stresses, and thereby imparts superior wear resistance. The effective compositional tuning of the ZnAl_1.98Bi_0.20O₄ composite coating facilitates stress accommodation and lattice-scale adaptability owing to the combined effects of defect engineering and AO₄–BO₆ polyhedral incommensurability. This structural configuration promotes controlled interfacial shear and minimizes stress concentration on the coating surface, resulting in an exceptionally low wear rate of 0.889 × 10⁻⁵ mm³ N⁻¹ m⁻¹ and a coefficient of friction of 0.19, which is lower than that of the undoped ZnAl₂O₄ coating (0.72). The tribo-induced transformation of the surface spinel structure, facilitated by bismuth redox cycling and oxygen-vacancy dynamics, further stabilizes the evolving tribofilm through the formation of mechanically robust oxide phases. This structural adaptation contributes significantly to the exceptional tribological performance of the coating.

The present findings establish a strategic pathway for overcoming the persistent challenge of wear resistance in demanding tribological environments. The specifically engineered spinel coating, with its exceptional combination of defect chemistry and structural incommensurability, exhibits outstanding tribological performance, holding strong potential as a high-performance protective coating.

Author contributions

S. Sameera: methodology, investigation, formal analysis, writing – original draft, visualization, validation, review and editing; Najiya Nasirin: formal analysis, writing – original draft, visualization, validation, review and editing; A. H. Dalaver: validation, resources, investigation, review and editing; S. M. A. Shibli: conceptualization, methodology, supervision, review and editing.

Conflicts of interest

There are no conflicts to declare.

Data availability

The data supporting this article have been included as part of the supplementary information (SI). Supplementary information: XPS survey spectrum of ZA, high resolution images of Zn 2p, and Al 2p; stacked XRD spectra of all prepared samples; graphical output of the Rietveld refinement of other samples; W–H plot of other samples; XPS survey spectra and high resolution spectra of ZAB_0.20; EDX spectrum of ZA and ZAB_0.20; selected area for mapping and element overlay of ZA and ZAB_0.20; FESEM micrographs, EDX spectra and corresponding mapping of ZAB_0.05, ZAB_0.10 andZAB_0.50; particle size distribution from TEM measurement measurement of ZA and ZAB_0.20; TEM-EDS elemental analysis ZA and ZAB_0.2; cross sectional SEM images of coatings; OSP images of outside and inside wear tracks of coatings; high resolution XPS image of Ni2p outside wear track of coatings; SEM image of ZAB_0.5 inside wear track; structure refinement values obtained by Rietveld refinement; TEM-EDS elemental composition of ZA and ZAB_0.20; comparison of wear parameters of the host matrix and tuned coatings; comparison with other wear resistive materials. See DOI: https://doi.org/10.1039/d5ta07478a.

Acknowledgements

The authors acknowledge the Department of Chemistry & Centre for Renewable Energy and Materials (CREM), University of Kerala, Thiruvananthapuram, India, for providing the facilities to carry out the work. The authors thank Core Technology Platform (CTP) resources at New York University (NYU) Abu Dhabi for TEM experiments. The authors also acknowledge the CLIF-University of Kerala, Thiruvananthapuram, India for extending their material characterization facilities. Najiya Nasirin acknowledges DST-INSPIRE for a Junior Research Fellowship (Ref. No. IF 220634).

References

1. H. Liu, X. Wang, L. Zhao, S. Yang and X. Zhang, ACS Appl. Mater. Interfaces, 2024, 16, 59096–59108.
2. H. Sun, S. Wan, G. Yi, S. T. Pham, S. M. Collins, B. Xin, J. Chen, J. Yang and W. Wang, Wear, 2024, 538–539, 205189.
3. P. Uniyal, P. Gaur, J. Yadav, T. Khan and O. S. Ahmed, ACS Omega, 2024, 9, 12436–12456.
4. Z. Zhang, E. Hershkovitz, Q. An, L. Liu, X. Wang, Z. Deng, G. Baucom, W. Wang, J. Zhao, Z. Xin, L. Moore, Y. Yao, M. R. U. Islam, X. Chen, B. Cui, L. Li, H. Xin, L. Li, H. Kim and W. Cai, Nat. Commun., 2024, 15, 10039.
5. Q. Shen, J. Hao, Z. Zhu, S. Bai, J. Wu and S. Xuan, Ceram. Int., 2025, 51(6), 8107–8115.
6. A. Erdemir, Surf. Coat. Technol., 2005, 200, 1792–1796.
7. W. Gulbiński and T. Suszko, Surf. Coat. Technol., 2006, 201, 1469–1476.
8. S. M. Aouadi, Y. Paudel, B. Luster, S. Stadler, P. Kohli, C. Muratore, C. Hager and A. A. Voevodin, Tribol. Lett., 2008, 29, 95–103.
9. V. Ageh, H. Mohseni and T. W. Scharf, Surf. Coat. Technol., 2013, 237, 241–247.
10. P. Gonzalez-Rodriguez, K. J. H. van den Nieuwenhuijzen, W. Lette, D. J. Schipper and J. E. ten Elshof, ACS Appl. Mater. Interfaces, 2016, 8, 7601–7606.
11. Z. Chen, P. Shao, P. Xiong, Y. Xiao, B. Liu, Z. Wang, S. Wu, D. Jiang, K. Chen, J. Gan, D. Chen and Z. Yang, ACS Appl. Mater. Interfaces, 2024, 16(27), 35279–35292.
12. X. Shao, Y. Ren, K. Guo, C. Du, X. Wang, J. Chen, J. Liu, L. Wang and J. Yang, Tribol. Int., 2025, 206, 110577.
13. C. Domínguez-Ríos, A. Hurtado-Macias, R. Torres-Sánchez, M. A. Ramos and J. González-Hernández, Ind. Eng. Chem. Res., 2012, 51, 7762–7768.
14. Y. Wang, G. Zhang, D. Yang, Y. Li, W. Yin and J. Xu, Langmuir, 2024, 40(8), 4306–4313.
15. M. Ni, S. Wang, W. Li and W. Huang, Mater. Res. Express, 2019, 6, 116413.
16. Y. He, S. C. Wang, F. C. Walsh, Y.-L. Chiu and P. A. S. Reed, Surf. Coat. Technol., 2016, 307, 926–934.
17. S. Li, S. Pu, Z. You, C. Sun, S. Li and J. Zhang, Trans. IMF, 2020, 98, 21–29.
18. K. Shahzad, A. B. Radwan, O. Fayyaz, R. A. Shakoor, M. Uzma, M. A. Umer, M. N. Baig and A. Raza, Ceram. Int., 2021, 47, 19123–19133.
19. A. R. K. Rana and Z. Farhat, Surf. Coat. Technol., 2019, 369, 334–346.
20. A. P. Meshram, M. K. Punith Kumar and C. Srivastava, Diamond Relat. Mater., 2020, 105, 107795.
21. F. Du, C. Li, D. Li, X. Sa, Y. Yu, C. Li, Y. Yang and J. Wang, Lubricants, 2022, 10, 196.
22. M. Cichomski, W. Stanek, R. Stanecka-Badura, D. Świątczak, M. Małecka, J. Kowalczyk, M. Madej and M. Dudek, Appl. Surf. Sci., 2025, 710, 164014.
23. H. Gao, A. Otero-de-la-Roza, J. Gu, D. Stone, S. M. Aouadi, E. R. Johnson and A. Martini, ACS Appl. Mater. Interfaces, 2015, 7, 15422–15429.
24. S. Wang, D. Chen, Q. Hong, Y. Gui, Y. Cao, G. Ren and Z. Liang, J. Mol. Liq., 2023, 389, 122821.
25. X. Song, S. Zheng, J. Zhang, W. Li, Q. Chen and B. Cao, Mater. Res. Bull., 2012, 47, 4305–4310.
26. T. Thi Ngoc Nha, D. N. Toan, P. H. Nam, D. H. Manh, D. T. Khan and P. T. Phong, J. Alloys Compd., 2024, 996, 174773.
27. M. G. Brik, A. Suchocki and A. Kaminska, Inorg. Chem., 2014, 53, 5088–5099.
28. S. Sameera, V. Vidyadharan, S. Sasidharan and K. G. Gopchandran, J. Sci.:Adv. Mater. Devices, 2019, 4, 524–530.
29. J. Gao, J. Xue, S. Jia, Q. Shen, X. Zhang, H. Jia, X. Liu, Q. Li and Y. Wu, ACS Appl. Mater. Interfaces, 2021, 13, 18758–18771.
30. N. Srinatha, S. Satyanarayana Reddy, K. Gurushantha, M. Al-Dossari, S. O. Manjunatha and A. Madhu, Ceram. Int., 2024, 50, 37742–37753.
31. C. Guan, X. Yue and Q. Xiang, Adv. Mater., 2025, 37, 2501209.
32. S.-W. Chan and W. Wang, Mater. Chem. Phys., 2021, 273, 125091.
33. Y.-X. Chen, D. Zhou, X. Yang, X.-J. Guo, J.-Y. Zhang and J.-K. Liu, Ind. Eng. Chem. Res., 2021, 60, 10064–10075.
34. Q. Ren, Y. He, H. Wang, Y. Sun and F. Dong, ACS Catal., 2022, 12, 14015–14025.
35. J. Zhao, M. Zhou, L.-Y. Xiao, Z.-H. Zhou, X.-W. Du and J. Yang, ACS Appl. Energy Mater., 2023, 6(5), 3147–3154.
36. V. Babin, V. Hreb, A. Krasnikov, V. Stadnik, L. Vasylechko and S. Zazubovich, Mater. Res. Bull., 2026, 193, 113710.
37. V. Laguta, M. Buryi, P. Arhipov, O. Sidletskiy, O. Laguta, M. G. Brik and M. Nikl, Phys. Rev. B, 2020, 101, 024106.
38. S. Jyoti, A. Vijay, U. Terranova, S. K. Gupta, K. Sudarshan and S. Vaidya, Phys. Chem. Chem. Phys., 2024, 26, 17324–17333.
39. X.-P. Wei, Y.-T. Yang, S.-Y. Li, Y. Meng, W.-B. Yuan and H.-G. Ni, Environ. Res., 2025, 282, 122042.
40. V. Singh, R. P. S. Chakradhar, J. L. Rao and D.-K. Kim, J. Lumin., 2008, 128, 394–402.
41. Z. Wu, M. Guo, Y. Yan, H. Dou, W. Zhao, Y. Zhang, S. Li, J. Wu, X. Bin, X. Zhao, X. Yang and D. Ruan, ACS Sustainable Chem. Eng., 2021, 9, 7422–7430.
42. L. E. Sotelo Martin, N. M. O'Shea, J. K. Mason and R. H. R. Castro, J. Phys. Chem. C, 2023, 127, 4239–4250.
43. Z. Zhang, H. Liu, Y. Hua, B. Zhang, G. Wang, Z. Hou, W. Li and K. Song, Mater. Sci. Eng., A, 2024, 916, 147345.
44. C. Orgeldinger, T. Rosnitschek and S. Tremmel, Wear, 2024, 552, 205437.
45. Y. Meng, J. Xu, Z. Jin, B. Prakash and Y. Hu, Friction, 2020, 8, 221–300.
46. Q. Chen, S. Xu, Y. He, S. Yan, Y. Fan, Z. Li, H. Zhou, H. Li, K. Wei, X. Gong and X. Cheng, Langmuir, 2024, 40, 16400–16418.
47. H. Liu, Z. Li, X. Zhang, H. Chen, S. Xu, Y. Fan, Q. Yuan, Y. Sun, X. Hou and Y. He, J. Phys. Chem. C, 2024, 128, 20679–20692.
48. H. Zhang, K. Le, W. Yang, Y. Zhou, J. Jin, Y. Sun, S. Zhang and S. Xu, Langmuir, 2025, 41(26), 17209–17218.
49. X. Li, W. Tong, J. Shi, Y. Chen, Y. Zhang and Q. An, J. Mater. Chem. A, 2023, 11, 4458–4472.
50. C. Xu, Y. Zi, A. C. Wang, H. Zou, Y. Dai, X. He, P. Wang, Y.-C. Wang, P. Feng, D. Li and Z. L. Wang, Adv. Mater., 2018, 30, 1706790.
51. Y. Xu, R. Yin, Y. Zhang, B. Zhou, P. Sun and X. Dong, Langmuir, 2022, 38(46), 14153–14161.
52. Q. Zhang, Y. Jia, X. Wang, L. Zhang, G. Yuan and Z. Wu, Sep. Purif. Technol., 2023, 307, 122846.
53. A. K. Singh, N. Shukla, Kavita, D. K. Verma, K. D. Mandal and R. B. Rastogi, Ind. Eng. Chem. Res., 2024, 63(46), 20034–20046.
54. K. Yu, L.-L. Lou, S. Liu and W. Zhou, Adv. Sci., 2020, 7, 1901970.
55. C. Liu, Z. Li, W. Lu, Y. Bao, W. Xia, X. Wu, H. Zhao, B. Gault, C. Liu, M. Herbig, A. Fischer, G. Dehm, G. Wu and D. Raabe, Nat. Commun., 2021, 12, 5518.
56. T. J. Rupert and C. A. Schuh, Acta Mater., 2010, 58, 4137–4148.
57. J. F. Curry, M. A. Wilson, H. S. Luftman, N. C. Strandwitz, N. Argibay, M. Chandross, M. A. Sidebottom and B. A. Krick, ACS Appl. Mater. Interfaces, 2017, 9, 28019–28026.
58. H. Algul, M. Uysal and A. Alp, Appl. Surf. Sci. Adv., 2021, 4, 100089.
59. D. R. Dhakal, Y. K. Kshetri, G. Gyawali, T.-H. Kim, J.-H. Choi and S. W. Lee, Appl. Surf. Sci., 2021, 541, 148403.
60. L. Savio, K. B. Bhavitha, G. Bracco, G. Luciano, D. Cavallo, G. Paolini, S. Passaglia, G. Carraro, L. Vattuone, R. Masini and M. Smerieri, Appl. Surf. Sci., 2021, 542, 148574.
61. P. H. Citrin, P. Eisenberger and D. R. Hamann, Phys. Rev. Lett., 1974, 33(16), 965–969.
62. H. Asgarimoghaddam, Q. Chen, F. Ye, A. Shahin, B. Song and K. P. Musselman, Small Methods, 2024, 8, 2300995.
63. Q. Huang, G.-J. Xia, B. Huang, D. Xie, J. Wang, D. Wen, D. Lin, C. Xu, L. Gao, Z. Wu, J. Wu, F. Xie, W. Guo and R. Zou, Energy Environ. Sci., 2024, 17(14), 5260–5272.
64. F. M. Elam, F.-C. Hsia, S. van Vliet, R. Bliem, L. Yang, B. Weber and S. E. Franklin, Carbon, 2021, 179, 590–599.
65. D. Dinesh Kumar, S. Hazra, K. Panda, P. Kuppusami, T. Stimpel-Lindner and G. S. Duesberg, ACS Appl. Mater. Interfaces, 2022, 14, 26148–26161.
66. Y. Wan, H. Zhou, M. Zheng, Z.-H. Huang, F. Kang, J. Li and R. Lv, Adv. Funct. Mater., 2021, 31(30), 2100300.
67. J. Pandey, N. Yadav, P. Yadav, S. Rao, K. Malla, P. Koirala, M. P. Ghimire and R. Nagarajan, Inorg. Chem., 2023, 62(24), 9471–9483.
68. M. Sedlaček, B. Podgornik and J. Vižintin, Tribol. Int., 2012, 48, 102–112.
69. M. Chen, H. Che and S. Yue, Surf. Coat. Technol., 2022, 450, 128993.
70. A. ur R. Khan, M. Ramzan, S. J. F. Alanazi, A. M. Al-Mohaimeed, S. Ali, M. Imran, M. A. Majid and M. H. Sarfraz, ACS Omega, 2024, 9, 22650–22659.
71. N. Wang, M. Yi, M. Dai, H. Wang, N. Zhang, T. Zhang, J. Zhang and B. Zhang, J. Alloys Compd., 2025, 1020, 179493.
72. P. Wang, Y. Ye, G. Zhao, J. Hou, Y. Liu, X. Qiao, Z. Wang, B. Li and Y. Fang, J. Non-Cryst. Solids, 2024, 628, 122848.
73. S. C. Lims, M. Jose, S. Aswathappa, S. S. J. Dhas, R. S. Kumar and P. V. Pham, Mater. Charact., 2024, 216, 114294.
74. N. Garg, S. Bera, G. Mangamma, V. K. Mittal, R. Krishnan and S. Velmurugan, Surf. Coat. Technol., 2014, 258, 597–604.
75. P. Loiko, A. Belyaev, O. Dymshits, I. Evdokimov, V. Vitkin, K. Volkova, M. Tsenter, A. Volokitina, M. Baranov, E. Vilejshikova, A. Baranov and A. Zhilin, J. Alloys Compd., 2017, 725, 998–1005.
76. S. T. Pham, S. Wan, K. A. Tieu, M. Ma, H. Zhu, H. H. Nguyen, D. R. G. Mitchell and M. Nancarrow, Adv. Mater. Interfaces, 2019, 6, 1901081.
77. C. Lichtenberg, Chem. Commun., 2021, 57, 4483–4495.
78. Magnetic Spinels: Synthesis, Properties and Applications, ed. M. Seehra, BoD–Books on Demand, 2017.

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Footnote

† These authors contributed equally to this work and share first authorship.

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